diff --git a/.gitignore b/.gitignore new file mode 100644 index 0000000..f552a67 --- /dev/null +++ b/.gitignore @@ -0,0 +1,17 @@ +# Local/generated build dependencies +node_modules/ + +# SMA 2026 generated deck outputs and intermediate work +金刚石散热项目/SMA2026_pitch_deck/ +金刚石散热项目/SMA2026_image_deck/ + +# Local extraction and helper scripts +金刚石散热项目/_extracted/ +金刚石散热项目/_scripts/ + +# Mineru conversion byproducts; keep Markdown files and README tracked +金刚石散热项目/项目前期工作_MD/**/*_origin.pdf +金刚石散热项目/项目前期工作_MD/**/*_layout.pdf +金刚石散热项目/项目前期工作_MD/**/*_model.json +金刚石散热项目/项目前期工作_MD/**/*_middle.json +金刚石散热项目/项目前期工作_MD/**/*_content_list*.json diff --git a/Mineru_api_client-V2.py b/Mineru_api_client-V2.py new file mode 100644 index 0000000..b560f7c --- /dev/null +++ b/Mineru_api_client-V2.py @@ -0,0 +1,129 @@ +import os +import requests +import zipfile +import io +import shutil # 【新增】用于删除非空文件夹 +import argparse # <--- 1. 确保这里导入了 + +def main(): + # 1. 配置命令行参数解析 + parser = argparse.ArgumentParser(description="Mineru PDF 转 Markdown 批量处理工具") + + # 添加参数 + parser.add_argument("-s", "--source", type=str, default="./Papers/ORI_PDF", + help="源 PDF 文件夹路径 (默认: ./Papers/ORI_PDF)") + parser.add_argument("-t", "--target", type=str, default="./Papers/ORI_MD", + help="目标 Markdown 文件夹路径 (默认: ./Papers/ORI_MD)") + parser.add_argument("-u", "--url", type=str, default="http://10.168.1.103:4000/extract", + help="API 服务端地址 (默认: http://10.168.1.103:4000/extract)") + parser.add_argument("--request-timeout", type=int, default=600, + help="单个 PDF 请求超时时间,单位秒 (默认: 600)") + parser.add_argument("--sync", action="store_true", + help="任务完成后是否开启反向同步清理(删除多余的输出文件夹)") + + args = parser.parse_args() + + # 使用解析后的参数 + url = args.url + source_dir = args.source + target_dir = args.target + + # 2. 确保源目录存在,避免报错 + if not os.path.exists(source_dir): + print(f"错误: 找不到源文件夹 '{source_dir}'") + exit(1) + + # 确保目标主文件夹存在 + os.makedirs(target_dir, exist_ok=True) + + # 3. 遍历源文件夹下的所有文件进行上传处理 + for filename in os.listdir(source_dir): + # 只处理 PDF 文件 + if not filename.lower().endswith(".pdf"): + continue + + file_path = os.path.join(source_dir, filename) + # 获取去掉 .pdf 后缀的文件名 + pdf_name_no_ext = os.path.splitext(filename)[0] + + # 对应的输出文件夹路径 + output_folder = os.path.join(target_dir, pdf_name_no_ext) + + # ========================================== + # 检查是否已经处理过 + # 如果输出文件夹已存在,且内部有文件,则视为已转换,直接跳过 + # ========================================== + if os.path.exists(output_folder) and len(os.listdir(output_folder)) > 0: + print(f"⏩ 已存在转换结果,跳过处理: {filename}") + continue + + print(f"正在上传并处理 {filename}...") + + try: + # 4. 发送请求 + with open(file_path, "rb") as f: + files = {"file": (filename, f, "application/pdf")} + response = requests.post(url, files=files, timeout=args.request_timeout) + + # 5. 处理响应结果 + if response.status_code == 200: + # 检查服务端是否返回了包含报错信息的 JSON + if response.headers.get('content-type') == 'application/json': + error_msg = response.json() + print(f"❌ 服务端处理失败 ({filename}):{error_msg.get('message', '未知错误')}") + continue # 跳过解压,处理下一个 + + # 确保该 PDF 专属的输出文件夹存在 + os.makedirs(output_folder, exist_ok=True) + + # 核心:使用 io.BytesIO 直接在内存中读取 zip 内容并解压 + try: + with zipfile.ZipFile(io.BytesIO(response.content)) as zip_ref: + zip_ref.extractall(output_folder) + print(f"✅ 成功!已解压并保存至文件夹: {output_folder}") + except zipfile.BadZipFile: + print(f"❌ 失败!{filename} 返回的内容不是有效的 ZIP 格式。") + else: + print(f"❌ 失败!{filename} 状态码: {response.status_code}, 报错: {response.text}") + + except requests.exceptions.RequestException as e: + print(f"❌ 网络请求异常 ({filename}): {e}") + except Exception as e: + print(f"❌ 处理 {filename} 时发生未知错误: {e}") + + print("-" * 30) + print("转换任务结束,开始进行文件夹同步清理...") + + # ========================================== + # 【新增逻辑】:反向比对,清理多余的 MD 文件夹 + # ========================================== + # 1. 收集当前 ORI_PDF 中所有有效的 PDF 名字(不含后缀) + valid_pdf_names = set() + for filename in os.listdir(source_dir): + if filename.lower().endswith(".pdf"): + valid_pdf_names.add(os.path.splitext(filename)[0]) + + # 2. 遍历 ORI_MD 文件夹 + if os.path.exists(target_dir): + for folder_name in os.listdir(target_dir): + folder_path = os.path.join(target_dir, folder_name) + + # 仅处理文件夹(以防里面有意外的独立文件) + if os.path.isdir(folder_path): + # 3. 如果这个文件夹的名字不在有效的 PDF 列表中,说明是被删掉的“孤儿” + if folder_name not in valid_pdf_names: + print(f"🧹 发现多余文件,正在清理: {folder_name}") + try: + # 使用 shutil.rmtree 删除整个文件夹及其内部所有内容 + shutil.rmtree(folder_path) + except Exception as e: + print(f"❌ 删除 {folder_name} 时发生错误: {e}") + + print("-" * 30) + print("所有批量任务彻底完成!") + +if __name__ == "__main__": + try: + main() + except KeyboardInterrupt: + print("\n检测到用户中断,程序退出。") diff --git a/博志金钻项目基本情况_知识库/wiki/Knowledge/SMA2026路演素材包.md b/博志金钻项目基本情况_知识库/wiki/Knowledge/SMA2026路演素材包.md new file mode 100644 index 0000000..7847102 --- /dev/null +++ b/博志金钻项目基本情况_知识库/wiki/Knowledge/SMA2026路演素材包.md @@ -0,0 +1,56 @@ +# SMA 2026 路演素材包 + +本页面向 `Samsung Mobile Advance 2026` 金刚石项目路演,汇总公司基础口径、可快速调用的图片 ID、以及与三星移动终端场景最相关的证据链。正式外发前,仍需回看对应 `Sources/PPT` 原页确认措辞和图片使用边界。 + +## 路演定位 + +建议将博志金钻描述为高功率芯片热管理封装材料与器件公司,核心能力不是单一材料供应,而是围绕高导热基体、金属化膜层、图形化、电镀/化镀、切割与检测形成的“材料-界面-封装器件”工程化平台。 + +SMA 2026 场景下,金刚石项目应聚焦: + +- Advanced Materials: diamond-based heat spreading and packaging substrates. +- Power / AI / XR / Wearable: high heat-flux, limited-space thermal paths. +- Sensors / Camera / Audio modules: thermal stability and drift reduction. +- Foldable devices: thin, shaped heat spreaders for constrained mechanical envelopes. + +## 公司基础口径 + +| 主题 | 推荐表述 | 证据入口 | +|---|---|---| +| 公司定位 | High-power chip thermal-management packaging materials and devices | [公司身份与定位](公司身份与定位.md) | +| 产品基础 | Stable ceramic substrate production, with diamond-metal composites in pilot stage | [产品与技术管线](产品与技术管线.md) | +| 工程能力 | PVD, photolithography, electroplating, electroless plating, cutting, inspection and process control | [制造能力与设备](制造能力与设备.md) | +| 路演主张 | Turn diamond thermal materials into evaluable packaging modules for Samsung Mobile PoC | [金刚石项目路演关联点](金刚石项目路演关联点.md) | + +## 快速图片调用 + +使用口令: + +```text +请调用 Image ID `{Image ID}`,读取对应图片笔记,并结合 Source PPT note 的 slide 文本确认上下文后用于 SMA 2026 路演。 +``` + +| SMA 用途 | Image ID | 建议用法 | +|---|---|---| +| 公司封面 / 品牌起页 | `SLIDE-ppt03_company_intro_20250412-S001` | 可作为公司介绍背景或素材参考,建议在新PPT中重绘成更国际化封面。 | +| 公司能力概览 | `SLIDE-ppt03_company_intro_20250412-S002` | 提取苏州工厂、南通工厂、研发实验室和认证信息。 | +| 热管理痛点 | `SLIDE-ppt03_company_intro_20250412-S003` | 用于说明芯片小型化、高性能带来的散热压力。 | +| 封装热路径 | `SLIDE-ppt03_company_intro_20250412-S004` | 可重绘为三星移动端 SoC / VC / diamond layer 的系统热路径图。 | +| 产品管线 | `SLIDE-ppt03_company_intro_20250412-S005` | 用于证明金刚石项目处在公司既有产品路线中。 | +| DPC / 金刚石载板流程 | `SLIDE-ppt03_company_intro_20250412-S007` | 用于支撑“材料进入器件”的工艺链。 | +| 金刚石-金属复合材料 | `SLIDE-ppt03_company_intro_20250412-S009` | 用于金刚石铜、金刚石铝、散热垫片和新一代封装材料说明。 | +| 生产能力矩阵 | `SLIDE-ppt03_company_intro_20250412-S010` | 用于展示机加工、激光加工、PVD、粉末烧结、光刻、电镀、化镀、切割等能力。 | +| 技术优势 | `SLIDE-ppt03_company_intro_20250412-S011` | 用于提炼界面结合、膜系结构、合金相、高精度图形化和薄膜形态控制。 | +| 设备与检测 | `SLIDE-ppt03_company_intro_20250412-S012`, `SLIDE-ppt03_company_intro_20250412-S013` | 用于证明样品制造、检测和迭代闭环。 | + +完整图片库见 [图片快速索引](../Sources/Images/图片快速索引.md),精选清单见 [路演图片调用清单](路演图片调用清单.md)。 + +## 建议外部话术 + +> Bozhi Jinzhuan is developing diamond-based thermal packaging modules for next-generation mobile devices, combining ultra-high thermal conductivity materials with metallization, patterning and device-level integration. For Samsung Mobile Advance 2026, the proposed PoC will focus on validating where diamond heat spreaders or diamond-metal composite submounts can reduce hotspot temperature, stabilize modules and fit into space-constrained mobile architectures. + +## 使用边界 + +- 2024 英文验厂材料中的客户项目细节仅作内部工艺能力证据,不建议直接外显客户名称或问题细节。 +- 研发人数、专利数量、估值等口径在不同PPT中存在版本差异,正式提交前需统一。 +- 原PPT截图适合内部定位素材;正式路演建议重绘成英文、少字、高可信图示。 diff --git a/博志金钻项目基本情况_知识库/wiki/index.md b/博志金钻项目基本情况_知识库/wiki/index.md index d3d6d12..8f5f00c 100644 --- a/博志金钻项目基本情况_知识库/wiki/index.md +++ b/博志金钻项目基本情况_知识库/wiki/index.md @@ -10,6 +10,7 @@ - [Sources/PPT](Sources/PPT/):4 份 PPT 的逐页来源笔记 - [图片快速索引](Sources/Images/图片快速索引.md):70 张幻灯片预览图 + 419 张 PPT 内嵌图片 - [幻灯片缩略图总览](Sources/Images/幻灯片缩略图总览.md):4 张 contact sheet,先看图再定位 Image ID +- [SMA2026路演素材包](Knowledge/SMA2026路演素材包.md):面向三星路演的公司口径与图片调用入口 - [路演图片调用清单](Knowledge/路演图片调用清单.md):精选路演可复用图片 - [证据台账](Knowledge/证据台账.md) - [版本差异与待核对](Knowledge/版本差异与待核对.md) @@ -21,6 +22,7 @@ - [制造能力与设备](Knowledge/制造能力与设备.md) - [市场与基地规划](Knowledge/市场与基地规划.md) - [金刚石项目路演关联点](Knowledge/金刚石项目路演关联点.md) +- [SMA2026路演素材包](Knowledge/SMA2026路演素材包.md) - [路演图片调用清单](Knowledge/路演图片调用清单.md) - [证据台账](Knowledge/证据台账.md) - [版本差异与待核对](Knowledge/版本差异与待核对.md) diff --git a/博志金钻项目基本情况_知识库/wiki/log.md b/博志金钻项目基本情况_知识库/wiki/log.md index a61ea99..ae6ce54 100644 --- a/博志金钻项目基本情况_知识库/wiki/log.md +++ b/博志金钻项目基本情况_知识库/wiki/log.md @@ -7,3 +7,11 @@ - Evidence records: 见 `Knowledge/证据台账.md`。 - Images: 抽取 70 张幻灯片预览图、419 张 PPT 内嵌图片,并生成 4 张 contact sheet。 - Open questions: 2024 与 2025 材料中的研发人数、专利数量等指标存在口径差异,已写入 `Knowledge/版本差异与待核对.md`。 + +## [2026-06-11] promote | SMA 2026 路演素材包 + +- Source: `Samsung Mobile Advance 2026` 活动信息、公司基本情况 PPT 知识库。 +- Pages updated: `Knowledge/SMA2026路演素材包.md`, `index.md`。 +- Evidence records: 引用既有 `Knowledge/公司身份与定位.md`, `Knowledge/产品与技术管线.md`, `Knowledge/制造能力与设备.md`, `Knowledge/金刚石项目路演关联点.md`。 +- Images: 汇总 SMA 路演可快速调用的 Image ID,实际调用仍需读取 `Sources/Images/图片快速索引.md`。 +- Open questions: 正式提交前需统一公司人数、专利数量、估值等对外口径。 diff --git a/金刚石散热项目/SMA2026英文路演PPT大纲.md b/金刚石散热项目/SMA2026英文路演PPT大纲.md new file mode 100644 index 0000000..1c45a31 --- /dev/null +++ b/金刚石散热项目/SMA2026英文路演PPT大纲.md @@ -0,0 +1,181 @@ +# SMA 2026 English Pitch Deck Outline + +Deck goal: answer Samsung Mobile Advance judges' three questions: innovation, value added to Samsung Mobile devices, and feasibility for a six-month PoC. + +Recommended title: + +> Diamond Thermal Packaging Modules for Next-Generation Samsung Mobile Devices + +## Slide 1 - Cover + +Title: Diamond Thermal Packaging Modules for Next-Generation Samsung Mobile Devices + +Subtitle: Samsung Mobile Advance 2026 Proposal + +Company: Suzhou Bozhi Jinzhuan Technology Co., Ltd. + +Key message: from diamond thermal materials to evaluable mobile packaging modules. + +## Slide 2 - Project Summary + +Headline: A six-month PoC to validate diamond-based heat spreading in mobile-scale packages. + +Copy: + +- We propose to develop and test diamond-based thermal packaging modules for Samsung Mobile use cases. +- The PoC will compare diamond heat spreaders / diamond-metal submounts against conventional thermal stacks under Samsung-defined power and space constraints. +- Final deliverables include sample coupons, one demonstrator module, test report, and integration recommendations. + +## Slide 3 - Why Samsung Mobile + +Headline: Mobile innovation is becoming thermally constrained. + +Three pressure points: + +- On-device AI and high-performance APs increase local heat flux. +- XR, foldable and wearable devices have tighter thermal envelopes. +- Camera, sensor and optical modules require thermal stability, not only peak cooling. + +SMA fit: AI, XR, Power, Foldable, Wearable, Sensors, Materials. + +## Slide 4 - Company Information + +Headline: Bozhi Jinzhuan is an engineering platform for high-power chip thermal packaging. + +Content: + +- Core products: high-power chip thermal-management packaging materials and devices. +- Existing products: ceramic substrates, TEC substrates, laser heat sinks, optical communication substrates, sensor substrates. +- Manufacturing base: over 10,000 m2 manufacturing footprint in company PPT materials. +- Capability foundation: PVD, photolithography, electroplating, electroless plating, cutting, inspection and customized metallization. + +Note: employee count, valuation and patent quantity should be checked before final external submission. + +## Slide 5 - Core Technology + +Headline: The differentiator is the interface, not only the diamond. + +Technology stack: + +- Diamond / diamond-metal thermal material. +- Surface activation and metallization. +- Fine patterning and solder / bonding compatibility. +- Package-level thermal path design. +- Inspection, reliability feedback and manufacturing iteration. + +Evidence points: + +- Company process capability includes Ti, TiW, Ni, Cr, Cu, Ta sputtering and 5/5 um line/space capability in existing audit materials. +- Diamond project materials include ultra-thin CVD diamond heat spreaders, diamond package substrates and diamond-Cu composites. + +## Slide 6 - Two Integration Routes + +Headline: Two practical routes for Samsung Mobile evaluation. + +Route A: Chip-on-diamond submount + +- Die mounted directly on metallized diamond substrate. +- Suitable for XR optical engine, laser/LED, sensor, high-power module. +- Output: diamond submount coupon with solder / wire-bond compatible metallization. + +Route B: Diamond layer in thermal stack + +- Diamond heat spreader or diamond-metal insert integrated with VC / graphite / TEC / frame. +- Suitable for AP hotspot, foldable thin space, wearable module. +- Output: thin diamond heat spreader coupon and stack-level thermal comparison. + +## Slide 7 - Performance Evidence + +Headline: Literature and company data support a measurable thermal advantage. + +Evidence: + +- Company material deck: ultra-thin CVD diamond heat spreader, 10-150 um, 1600-1800 W/mK. +- Company material deck: diamond-Cu composite, 600-800 W/mK, CTE 5-10 ppm/K. +- Literature MD: copper-coated diamond/Cu SLM composite reached 336 W/mK at 1 vol.% diamond under optimized process parameters. +- Literature MD: CVD diamond membranes can be fabricated as free-standing films; metallization literature supports Au/Cr and Au/Ti-W adhesion and thermal processing windows. +- Competitor benchmark: Element Six and Coherent both position CVD diamond as a high-power-density thermal management solution. + +Use caution: literature values are not direct product guarantees; they support feasibility and mechanism. + +## Slide 8 - Application Concept 1: AP / XR Hotspot + +Headline: Reduce hotspot temperature without redesigning the entire thermal system. + +Scenario: + +- Heat source: AP / NPU / XR optical module. +- Constraint: high heat flux, thin z-height, limited spreading area. +- Proposed module: metallized diamond heat spreader or diamond-Cu submount between die and existing spreader path. + +PoC test: + +- Dummy heat source + baseline stack vs. diamond stack. +- Metrics: peak temperature, thermal resistance, temperature uniformity, thickness impact. + +## Slide 9 - Application Concept 2: Sensor / Camera / Wearable Stability + +Headline: Use high thermal conductivity and low CTE to stabilize precision modules. + +Scenario: + +- Heat from AP, PMIC, camera actuator, optical engine or sensor readout can cause drift. +- Diamond package substrate can spread heat and reduce local thermal gradients. + +PoC test: + +- Sensor/camera module coupon with controlled heat source. +- Metrics: thermal gradient, module drift proxy, thermal cycling stability, warpage. + +## Slide 10 - Six-Month PoC Plan + +Headline: A focused PoC with three evaluation gates. + +Milestone 1 - Detailed design and review, Month 1-2: + +- Select target use case with Samsung. +- Freeze baseline stack, dimensions, power condition and KPIs. +- Deliver design review package and sample drawings. + +Milestone 2 - Sample build and early working review, Month 3-4: + +- Fabricate diamond heat spreader / submount coupons. +- Complete metallization and assembly process. +- Deliver early thermal data and manufacturability feedback. + +Milestone 3 - TRL 6/7 demonstrator, Month 5-6: + +- Build final demonstrator module. +- Complete thermal comparison, reliability screening and integration recommendation. +- Demo to Samsung Mobile R&D. + +## Slide 11 - Samsung Support Required + +Headline: We can execute independently; Samsung input will sharpen the PoC. + +Requested support: + +- Target device class and rough mechanical envelope. +- Baseline material stack or reference thermal coupon. +- Power map / heat source assumption for dummy testing. +- Loan device or non-confidential module sample if available. +- Review access to Samsung R&D thermal / packaging owner at milestone gates. + +No dependency: + +- The PoC can proceed with dummy heat source and public mechanical assumptions if device data is not available. + +## Slide 12 - Expected Value + +Headline: A fast path to evaluate diamond thermal packaging inside Samsung Mobile architectures. + +Expected Samsung value: + +- Lower hotspot temperature under thin, constrained mobile geometry. +- More stable sensors, optical modules and power modules. +- New material option for AI, XR, foldable and wearable roadmaps. +- A manufacturing partner able to move from material coupons to package-level modules. + +Closing: + +> We invite Samsung Mobile to co-define the most valuable target module and validate diamond thermal packaging within a six-month PoC. diff --git a/金刚石散热项目/[SMA 2026] Samsung Mobile Advance 2026-介绍.jpg b/金刚石散热项目/[SMA 2026] Samsung Mobile Advance 2026-介绍.jpg new file mode 100644 index 0000000..ede9633 Binary files /dev/null and b/金刚石散热项目/[SMA 2026] Samsung Mobile Advance 2026-介绍.jpg differ diff --git a/金刚石散热项目/[SMA 2026] 信息.pptx b/金刚石散热项目/[SMA 2026] 信息.pptx new file mode 100644 index 0000000..15912c7 Binary files /dev/null and b/金刚石散热项目/[SMA 2026] 信息.pptx differ diff --git a/金刚石散热项目/[SMA 2026] 路演模板.pptx b/金刚石散热项目/[SMA 2026] 路演模板.pptx new file mode 100644 index 0000000..5644250 Binary files /dev/null and b/金刚石散热项目/[SMA 2026] 路演模板.pptx differ diff --git a/金刚石散热项目/※苏州博志金钻-金刚石专项介绍.pdf b/金刚石散热项目/※苏州博志金钻-金刚石专项介绍.pdf new file mode 100644 index 0000000..44f1a9a Binary files /dev/null and b/金刚石散热项目/※苏州博志金钻-金刚石专项介绍.pdf differ diff --git a/金刚石散热项目/前期沟通.md b/金刚石散热项目/前期沟通.md new file mode 100644 index 0000000..997766a --- /dev/null +++ b/金刚石散热项目/前期沟通.md @@ -0,0 +1,145 @@ +说话人1 +呃,就是你结论里面的第一条。 +说话人1 +和。第二条就是你看你第二条的那个膜层结构,不是有钛铂金嘛,对不对?钛、铂金和泰乌的这些这些。呃,金属化体系。然后咱们自己的。咱们自己的钛伯金和泰头镍巴金,就是在我发给你的那个专项的介绍里面是不一样的。就是你钛伯金是用到一些更尺寸更薄、可以柔性的。呃,那个就叫超薄的那个器件里面。我给你截图,你能听见我说话吗?啊,我给你截图。 +说话人1 +你看我给你截图的这两个。就是,它都是。呃,做的金属化,但是它的方案不一样,然后衬底也不一样,然后用的场景也不一样。你看,第二个,它那个一般是我们做的是太铜镍巴金的方案,就是它那个铜会稍微厚一点,所以它耐电压电流会稍微高一点。然后尺寸会比较大。哦,它的机体呢,就是生长出来的比较厚的,相对比较厚啊,比如说0.3μm的。这个金刚石的再办,但是像第一个的这种。它就是它的,首先它的机体可以是我刚刚说的那种0.3,应该是0.33厚度的,然后也可以是那种直接在硅基上面生长出来的金刚石。 +说话人1 +然后他现在做那种可撕拉的,它那种。厚度应该是只有0.1μm左右吧,就反正撕下来就像一张纸一样的那种。薄膜,所以它叫超薄的。然后它的金属化方案是啊,钛铂金的方案。 +说话人1 +对钛铂金的话,就没有铜嘛,所以它耐电流就是比较能力比较小,但是它尺寸可以做到很小,然后呃,线路的精细程度会很高。 +说话人1 +所以这两个要区分一下,这两个要区分一下。 +说话人1 +这是我让你查的呀。就是。 +说话人1 +对对呀对呀,他不是给了一个proposal的模板吗?对不对?然后他还给你列举了一些应用场景。 +说话人1 +你发我的啥哦?我看一下啊。给我两分钟。 +说话人1 +王硕,我觉得你这个写的。 +说话人1 +反方向是对的呀。你这是。喂。哈喽,能听见吗?你这方向是对的,你这是check GPT出来的,还是你搜论文搜出来的? +说话人1 +哦,有有有有那个呃,数据支撑吗?有实验支撑吗? +说话人1 +嗯。 +说话人1 00:00:00 +喂,师姐,可以听到,不好意思,有点耽误耽搁了。 +说话人2 00:00:04 +没事儿没事儿没事儿。我给你发个图,你不要发给我,发给其他人。 +说话人1 00:00:16 +我看一下。 +说话人2 00:00:18 +没事儿,我给你发个图,你不要发给其他人。就是这个,是我们给一家客户做的,我们的产品在他们里面的用的。位置就是我红色圈起来的那个,你能看见吗? +说话人1 00:00:31 +对,了解。 +说话人2 00:00:35 +其实我的理想的就是这个proposal,我理想的就是能五九确认到我们在它各个的使用场景里面,它这个模组。就是它不是有好几个使用的场景嘛,大的使用场景。就他自己的产品类型。什么头衔?手那个耳手机指环耳机什么的,就它这个这几个大的应用场景里面它的。 +就是可能相关的。模块里面,它里面的组成是什么,然后我们能给它做什么的优化,就是我的理想的状态是这样,但是我觉得可咱可能也做不到,因为不知道它里面的内部结构是啥。 +说话人1 00:01:20 +对反正就是说是我们有类似的这种就是供货,或者说这种。 +说话人2 00:01:27 +所以你就你先尽可能的去查一下,如果查不到的话。我来如果查不到的话,就proposal里面他不是说要需要怎么支持吗?就给他问一下。 +说话人1 00:01:40 +行,那我们就直接是说,那直接就是说是我们这个可以是在手机这个,比如说这个模块散热之类的这边使用。然后就相当于这种的话,可以画一个用一个爆炸图或者说其他的形式表示爆炸图是啥,就就相当于是一个就是把它层层层级分开嘛,就相当于比如说一个棚,是咱们三热芯片下面他们自己的芯片就就就就就也行。 +说话人2 00:02:10 +我刚给你发的这几个就是我觉得你整个的路线你可以分为两种,就是第一种是芯片直接贴在咱们上面的,就是我发给你的这个图,你能看到它直接从咱们的甲板上面打线出来了吗?那个是个VC的芯片的。 +七,那俩小黄色的,中间那俩小黄色的,就是咱给他整的仔吧,明白? +说话人1 00:02:39 +我姐。行那之后反正就是做一下这个PPT。 +说话人2 00:02:46 +算是就是就是这个是第一条路线,就是它那个芯片直接在我们上面用。然后第二条路线是我们作为其他器件的组件,就是我们自己不生产。比如说那个T,你也了解的吧。一看。直接用到头衔里面,然后我们作为tec的再版,这是两条两两个大的路线。 +这个是单晶或者是多晶金刚石上面覆铜,然后还有那个金刚石铜的就是这个方案现在比较成熟了。就是手机上面的散热,就是手机的SOC的就是手机的上半部分是个SOC,然后在中间是个VC的均热板,然后把热都从VC均热板均均过来,然后手机的底部是一个T去做导热,这个方案其实算比较成熟了。 +然后还有一个。荣耀的方案是,它的那个摄像的模组是可移动式的。我等会儿发你一个视频号,它那个可移动下面的TC也是要用金刚石的。 +说话人1 00:03:56 +可移动设备的了解。 +说话人3 00:04:00 +对。 +说话人2 00:04:03 +然后他他那个AR眼镜儿。眼镜上面的那个投影的那个模组也是要用的,就是它在正中间的。那个应该就是一个激光器芯片。然后。然后,直接是我们的再板。然后,之前是氮化铝的,现在可以换成金刚石的。 +就是,这是几个例子,这是我知道的几个例子。 +说话人1 00:04:33 +所以你得找一下其他的有没有案例之类的。我找找发给你行没没没问题,明天我也找一下案例,然后之后就主要是就相当于看他之前的话应用领域。 +说话人2 00:04:50 +是的。 +说话人1 00:04:57 +然后是最后那个PPT的话。 +说话人2 00:04:58 +它是一些还有一些场景,我也不太熟的。就是远志之前跟其他几家聊的情况,我也没问,就这几个是我知道的。我发给你行。 +说话人1 00:05:11 +让我想一想后续。 +说话人2 00:05:15 +你先看着弄。 +说话人1 00:05:17 +行,那最后是这个,还是以英文的模式来去呈现。 +说话人2 00:05:20 +肯定是英文。然后。 +说话人3 00:05:23 +我想的是。 +说话人2 00:05:28 +就是第一个要呈现我刚跟你说的,我们的产品在它们结构里面是什么位置。然后,第二个是给它做一些优化,性能提升,能提升多少。但是,但是这个可能需要一些热仿真。热仿真的话,你可以找那个谁,就是陶老师他们下面那个叫啥来着,树峰之前给我们做的。 +然后,但是我们没有他们的实际的应用环境,好像也没办法做热仿真。 +说话人1 00:05:57 +一看多拟没有,就算是就是他做一个模拟。 +说话人2 00:06:01 +对之前热仿真出来的结果还是挺好的,就是跟真实的比较接近。所以我觉得最好能给一个热仿真的结果。 +说话人1 00:06:09 +我觉得这个仿真的话,主要就相当于是你只要有图就行嘛,就相当于图好看一点,就相当于是唬一唬人,那其实就是就差不多了。 +说话人2 00:06:19 +我们也可以假设一些条件嘛,就是反正先提过去,然后让他们去改一一件。 +说话人1 00:06:24 +我们再就是相当于是他们是不是已经有这些热仿真的一些内容了?反正我觉得要不然就直接是迁移过来,或者说直接说就是现在是不是博这边已经有一些热方针的结果了?是还是说这些结果,因为是其他的现有产品是不能用。 +说话人2 00:06:42 +一肯定是能用的,这结果你得让远志发你。我没有结果。 +说话人1 00:06:50 +了解了行,那我这边反正先先留着他上次做的是哪个产品呢? +说话人2 00:06:56 +我不记得了。你你你你你你你你要不先找找看有没有其他的。 +说话人1 00:07:02 +就是类似的模式图或者说之类的,反正先放着,然后呢。 +说话人2 00:07:05 +现在有个示意对行好的。三六反正这两种吧,就有就就这两方面吧。如果能有的话,最好,如果没有的话,那只能反正我打算先20多号草稿提出来,然后去他们南京的技术总部去聊一下。但是我之前看见他们南京的技术应该是sourcing,他们还是技术可能在韩国那边,所以我不知道20多号能不能聊出个什么东西来。 +说话人1 00:07:37 +我感觉就是这就是这个上面的话,主要就相当于是希望能成为他们这个供应链的一环,算是可能这个钱不钱的话就是无所谓是吧?对对了解了。好呀。行我了解了,然后之后的话,我这边是把这P1T1弄。 +行,然后那之后,那我记得之前不是有给苹果的那边的一个,那个相当于相关的一些内容嘛。 +说话人2 00:08:07 +就是相当于也是是我发你的这个图就是他们的。 +说话人1 00:08:12 +我我我知道我知道,然后。 +说话人2 00:08:14 +然后,你如果有那个资料的话,你也可以参考。 +说话人1 00:08:18 +对,我就相当于把那个反正这两个材材料全都综合一下,然后就来去处理一下。 +说话人2 00:08:27 +好。 +说话人1 00:08:28 +然后就这个图标出去,我肯定没问题。就是相当于相比于其他的这种的话,它的主要的这个特色在于就是咱们的工艺上面的话,就相当于是可以做一些后处理,还是说是相当于通孔,还是就是哪哪一部分和就是相当于现有的这种,就是首先,咱们不是做这个纯CBD。 +咱们CBD的话,只是咱们上游CBD金装石的话,只是咱们上游。然后的话呢,咱们是相当于是相比于是这种是纯切割或者说之类的,咱们的话就相当于增加的步骤就相当于是定制化还有表面改性,然后的话就相当于是就就就是这这这两点是吧,还有哪些就是比如说可以提供的服务。 +或者说特色在哪儿,就相当于和其他的这种是除了这种CBD厂商。 +说话人2 00:09:15 +你那个总结上写的是对的,你就按你总结的来写。 +说话人1 00:09:18 +行行没问题。 +说话人2 00:09:19 +了解了通孔也可以,但是现在金刚石通孔的很少。 +说话人1 00:09:25 +对,因为毕毕竟硬嘛,对行。 +说话人2 00:09:29 +而且他应该也用不到瞳孔的。 +说话人1 00:09:32 +因为它整体都是导电的是吧? +说话人2 00:09:37 +反正现在没预料过。 +说话人1 00:09:39 +明白了。行,辛苦学姐了。好。 +说话人2 00:09:41 +那我这边如果有这一方面资料的话,你也可以发来给我看一下。 +说话人1 00:09:45 +好嘞,辛苦。 +说话人2 00:09:47 +好好好。 +说话人1 00:09:49 +谢谢,OK, +说话人2 00:09:50 +拜拜。 \ No newline at end of file diff --git a/金刚石散热项目/前期调研.md b/金刚石散热项目/前期调研.md new file mode 100644 index 0000000..5652bb7 --- /dev/null +++ b/金刚石散热项目/前期调研.md @@ -0,0 +1,85 @@ +# 金刚石散热项目 - 前期调研补充 + +更新时间:2026-06-11 + +## 1. 面向 SMA 2026 的结论 + +Samsung Mobile Advance 2026 要求提交 6 个月 PoC proposal,重点看项目对 Samsung Mobile 产品的潜在影响,并明确关注 AI、XR、Power、Health、Camera、Audio、Foldable、Wearable、Sensors、Materials、Connectivity、Security 等方向。 + +对本项目而言,建议主线不是“我们能做金刚石材料”,而是: + +> Bozhi Jinzhuan provides diamond-based thermal packaging modules that can be evaluated in Samsung mobile devices, combining ultra-high thermal conductivity materials with metallization, patterning and module-level integration. + +评审关心的三件事应逐页回答: + +1. Innovation and uniqueness: 金刚石材料 + 表面金属化 + 图形化/封装集成能力,而不是单纯热导率指标。 +2. Value added to Samsung Mobile devices: 降低热点温度、提高热稳定性、缩小散热路径厚度、提升高功率模块可靠性。 +3. Feasibility to integrate: 6 个月内交付可测试样件、结构建议、热测试/仿真对比和 Demo package,而不是开放式研发。 + +官方信息入口: + +## 2. 应用场景优先级 + +| 优先级 | 场景 | 对应 SMA 领域 | 金刚石方案切入点 | PoC 可交付 | +|---|---|---|---|---| +| P0 | On-device AI / AP hotspot | AI, Power, Materials | SoC 顶部或近芯片热扩散层,减少局部热点 | Dummy chip + diamond heat spreader thermal comparison | +| P0 | Galaxy XR / AR optical engine | XR, Power, Wearable | 激光/微投影芯片 submount,以金刚石替代 AlN/Al2O3 | Laser/LED submount sample + junction temperature comparison | +| P1 | Foldable thin thermal path | Foldable, Materials | 薄型/异形金刚石散热片用于折叠结构有限空间 | Thin/shaped heat spreader coupon + bending/assembly review | +| P1 | Camera movable module / high-power imaging | Camera, Sensors | 模组热漂移控制,提高传感器/执行器稳定性 | Module-level thermal stability mock-up | +| P1 | Wearable / health sensor module | Wearable, Health, Sensors | 低热膨胀、高导热载板降低传感器基准漂移 | Sensor substrate coupon + thermal cycling data | +| P2 | Fast charging / PMIC module | Power | 高热流电源管理模块热沉或封装载板 | PMIC dummy package thermal resistance test | + +领导沟通中提到的两条路线可以作为 PPT 的方案骨架: + +- Route A: chip directly mounted on diamond substrate / diamond-metallized submount. +- Route B: diamond component as part of another thermal module, such as TEC/VC/thermal spreader stack. + +## 3. 竞品与供应链地图 + +| 类型 | 公司 / 产品 | 已知公开信息 | 对我们的启发 | +|---|---|---|---| +| CVD diamond heat spreader | Element Six | 官方称 CVD diamond heat spreaders 用于更高功率密度,热导率可按需求提供 1000-2200 W/(mK),应用包括 RF、GPU、HPC、AI accelerators 等。 | 国际头部强调“系统热瓶颈”和“可定制热要求”;我们要避开纯材料规模竞争,突出移动端封装集成。 | +| CVD diamond substrate | Coherent | 官方 CVD diamond substrate 页面强调高热导、抗热冲击等特性;2026-01-19 发布 bondable diamond thermal management,主张通过直接键合大幅降低界面热阻。 | “界面热阻”已经成为国际产品叙事重点;我们的 PPT 应把金属化、表面处理、键合/焊接作为差异化,而不是只说 TC。 | +| CVD diamond / Cu-diamond heat spreader | A.L.M.T. / Sumitomo Electric | 官方列出 CVD-Diamond 厚度 0.2-0.4 mm,热导率 >1000 W/(mK),CTE 2.3 ppm/K;Cu-Diamond 用于半导体激光 submount、功率晶体管基板等。 | 竞品已把 CVD diamond 和 Cu-diamond 都产品化,应用多在激光、功率器件;三星移动端 PoC应强调轻薄小型化和模组适配。 | +| Single crystal diamond submount | Sumitomo Electric | 官方称单晶金刚石可用于 semiconductor laser submount,并可形成各类金属化薄膜用于 chip mounting 和 wire bonding。 | 金属化薄膜 + 芯片安装/打线是成熟叙事,可迁移到 XR 光机、摄像/传感器高稳定模组。 | +| Diamond metal composites | Hi-Rel / DiaCool | Hi-Rel 展示 DiaCool-Ag700、Ag850、Al500、Cu600 等 diamond metal composite thermal spreaders;其 WDAM 合作页面称组合覆盖 500 W/mK 到 2200 W/mK 以上。 | 竞品已经按材料系列包装解决方案;我们需要把“可加工、可金属化、可交付样件”讲清楚。 | +| Ag-Dia heat spreader | Sumitomo Electric U.S.A. | 官方称 Ag-Dia 通过银与金刚石粉末专有烧结工艺实现 600 W/mK 或更高热导率,可用于更大尺寸半导体器件散热基板。 | 金刚石-金属复合路线商业化空间明确;我们可把金刚石铜/铝/银作为材料平台而非单一产品。 | + +## 4. 博志金钻差异化表述 + +不建议说: + +- “首次提出金刚石用于芯片散热。” +- “单靠金刚石材料即可解决三星终端散热。” +- “已确定可进入三星某具体产品结构。” + +建议说: + +- We focus on the last-mile integration of diamond thermal materials into package-level modules. +- Our differentiation is not only thermal conductivity, but interface engineering, metallization, patterning, solder/bonding compatibility and manufacturable samples. +- The PoC will validate where a diamond layer or diamond-metal submount creates measurable thermal value in Samsung mobile architectures. + +可强调的公司能力: + +- CVD / diamond thin film route: 10-150 um 超薄金刚石散热片,1600-1800 W/mK 公司材料口径。 +- Diamond package substrate route: Ti/Pt/Au thin film for fine features; Ti/Cu/Ni/Au thick film for higher current/voltage packaging. +- Diamond-Cu composite route: 600-800 W/mK、CTE 5-10 ppm/K 公司材料口径。 +- Process integration: PVD, photolithography, electroplating, electroless plating, cutting, inspection, customized metallization. + +## 5. 建议 PoC KPI + +| KPI 类别 | 建议指标 | 说明 | +|---|---|---| +| Thermal | hotspot temperature drop vs. baseline | 以 AlN / Cu / graphite / VC reference coupon 作对照,目标可先写 “target: measurable reduction under Samsung-defined test condition”。 | +| Interface | thermal resistance across bonding / metallization stack | 强调不是裸材料 TC,而是堆叠后的系统热阻。 | +| Mechanical | thickness, flatness, warpage, dicing edge quality | 对移动端轻薄空间很关键。 | +| Electrical / packaging | metallization adhesion, solderability, wire bonding compatibility | 对芯片直贴和模块集成重要。 | +| Reliability | thermal cycling, humidity/heat aging, power cycling | 可作为后续 Samsung feedback 后确认的测试包。 | +| Deliverable | sample coupons + one demo module + test report | 对应 6 个月 PoC,更容易评估。 | + +## 6. 仍需补充 + +- Samsung 具体产品结构需通过公开 teardown / 专利 / 产业链资料继续核对,不能在 proposal 中假设其内部结构。 +- 如能拿到远志/陶老师过往热仿真结果,可放入一页“validated simulation workflow”,比临时虚构仿真更可信。 +- 需要统一公司英文名、联系人、专利数量、员工人数、估值等外部口径。 +- 若提交前允许与 Samsung Open Innovation 团队沟通,应把两个 PoC 方向拆成可选项,让对方选择最感兴趣的目标产品:AP hotspot / XR optical engine / wearable sensor / charging PMIC。 diff --git a/金刚石散热项目/项目前期工作-整理.md b/金刚石散热项目/项目前期工作-整理.md new file mode 100644 index 0000000..f033205 --- /dev/null +++ b/金刚石散热项目/项目前期工作-整理.md @@ -0,0 +1,105 @@ +# 苏州博志金钻金刚石专项0前期合作单位共同成果产出介绍 + +## 1. 结论 + +主要覆盖以下三条主线: + +1. **CVD 金刚石薄膜/散热片路线**:在硅或其他基底上制备大面积 CVD 多晶金刚石,随后形成独立薄膜、热扩散片或热管理基板。这与 PDF 第 4-5 页“CVD 长晶、晶圆切割、剥离、金属化、超薄金刚石散热片/薄膜器件”高度相似。 +2. **金属化金刚石载板/热沉路线**:CVD 金刚石或金刚石片经过 Ti/Pt/Au、Au/Cr、AuSn、Ti-W 等金属化体系,用于激光器、InP/GaN/高功率器件封装和键合。这与 PDF 第 5-6 页“金属功能层、薄膜电阻、AuSn 共晶焊料、通孔、金属化、封装载板”高度相似。 +3. **改性金刚石粉末到 Cu/diamond 复合热沉路线**:Ti、W、Cr 等界面层或活性元素改善金刚石与铜的润湿/界面热阻,再通过气压浸渗、粉末冶金、烧结、激光选区熔化等工艺制备高导热低膨胀 Cu/diamond 复合材料。这与 PDF 第 7-8 页“磁控溅射改性金刚石粉末、TIM 增强相、金刚石铜复合材料、液相烧结”高度相似。 + +因此,**该专项路线有明确的文献基础**。如果用于 PoC 论证,建议把创新点放在“大面积/超薄/可剥离量产能力、金属化与器件级封装集成、特定应用场景验证、供应链和制造良率”上,而不是把 CVD 金刚石热管理或 Cu/diamond 复合材料本身表述为全新概念。 + +## 2. PDF 中的技术路线摘要 + +从 10 页 PDF 可抽取出的核心路线如下: + +| PDF 页码 | 产品/路线 | 关键工艺 | 关键指标/用途 | +|---|---|---|---| +| 第 4 页 | 超薄金刚石散热片 | CVD 在硅基体异质外延多晶金刚石,晶圆切割,剥离,表面处理,Ti/Pt/Au 金属化 | 6/8 英寸,10-150 um,热导率 1600-1800 W/mK,用于有限空间和特殊工况芯片散热 | +| 第 5 页 | 超薄金刚石薄膜器件 | 硅-金刚石衬底上制备导电层、薄膜电阻、AuSn 焊料;剥离切割;或硅面加工芯片、金刚石面做通孔与金属化 | 2/4/6 英寸,10-150 um,TaN 50/100 ohm,AuSn 75/25,共晶焊 | +| 第 6 页 | 单/多晶金刚石封装载板 | 大尺寸金刚石片金属化、光刻、通孔垂直互联、薄膜电阻集成、金锡焊料预制 | 2/3 英寸,Ti/Pt/Au 薄膜,TiCuNiAu 厚膜,面向超大功率芯片 | +| 第 7 页 | 改性金刚石粉末 | 磁控溅射 Ti/Cu/Cr/W 表面包覆 | 用作 TIM 增强相,或用于金刚石-金属复合材料,改善润湿性并降低热阻 | +| 第 8 页 | 金刚石-铜复合材料 | 铜熔化、金刚石颗粒重排、烧结/表面处理、异形加工 | 热导率 600-800 W/mK,CTE 5-10 ppm/K,适用于算力模块、射频、大功率激光器等 | + +## 3. Google Scholar 检索式与返回方向 + +本轮使用的 Scholar 检索式如下。Scholar 页面可返回结果标题,但最终论文元数据以 DOI/Crossref/出版社或开放页面核对。 + +| 检索式 | 返回的代表性标题方向 | +|---|---| +| `CVD diamond heat spreader silicon substrate lift off thermal management` | 大面积 CVD 金刚石热管理基板、CVD 金刚石热扩散片、GaN 热点热管理 | +| `metallized CVD diamond substrate AuSn solder heat spreader laser diode` | InP 激光器到 CVD 金刚石热沉键合、CVD 金刚石 submount 金属化、AuSn/Cr/W 金属化热沉 | +| `diamond copper composite Ti coated diamond particles gas pressure infiltration thermal conductivity` | Ti 涂覆金刚石颗粒、气压浸渗 Cu/diamond、Cu-Ti/diamond 复合材料 | +| `magnetron sputtering tungsten coating diamond particles copper composite thermal conductivity` | W 磁控溅射涂覆金刚石颗粒、W 界面层提升 Cu/diamond 热导率 | +| `copper diamond composite carbide forming additives W Mo Cr Ti thermal conductivity` | W/Mo/Cr/Ti 碳化物形成元素对 Cu/diamond 界面和热导率影响 | +| `diamond copper composite heat sink high power electronics thermal management` | Cu/diamond 热沉、SiC 功率器件微通道热沉、先进热管理材料 | + +## 4. 重点论文对照表 + +匹配等级说明: + +- **高度同路**:工艺链或器件目标与 PDF 某一产品线直接重合。 +- **相邻支撑**:研究对象不同,但支撑 PDF 的材料选择、界面设计或应用论证。 +- **背景综述**:适合写行业/技术基础,不宜单独作为工艺可行性证据。 + +| 序号 | 匹配等级 | 对应 PDF 路线 | 论文 | DOI/来源 | 与 PDF 的关系 | +|---:|---|---|---|---|---| +| 1 | 背景综述 | CVD 金刚石热扩散片/GaN 高功率散热 | Liwen Sang. **Diamond as the heat spreader for the thermal dissipation of GaN-based electronic devices**. *Functional Diamond*, 2021. | https://doi.org/10.1080/26941112.2021.1980356 | 综述金刚石作为 GaN 器件热扩散层/热沉的路线,可支撑 PDF 中“高功率芯片热管理封装材料”的应用动机。 | +| 2 | 高度同路 | CVD 金刚石薄膜热扩散 | M. Seelmann-Eggebert et al. **Heat-spreading diamond films for GaN-based high-power transistor devices**. *Diamond and Related Materials*, 2001. | https://doi.org/10.1016/S0925-9635(00)00562-8 | 直接讨论高功率 GaN 器件用金刚石热扩散薄膜,与 PDF 的 CVD 金刚石散热片/器件封装目标吻合。 | +| 3 | 高度同路 | 大面积 CVD 金刚石基板 | W. D. Brown et al. **State-of-the-art synthesis and post-deposition processing of large area CVD diamond substrates for thermal management**. *Surface and Coatings Technology*, 1996. | https://doi.org/10.1016/S0257-8972(96)03005-8 | 覆盖大面积 CVD 金刚石基板合成与后处理,直接对应 PDF 的 6/8 英寸 CVD 金刚石散热片制造路线。 | +| 4 | 高度同路 | 独立/剥离金刚石膜 | M. C. Salvadori et al. **Fabrication of free-standing diamond membranes**. *Thin Solid Films*, 1996. | https://doi.org/10.1016/S0040-6090(96)09010-4 | 直接对应 PDF 中“CVD 长晶后剥离金刚石薄膜”的独立膜路径。 | +| 5 | 相邻支撑 | 多层金刚石热扩散器 | K. Jagannadham. **Multilayer diamond heat spreaders for electronic power devices**. *Solid-State Electronics*, 1998. | https://doi.org/10.1016/S0038-1101(98)00216-0 | 支撑电子功率器件中多层金刚石热扩散结构,与 PDF 的超薄散热片和封装载板相邻。 | +| 6 | 高度同路 | 金刚石载板金属化 | Ilango Meyyappan et al. **Au/(Ti-W) and Au/Cr metallization of chemically vapor-deposited diamond substrates for multichip module applications**. *Thin Solid Films*, 1994. | https://doi.org/10.1016/0040-6090(94)90357-3 | 直接对应 PDF 的 Ti/Pt/Au、金属化、封装载板思路;虽金属体系不完全一致,但目的和路线一致。 | +| 7 | 高度同路 | 激光器与 CVD 金刚石热沉键合 | A. Katz et al. **Advanced metallization schemes for bonding of InP-based laser devices to CVD-diamond heatsinks**. *Materials Chemistry and Physics*, 1994. | https://doi.org/10.1016/0254-0584(94)90169-4 | 直接对应 PDF 的大功率激光器、金属化、AuSn/键合热沉应用方向。 | +| 8 | 相邻支撑 | 金属化可靠性 | C. D. Iacovangelo. **Thermal stability of metallized CVD diamond**. *Thin Solid Films*, 1996. | https://doi.org/10.1016/S0040-6090(95)08236-0 | 可支撑 PDF 中金属化金刚石载板可靠性论证,但不是完整器件路线。 | +| 9 | 高度同路 | 激光二极管金刚石 submount | E. E. Ashkinazi et al. **Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition**. *Quantum Electronics*, 2012. | https://doi.org/10.1070/QE2012V042N11ABEH015042 | 与 PDF 中“大功率激光器/光通信/高功率芯片载板”的应用高度贴近,强调 CVD 金刚石 submount 对输出功率的提升。 | +| 10 | 背景综述 | Cu/diamond 复合热管理材料 | S. Q. Jia and F. Yang. **High thermal conductive copper/diamond composites: state of the art**. *Journal of Materials Science*, 2021. | https://doi.org/10.1007/S10853-020-05443-3 | 可作为 Cu/diamond 复合材料路线综述,支撑 PDF 第 8 页金刚石铜复合热沉的技术背景。 | +| 11 | 高度同路 | Ti 涂覆金刚石颗粒 + Cu 复合 | Jianwei Li et al. **Microstructure and thermal conductivity of Cu/diamond composites with Ti-coated diamond particles produced by gas pressure infiltration**. *Journal of Alloys and Compounds*, 2015. | https://doi.org/10.1016/J.JALLCOM.2015.06.062 | 与 PDF 第 7 页 Ti 改性金刚石粉末和第 8 页 Cu/diamond 复合材料直接吻合。 | +| 12 | 高度同路 | W 磁控溅射涂覆金刚石颗粒 | Jinhao Jia et al. **Enhanced thermal conductivity in diamond/copper composites with tungsten coatings on diamond particles prepared by magnetron sputtering method**. *Materials Chemistry and Physics*, 2020. | https://doi.org/10.1016/J.MATCHEMPHYS.2020.123422 | 直接对应 PDF 的“磁控溅射 W 改性金刚石粉末”路径,是本轮最贴合第 7 页表述的论文之一。 | +| 13 | 高度同路 | Cu-Ti/diamond 气压浸渗 | Jianwei Li et al. **Optimized thermal properties in diamond particles reinforced copper-titanium matrix composites produced by gas pressure infiltration**. *Composites Part A*, 2016. | https://doi.org/10.1016/J.COMPOSITESA.2016.10.005 | 直接支撑 PDF 中“活性金属改善金刚石-金属界面、降低热阻”的机理。 | +| 14 | 高度同路 | 涂层金刚石 + 粉末冶金 Cu/diamond | Shubin Ren et al. **Effect of coating on the microstructure and thermal conductivities of diamond-Cu composites prepared by powder metallurgy**. *Composites Science and Technology*, 2011. | https://doi.org/10.1016/J.COMPSCITECH.2011.06.012 | 与 PDF 第 7-8 页“金刚石粉末改性后进入金刚石铜复合材料”的路线一致。 | +| 15 | 高度同路 | W/Mo/Cr/Ti 碳化物形成元素 | Arina V. Ukhina et al. **The Influence of the Carbide-Forming Metallic Additives (W, Mo, Cr, Ti) on the Microstructure and Thermal Conductivity of Copper-Diamond Composites**. *Journal of Composites Science*, 2023. | https://doi.org/10.3390/JCS7060219 | 与 PDF 写明的 Ti/Cu/Cr/W 改性元素高度相关,适合说明界面化合物层和热导率之间的关系。 | +| 16 | 相邻支撑 | Cu/diamond 界面工程 | Bin Xu et al. **Scalable monolayer-functionalized nanointerface for thermal conductivity enhancement in copper/diamond composite**. *Carbon*, 2021. | https://doi.org/10.1016/J.CARBON.2021.01.018 | 不同于 PDF 的金属溅射路线,但同样解决金刚石/Cu 界面热阻问题,可作为界面工程相邻路线。 | +| 17 | 相邻支撑 | Cu/diamond 热沉器件化 | Kangyong Li et al. **Diamond/Cu composites microchannel heat sink for effective thermal management of SiC power devices**. *Applied Thermal Engineering*, 2026. | https://doi.org/10.1016/J.APPLTHERMALENG.2025.129116 | 体现 Cu/diamond 从材料向功率器件热沉集成的近年方向,适合支撑 PDF 的 SiC/射频/算力模块等应用延展。 | +| 18 | 相邻支撑 | Ti/Cu 涂覆金刚石 + 成形工艺 | Lu Zhang et al. **Fabrication of Titanium and Copper-Coated Diamond/Copper Composites via Selective Laser Melting**. *Micromachines*, 2022. | https://doi.org/10.3390/MI13050724 | 与 PDF 的 Ti/Cu 涂覆金刚石粉末相近,但制备工艺为选择性激光熔化,不是 PDF 主要描述的液相烧结路径。 | + +## 5. 最值得优先阅读的 8 篇 + +如果时间有限,建议优先读下面 8 篇,因为它们最能支撑 PDF 的具体工艺链。 + +1. Brown et al., 1996:大面积 CVD 金刚石基板合成与后处理。 +2. Salvadori et al., 1996:独立金刚石膜制备。 +3. Seelmann-Eggebert et al., 2001:GaN 高功率器件金刚石热扩散薄膜。 +4. Meyyappan et al., 1994:CVD 金刚石载板 Au/Ti-W 和 Au/Cr 金属化。 +5. Katz et al., 1994:InP 激光器到 CVD 金刚石热沉的金属化键合。 +6. Li et al., 2015:Ti 涂覆金刚石颗粒 + Cu/diamond 气压浸渗。 +7. Jia et al., 2020:W 磁控溅射涂覆金刚石颗粒提升 Cu/diamond 热导率。 +8. Ukhina et al., 2023:W/Mo/Cr/Ti 等碳化物形成元素对 Cu/diamond 微结构和热导率的影响。 + +## 6. 可用于 PoC 材料的表述建议 + +可以这样表述: + +> 公开文献已经证明,CVD 金刚石薄膜/热扩散片、金属化金刚石载板、以及界面改性的 Cu/diamond 复合热沉,是高功率电子、光通信、激光器和功率半导体热管理中的成熟研究方向。本项目拟推进的重点不在于重新提出金刚石热管理概念,而在于面向量产封装场景,将大面积超薄金刚石膜制备、剥离和金属化工艺,与器件级焊接、薄膜电阻、通孔互联和金刚石铜复合热沉集成起来。 + +不建议这样表述: + +> 本项目首次提出金刚石用于芯片散热或金刚石铜复合热沉。 + +原因:以上文献已经覆盖这些基础概念和相当多的工艺路径。更稳妥的创新表述应落在规模化、良率、厚度控制、封装集成、客户应用验证、特殊结构和成本上。 + +## 7. 仍需补充的内容 + +1. **专利检索**:本轮只整理学术论文。若用于商业 PoC 或合作评估,应补充 Google Patents、CNIPA、WIPO 检索,重点看“CVD diamond heat spreader lift-off”“metallized diamond submount”“diamond copper composite sputtering coating”等关键词。 +2. **全文数据核对**:部分论文可能在 ScienceDirect 或 Taylor & Francis 付费墙后。本轮核对了题名、作者、期刊和 DOI,但没有逐篇下载全文。 +3. **指标对比表**:如果要进入技术尽调,建议新增一张表,把 PDF 的 1600-1800 W/mK、10-150 um、600-800 W/mK、5-10 ppm/K 等指标逐项对照论文中的实测值。 +4. **Samsung PoC 应用映射**:如果目标是三星电子资助计划,建议进一步把论文按应用端拆成光通信/激光器、GaN/SiC 功率器件、数据中心、RF 模块、先进封装五类。 + +## 8. 本轮核验边界 + +- 已读取本地 PDF 正文并抽取 10 页内容。 +- 已通过 Google Scholar 查询返回的标题判断论文方向。 +- 已用 DOI/Crossref 查询核对核心论文的题名、作者、年份和期刊。 +- 未声称完成系统性综述;本文件是“相同或类似技术路线”的第一版论文地图。 +- 未改动原始 PDF。 diff --git a/金刚石散热项目/项目前期工作/01_CVD金刚石薄膜与热扩散片/1_Diamond as the heat spreader for the thermal dissipation of GaN-based electronic devices.pdf b/金刚石散热项目/项目前期工作/01_CVD金刚石薄膜与热扩散片/1_Diamond as the heat spreader for the thermal dissipation of GaN-based electronic devices.pdf new file mode 100644 index 0000000..c288747 --- /dev/null +++ b/金刚石散热项目/项目前期工作/01_CVD金刚石薄膜与热扩散片/1_Diamond as the heat spreader for the thermal dissipation of GaN-based electronic devices.pdf @@ -0,0 +1,28776 @@ +%PDF-1.4 +% +1 0 obj +<< +/Type /Catalog +/Version /1.5 +/Pages 2 0 R +/PageLabels 3 0 R +/Metadata 4 0 R +/Names 5 0 R +/Outlines 6 0 R +>> +endobj +7 0 obj +<< +/Title (Diamond as the heat spreader for the thermal dissipation of GaN-based electronic devices) +>> +endobj +2 0 obj +<< +/Type /Pages +/Kids [8 0 R 9 0 R 10 0 R 11 0 R 12 0 R 13 0 R 14 0 R 15 0 R 16 0 R 17 0 R +18 0 R 19 0 R 20 0 R 21 0 R 22 0 R 23 0 R] +/Count 16 +>> +endobj +3 0 obj +<< +/Nums [0 24 0 R 1 25 0 R] +>> +endobj +4 0 obj +<< +/Length 4662 +/Subtype /XML +/Type /Metadata +>> +stream + + + + Adobe LiveCycle PDFG ES + + + + + + 2022-01-24T15:47:00+05:30 + Arbortext Advanced Print Publisher + 2022-01-24T15:47:00+05:30 + 2022-01-24T15:47:00+05:30 + + + uuid:8addd02c-2a8d-4824-a074-0375b52bef83 + uuid:9735c38c-787b-4f5b-aca7-e8422204bc96 + + + Journal + Functional Diamond + © 2021 The Author(s). 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dissipation of GaN-based electronic devices.md @@ -0,0 +1,382 @@ +# Diamond as the heat spreader for the thermal dissipation of GaN-based electronic devices + +Liwen Sang + +To cite this article: Liwen Sang (2021) Diamond as the heat spreader for the thermal dissipation of GaN-based electronic devices, Functional Diamond, 1:1, 174-188, DOI: 10.1080/26941112.2021.1980356 + +To link to this article: https://doi.org/10.1080/26941112.2021.1980356 + +![](images/091588d5a367d5de2e5d667682f3a64f8accc02ae7a155051e39bbea750840d6.jpg) + +![](images/e500f787007e33a5dfac9f2f520f05e3a3cb0bc3d48216361d848a98fb49f5d5.jpg) + +![](images/4157837989f5709dde5a42b00e1db0c7262982b89b1d98c2db996112962635df.jpg) + +![](images/12a84b5342af08ca08f805683eee0a3e0f9e949c7b5ee18a6337ecf3852977e8.jpg) + +![](images/977e6c88ccca94e35e1b33e98bdab5d006454147dcd3faecd210b023c4715886.jpg) + +![](images/740d21a21626750bdf408e923f7e312fb93ae36907f4358883cfcfdece817dde.jpg) + +![](images/f9a62662ddfdba4b7dc7293d2a7ab8de11f4587072b0722af0cd266f24178d82.jpg) + +© 2021 The Author(s). Published by Informa UK Limited, trading as Taylor & Francis Group, on behalf of Zhengzhou Research Institute for Abrasives & Grinding Co., Ltd. + +Published online: 15 Oct 2021. + +Submit your article to this journal + +Article views: 13833 + +View related articles + +View Crossmark data + +Citing articles: 63 View citing articles + +Review + +OPEN ACCESS + +C Check for updates + +# Diamond as the heat spreader for the thermal dissipation of GaN-based electronic devices + +Liwen Sanga,b + +a International Center for Materials Nanoarchitectonics (MANA ), National Institute for Materials Science (NIM S), Tsukuba, Ibaraki, Japan; bJST-PRESTO , The Japan Science and Technology Agency, Tokyo, Japan + +# ABSTRACT + +With the increasing power density and reduced size of the GaN-based electronic power converters, the heat dissipation in the devices becomes the key issue toward the real applications. Diamond, with the highest thermal conductivity among all the natural materials, is of the interest for integration with GaN to dissipate the generated heat from the channel of the AlGaN/ GaN high electron mobility transistors (HEMTs). Current techniques involve three strategies to fabricate the GaN-on-diamond wafers: bonding of GaN with diamond, epitaxial growth of diamond on GaN, and epitaxial growth of GaN on diamond. As a result of the large lattice mismatch and thermal mismatch, the integration of GaN-on-diamond wafer is suffered from stress, bow, crack, rough interfaces, and large thermal boundary resistance. The interfaces with transition or buffer layers impede the heat flow from the device channel and greatly influence the device performance. In this review, we summarize the three different techniques to achieve the GaN-on-diamond wafers for the fabrication of AlGaN/GaN HEMTs. The problems and challenges of each method are discussed. In addition, the effective thermal boundary resistance between GaN and diamond, which characterizes the heat concentration, is analyzed with regard to different integration and measurement methods. + +# ARTICLE HISTORY + +Received 13 June 2021 + +Accepted 26 August 2021 + +# KEYWORDS + +Semiconductor; + +Heat-related + +# 1. Introduction + +Benefitted from the high breakdown voltages (10 times higher than Si), high switching speed (over GHz), compact size, and tunable electronic architecture [1–7], III-V nitride semiconductor is becoming one of the best candidates for high-power electronics to enable the increasing power density and high conversion efficiency. The commercialized AlGaN/GaN high electron mobility transistors (HEMTs) have led to the entry into the mediumpower market, and play a central role for the RF and millimeter-wave applications [8–11]. In the applications of 5 G communications, radar, and electronic warfare, the HEMTs devices can offer more than 10 times higher power density than the existing Si technologies [12]. This giant power induces a huge amount of heat in the chip area, creating localized hot spots with fluxes above 10 kW/cm2 and package-level volumetric heat generation that can exceed 100 W/cm−3. The high-level power dissipation results in the challenges using conventional approaches for the thermal management. With the increased power density, self-heating inside the devices becomes an essential issue that accelerates the failure and poor reliability in the real application. Thermal + +dissipation through conventionally used approaches is no longer adequate. To achieve the effective thermal dissipation, the heat spreader with a much higher thermal conductivity is required. + +The current GaN wafers are typically grown on sapphie, silicon (Si), silicon carbide (SiC), or free-standing GaN substrates, whose thermal conductivities are 35, 150, 400, and 280 W/mK, respectively [13–15], which are far from the requirements. The ideal heat spreader would be a substrate that is both highly thermally conductive and electrically insulating. Diamond, with the thermal conductivity up to 2400 W/mK at room temperature for the single crystals, and approaching 2000 W/ mK for the polycrystals, is the best candidate as a heat spreader for GaN power transistors [16–19]. Early simulations and modelling showed that the passive thermal extraction by direct contact with diamond could dramatically reduce junction temperatures by 25-50% [20– 23]. However, as shown in Table 1, diamond and GaN exhibit widely mismatched properties, such as the crystalline structures, lattice constants and thermal expansion coefficients (TEC), making them challenging as bonded or growth pairs [24–28]. + +Over the past twenty years, a variety of methods have been developed to utilize diamond as the heat spreader for AlGaN/GaN power transistors. The developed wafer is therefore called “GaN-on-diamond wafer”. The approaches are summarized in Figure 1, which include: (1) bonding of diamond to GaN wafers or directly to the HEMT devices with/without an adhesion layer, (2) GaN epitaxial growth on single-crystal or poly-crystal diamond substrate, then fabrication of HEMT devices, and (3) nanocrystalline or poly-crystalline diamond growth on the frontside or backside of GaN or the HEMTs devices. For the three approaches, the thermal resistance + +Table 1. T he lattice constant and thermal expansion coefficient of GaN, sapphire, Si and diamond at room temperature. + +
SubstrateGaN (a-axis)SapphireSiDiamond
Lattice constant (Å)3.1894.7585.4203.567
TEC (×10-6K-1)5.64.5-5.82.6~1.1
+ +at the GaN/diamond, which is referred to the “effective thermal boundary resistance $( T B R _ { e f f } ) ^ { 3 }$ ”, is one of the factors that significantly increases the overall temperatures during device operation. Therefore, the optimizations on the integration technique and interface property are important. In this review, the state-of-the-art development on the GaN-on-diamond wafer and the modified fabrication process for HEMTs with regard to different integration methods will be presented. The strategies to improve the interfaces and reduce the $T B R _ { e f f }$ will be summarized and discussed. The device performances will be further shown with respect to different methods. + +# 2. Integration of the diamond to GaN or HEMTs through bonding technique + +The concept of harvesting the III-V epitaxial layers from one substrate and thermally bonding them to another + +![](images/b06ca0f4103825a09fcc534d1664d5353b720cbf178e159d2b7f8bbef425efb2.jpg) + +![](images/45ea2504a2588ba9cc4656d054853ecd4fb0c6cb802ab6ba010057ded1a97a34.jpg) +(a) Top:GaN wafer bonding to diamond then fabricate HEMT devices; Bottom:HEMTsdevices directly bondingto diamond substrate +(b) HEMT devices epitaxial grown on diamond substrate + +![](images/6c6cea8e3d8fb10dec6aa4a60c96f87b9d29a7263793b9e124d3d7c4b3f8109b.jpg) +(c) Top:Diamond grown on the top of HEMTs devices; Bottom:Diamond epitaxially grown on HEMTwafers +Figure 1. Fabrication of the GaN-on-diamond wafer for the HEMT devices. S: source, D: Drain, G: gate. + +substrate was proposed in the 1990s [29,30]. In 1997, Kelly et al. firstly demonstrated the lift-off of GaN film from sapphire substrate by illuminating the interface with a pulsed laser [31], which enabled the development of the GaN bonding technology. With the development of GaN-on-Si technology [32,33], the bonding process of GaN or HEMTs wafer became much easier due to the easy removal of the Si substrate. The advantage for the GaN bonding to diamond is that the crystal quality of both GaN and diamond can be guaranteed after the bonding. + +The formation of GaN-on-diamond wafer started with GaN or AlGaN/GaN HEMT epitaxial layers grown on a silicon substrate [34]. Other substrates, such as sapphire, silicon carbide, or aluminum nitride may also be used. In order to preserve the orientation of the epilayers, which is required to fabricate HEMTs on this structure, the epilayer was transferred twice (Figure 1 (a) top). The GaN-epilayer structure was firstly bonded to another sacrificial carrier. The growth substrate was then removed using either a wet chemical or dry etching process that is selective to GaN, leaving GaN epilayers flipped. An atomically flat dielectric layer was then deposited to bond with the diamond. The sacrificial carrier wafer was finally removed, leaving a composite wafer in which the GaN epilayers were attached to the diamond substrate. In this case, diamond was bonded to the N-polar HEMT devices. This process maintains the growth direction and the orientation of the built-in polarization fields of the GaN HEMTs for the realization of the 2D electron gas (2DEG). However, due to the bowing of the wafer, the optical lithography was unavailable for the process while an electron beam lithography was utilized for the device fabrication. Silicon nitride (SiN ), with a low thermal conductivity, was the typical adhesive dielectric layer, the heat barrier therefore was mainly located at the interface between GaN and diamond. + +Since the GaN used in this process can be grown on Si substrate, in principle, this process can be scaled to wafers of any size. But the wafer bow and wafer cracks resulted from different stresses during the formation of GaN-on-diamond wafer greatly limit the wafer size. Except for the lattice mismatch between the film and substrate, the thermal stress is also a well-known challenge. The thermal stresses arise due to the thermomechanical properties of the layers in the stack [35], the differences in the thermal expansion coefficients. The stresses bowed and/or warped the wafers as they cooled down from the process temperature to room temperature. By 2007, this process was improved and the 2-inch diameter GaN-on-diamond wafer was produced. In 2009, Francis et al. published the first demonstration of a 4-inch GaN-on-diamond wafer, as shown in Figure 2 [34, 36]. The interface between GaN and diamond through the adhesive layer was characterized by the acoustic properties using picosecond laser-ultrasonic probing. The measurement indicated a good adhesion + +of the interlayer to both GaN and diamond. The GaNon-diamond wafer technology has been demonstrated with 80 to 100 µm thick diamond substrates, which are mechanically stronger and flatter than that of the thin wafers. The developed GaN-on-diamond substrates were then utilized for the fabrication of devices, which demonstrated transition frequencies up to 85 GHz [37]. A comparison of device performances between GaNon-silicon and GaN-on-diamond was reported, which confirmed that the bonding process could avoid the damage to the active region of HEMTs [38]. The fabricated GaN HEMT-on-diamond transistor had a power density of 2.79 W/mm at 10 GHz. The device on SiC had a number of dimensional variables in its favor and demonstrated twice the thermal resistance of that on GaN-on-diamond. By using GaN-on-diamond as opposed to GaN-on-SiC, the operating junction temperature was reduced by 40-45% (Figure 3), and the thermal improvement has tripled the areal RF power density from a GaN transistor. However, the relatively low current densities in the GaN-on-diamond devices limited the output power compared to the devices on + +![](images/3e87cf3ff819153b54dade3fb3732d4bfde9e53faebb0ec92fe7104d3773b55c.jpg) +Figure 2. Photograph of a 4-inch GaN-on-diamond wafer [34]. + +![](images/0f8fc05dbcbf8b3b61d9e168a8849a268af1e2db2bf9203a375b3c2bd1b6141b.jpg) +Figure 3. T he operating temperature of the GaN-on-diamond and GaN-on-SiC HEMT s [38]. + +GaN-on-SiC (Figure 4), which may be attributed to the trapping effects rather than the heating effect. + +To avoid the heat barriers caused by the adhesive dielectric layer, the direct bonding without adhesive agent was developed. J. C. Kim et  al. used a spark plasma sintering process at $1 0 0 0 ^ { \circ } \mathrm { C }$ with a uniaxial pressure to fabricate the GaN-on-diamond wafers [39]. The highresolution transmission electron microscopy (HRTEM) confirmed the formation of GaN on diamond via the chemical bonding. Unfortunately, the localized hot molten zones of interlayer were observed during the spark plasma environment, bringing out the difficulty using this method. + +The high temperature process in the bonding approaches leads to the considerable stress and the wafer bowing due to the large mismatch of TEC. Besides, the high temperature process may also induce chemistry modification and damages of the bonded structure. To overcome these issues, a low temperature bonding technique was developed. The low temperature wafer bonding is successful in various applications such as the fabrication of silicon-on-insulator (SOI), heterogeneous integration, and advanced packaging at temperature below $4 0 0 ^ { \circ } \mathrm { C }$ [40,41]. One approach at room temperature is to employ the surface activated bonding machine, in which the Ar ion beam at high power was used to active both the surfaces and a Si nano-layer was sputtering deposited as the adhesion layer [42]. After the surface preparations, the two samples were bonded at room temperature by contact and press. But due to the highpower Ar ion deposition, an amorphous diamond layer was formed between diamond and the deposited Si as well as between the deposited Si and GaN, as shown in Figure 5 (a) and (b) by HR-TEM [43]. To experimentally obtain the $T B R _ { e f f }$ between GaN and diamond, the time-domain thermoreflectance (TDTR) was utilized, which is a pump-probe technique that can measure thermal properties of the nanostructures and epitaxial films. The $T B R _ { e f f }$ between GaN and diamond using surfaceactivated bonding (SAB) was estimated to be ∼18m2 K/GW, + +corresponding to a thermal boundary conductance (TBC) of 53 MW/m2 K. When the thickness of Si was reduced to 4 nm (Figure 5 (c) and (d)), the $T B R _ { e f f }$ was further reduced to ∼10 m2 K/GW. However, it is noted that, the yield of the direct bonding using the above methods for the formation of GaN-on-diamond is still low for the larger wafer diameters due to the challenge of uniform polishing of the GaN layer and the bonding process. Mu et  al. reported on the 1 cm × 1 cm sized polycrystalline diamond bonded to the GaN/sapphire [42]. Cheng et al. utilized the single-crystalline diamond substrate for the surface activated bonding to GaN and the size was even smaller [43]. + +The success of the low-temperature bonding of GaN to diamond enables the device transfer to the diamond technique. The device transfer has the direct advantage of utilizing the standard high yield GaN HEMT process. Chao et al. reported the GaN-on-SiC HEMTs devices transfer to the polycrystalline chemical vapor deposition (CVD) diamond substrate with a maximum drain current density of 1.2 A/mm and peak transconductance of 390 mS/mm [44]. In this approach, a thin layer of Si-based adhesion was utilized and the bonding process was at the temperature of $1 5 0 ^ { \circ } \mathrm { C }$ It was confirmed that the GaN HEMT-on-diamond maintained lower channel temperatures than the original GaN HEMT-on-SiC while delivering 3.6 times higher RF power within the same active area. In 2017, Liu et al. firstly achieved the 3-inch GaN-on-diamond HEMTs device transferred to the diamond substrate at the bonding temperature of $1 8 0 ^ { \circ } \mathrm { C }$ [45]. The 3-inch device wafer was coated with a thermosetting adhesive layer, then bonded face-down onto a 3-inch SiC temporary carrier wafer. A bonding adhesion layer with a thickness of 15-20 nm was deposited onto the exposed GaN as well as on the 3-inch polycrystalline diamond substrate. For a GaN HEMT at the power dissipation of 10 W/mm, the peak junction temperature of the device was decreased from 241 °C to191 °C after transferring to the diamond substrate. A maximum current density of ∼1 A/mm and a power + +![](images/62cac01ec90d6ca7c3343b1a77d4c52c271e31c9f1668373a8b30533458596ac.jpg) +Figure 4. Electrical characteristics. (a) I-V characteristics for GaN-on-diamond (solid) and GaN-on-SiC HEMT devices (hollow). (b) output power measured at 10 GHz CW and $V _ { \mathrm { D C } } = 2 0$ Vfor for GaN-on-diamond (solid) and GaN-on-SiC HEMT devices (hollow) [38]. + +![](images/c8d6fc9b6e1ef883119274b811b48640d4f7f7d9608e41efa4f5129cc5e2a588.jpg) + +![](images/b1081387a29394148c750a0733c221692cdfa1fdc655a8bb302d5bb0408c6333.jpg) + +![](images/24fb5f5833360360e2bb386285e4acd700b5a5cdbff648575c2d1bcf9b951418.jpg) + +![](images/573b98dde20909bd6c6b152336e810349c2d72f3eee2f542d7394ea0a8c76578.jpg) +Figure 5. (a) Cross-sectional HR-TEM images and (b) high-angle annular dark field scanning TEM images of the GaN/diamond interfaces using SAB bonding. (c) and (d) are for the sample with the improved interfaces [43]. + +density of 5.5 W/mm CW at 10 GHz with the power added efficiency (PAE) of 50.5% were achieved (Figure 6). By using a finite-element analysis, the $T B R _ { e f f } 0 \mathrm { f } 1 9 \mathrm { m } ^ { 2 } \mathrm { K } / \mathrm { G W }$ was obtained in the GaN-on-SiC device, while the $T B R _ { e f f }$ in the GaN-on-diamond was estimated to be 51 m2 K/GW. Further optimization on the thermal conductivity of the adhesion layer, thickness and bonding process is still needed. + +# 3. Gan epitaxially grown on the diamond substrate + +Another method to achieve the GaN-on-diamond is to grow the GaN and HEMT structures on the diamond substrate. The epitaxial growth is challenging due to the large lattice mismatch and thermal mismatch between GaN and diamond, as shown in Table 1. The lattice mismatch between GaN and diamond is 11.8% [46–49]. The thermal mismatch with different TECs also induces a high tensile strain in the epilayer during the cooling down of the sample after growth. There are a lot of efforts in the GaN deposited on different types of diamond substrates, such as single crystal diamond (SCD) with (110), (111) or (100) orientations, nano-crystalline diamond, + +polycrystalline diamond, or the highly misoriented diamond substrate [49–63]. The growth methods include metal organic chemical vapor deposition (MOCVD), molecular beam epitaxial (MBE), HVPE and resonance plasma enhanced MOCVD (ECR-MOCVD) [49–59]. + +The first trial of GaN deposition on the SCD substrate is in 2003. Since the GaN with the device quality can be deposited on sapphire substrates using a buffer layer, in spite of the lattice mismatch of nearly 16.1%, researchers also utilized an AlN nucleation layer to deposit GaN on the type IIa (110) SCD substrate by MOCVD [49]. However, the closely packed GaN grains instead of the smooth surface were observed on the surface. X-ray diffraction (XRD) showed the GaN layer was polycrystalline and hexagonal, with c-plane orientation perpendicular to the substrate. As a result of the carbon incorporation originating from the substrate, a very poor optical quality was determined by photoluminescence. + +In 2003, the epitaxial growth of AlN layer on diamond (100) substrate was improved by plasma-induced MBE [50]. The silicon-doped n-type AlN film on the natural boron-doped p-type diamond substrate formed a hetero-bipolar diode with good rectifying properties and + +![](images/6f519b9b445626ae9ce60fb965f94c944f54fc08c6ab564fd40196a84a1370c2.jpg) +(a) + +![](images/15683dec27e26e5710f92c1ce9b7c57b2524500ed7b38651334d7f7f0feb65af.jpg) +(b) + +![](images/e0073cedf3e9d4121d8fb302a95c4833d00d7e3ca6eef53ae5b41e8ecad3e4e8.jpg) +(c) + +![](images/ea761b7cc02b2421365fd68f3986dd459a356f5ca0738790ebc2b7be983df1c6.jpg) +(d) +Figure 6. (a) A 3-inch GaN-on-diamond wafer by substrate transfer process, (b) SEM image of air bridge of a multi-finger device. Pulsed (c) and DC (d) I-V characteristics of GaN HEMT before (dashed lines) and after (solid lines) substrate transfer. + +surprisingly efficient light emission in the spectral range from 2.7 to 4.8 eV under forward bias. The AlN was wurtzite with the [0001] direction on the diamond. However, the quality of the GaN was still unacceptable on the diamond (100) substrate. The XRD pole figure analysis established the presence of the two domains of epitaxial layer, namely (0001)<10-10 > GaN//(001)[110] diamond and (0001)<10-10 > GaN//(001)[1-10] diamond, which were 90° rotated with respect to each other [51]. The presence of these domains was explained by the occurrence of areas of (2 × 1) and (1 × 2) surface reconstruction of the diamond substrate. When applying highly misoriented diamond substrates toward the [110] diamond direction, one of the growth domains was suppressed and the crystalline quality of GaN was improved. However, the surface of the GaN was still rough and could not be utilized for the HEMT devices. + +The quality of the GaN was improved when the SCD (111) substrate was utilized. In 2009, Dussaigne et  al. reported the GaN grown on (111) SCD substrate by ammonia-source molecular beam epitaxy (NH -MBE) using an AlN buffer layer [52]. The reflectance high + +energy electron diffraction (RHEED) pattern demonstrated the good quality and smooth surface. The wurtzite-structured GaN grown on (111) diamond was with [0001] crystallographic direction. The root mean square roughness from atomic force microscopy (AFM) was 1.3 nm, while the surface was still with the grain boundary morphology. In addition, the cracks were observed for the 1 µm-thick GaN layer. The full-width at half maximum (FWHM) of the XRD rocking curve around (002)-plane was 1500 arcsec. For the GaN grown on the (111) orientated SCD substrate, the films has an in-plane epitaxial relationship [10-10] GaN//[110] diamond. The pre-treatments of the diamond surface were helpful to eliminate the formation of the amorphous layer or the inversion domains [53]. In 2010, Dussaigne et al. used a strain engineered interlayer to improve the surface morphology of the GaN grown on the (111) SCD substrate by MBE using ammonia as nitrogen source [54]. This strain engineered interlayer composed of a sequence of 200-nm-thick AlN layer and 200 nm-thick GaN layer. The rms of the 800 nm-thick GaN epitaxial layer was reduced to 0.6 nm, and no cracks were observed on the + +surface. The mobility of the 2DEG was improved to be 750 cm2 /Vs with the sheet carrier density of 1.4 × $1 0 ^ { 1 3 } \mathrm { c m } ^ { - 2 }$ from the AlGaN/GaN heterojunction, as shown in Figure 7. For the HEMT devices, the 200 nm gate length devices showed 0.73 A/mm maximum drain current + +![](images/93cb0858db06abfe3df4a051f6f4dc8c674b857a0f1a52ab26ef58d63214d643.jpg) + +![](images/706064e09ef7eb27ff2ba1fec5d716da36ef90a804d25ec799e40dd6684e161f.jpg) +Figure 7. T emperature dependent Hall effect measurements: (a) experimental data (opened squares) together with calculated electron mobility considering different limiting factors (blue dashed/dotted lines) and (b) sheet carrier density [54]. + +density and $f _ { T }$ and $f _ { m a x }$ cut-off frequency of 21 and 42 GHz [55]. + +The progress of GaN grown on the SCD (111) substrates was also achieved by the NTT Basic Research Laboratories. In 2011, Hirama et  al. reported the AlGaN/GaN HEMTs with a low thermal resistance grown on SCD (111) substrate by using MOCVD [56]. Benefitting from a high temperature cleaning process at $1 2 0 0 ^ { \circ } \mathrm { C }$ in the hydrogen ambient, the formation of the amorphous interfacial layer was prevented on the diamond substrate, which was a crucial process to obtain the atomically abrupt AlN/diamond heterointerface [57]. After the thermal cleaning, a 180 nm-thick AlN buffer was grown, followed by 20-period AlN/GaN multilayers, then the AlGaN/GaN heterostructure was grown with the GaN thickness of 600 nm. Although the surface morphology did not show atomic steps, the metal-polar was confirmed for the structure using convergent beam electron diffraction in this study. The two-dimensional electron gas (2DEG) was successfully achieved, with the sheet carrier density of $1 . 0 \times 1 0 ^ { 1 3 } \mathrm { c m } ^ { - 2 }$ and mobility of 730 cm2 /Vs. The AlGaN/GaN HEMTs with a 3 µm-gate length showed the maximum drain current of 220 mA/mm, cut-off frequency of 3 GHz and maximum frequency of oscillation of 7 GHz. The thermal resistance between HEMT and diamond is 4.1 K mm/W, as shown in Figure 8. The 0.4 µm-gate-length HEMT showed a dc drain-current density of 770 mA/ mm and breakdown voltage of 165 V [58]. The RF power density of 2.13 W/mm was obtained (Figure 9). This is the first report on the RF power operation of the AlGaN/GaN HEMTs epitaxially grown on the diamond substrate. + +The nano-crystalline and poly-crystalline diamond were also proposed as the template for the epitaxial growth of GaN to achieve the effective thermal + +![](images/99dd3ce9afde8dd08ef3ab7c49df471147d70305525eb3be0357e5f12fa62b8e.jpg) + +![](images/7edb135fe7affbfc54ba865b62d14699bd6fa41dbcf9cc44da8c97e9f7803ffa.jpg) +Figure 8. (a) Setup for measuring the temperature distributions from the side of the AlGaN/GaN HEMT s on the diamond or SiC substrates. Temperature distribution of the AlGaN/GaN HEMT s on (b) diamond and (c) SiC substrates at a dissipated power of 2W (3.2W/mm). (d) Dissipated power dependence of device temperature for AlGaN/GaN HEMT s on the diamond and SiC substrates. Closed and open squares indicate the temperatures for diamond and SiC substrates, respectively [56]. + +![](images/1856916dd3da01ffd0d119199e144c9bfa50532a2cf16a83af5647fdb9b20bc7.jpg) +Figure 9. RF large-signal characteristics of an AlGaN/GaN HEMT at 1 GHz [58]. + +dissipation [60–63]. The diamond was firstly deposited on the silicon substrate, therefore, there is no limitation in the substrate size. A low-temperature buffer layer is typically utilized before the deposition of GaN films. Unfortunately, the GaN layer was not pure wurtzite structure and the polycrystalline form was typically observed [60–62]. Recently, an epitaxial lateral overgrowth (ELO) of single-crystalline thick GaN films were reported on the polycrystalline diamond [63]. The polycrystalline diamond was deposited on the GaN/Si wafers using hot filament CVD. The GaN was grown on the window region between the diamond stripes. A lower pressure, higher V/III ratio, higher temperature, and GaN window mask openings along [11̅00] resulted in enhanced lateral growth of GaN. Complete lateral coverage and coalescence of GaN were achieved over a [11̅00]-oriented 5 μm-wide GaN window between 5 μm diamond stripes. The advantage of the ELO growth is no interlayer formation between diamond and GaN, which will be promising for the effective thermal dissipation. However, there is no device demonstration or the $T B R _ { e f f }$ results from this technique. + +The interface thermal property was analyzed for the HEMT device epitaxially grown on the SCD substrate by N-plasma MBE [64]. A transient interferometric method, in combination with a three-dimensional model, was used to describe a pulsed operation of a transistor-like heater, and a micro-Raman technique was used in a steady state. The thermal conductivity of the diamond was found to be 2200 W/mK, and a relatively lower $T B R _ { e f f } \mathrm { o f } < 1 0 \mathrm { m } ^ { 2 } \mathrm { K } / \mathrm { G W }$ was achieved. The temperature increase in the device was saturated after 1 µs from the start of the heat dissipation and the normalized device thermal resistance of about 3.5 K mm/W was achieved. + +# 4. Diamond epitaxially grown on GaN wafers + +The uniqueness of the direct CVD diamond growth on GaN is that the diamond can be deposited as close as possible to the Joule heating location of the HEMTs. This approach is highly effective for the thermal dissipation of HEMTs. However, there are three limitations that restrict this technique. First, the presence of hydrogen during the growth of CVD diamond requires a dielectric layer such as $\mathrm { S i N _ { x } }$ utilized to protect GaN layer, contributing to the large $T B R _ { e f f }$ for the devices [65]. Second, the diamond nucleating layer begins with many small grains, resulting in a poor thermal conductivity [66]. Third, to avoid the degradation of the GaN, low temperature deposition is needed. However, it has been shown that lowering temperature below $6 0 0 ^ { \circ } \mathrm { C }$ using conventional $\mathrm { H } _ { 2 } / \mathrm { C H } _ { 4 }$ based growth results in the growth of the poor-quality diamond [67]. In addition, the stress is another unavoidable problem. In general, diamond will only nucleate on the carbide forming materials like many refractory metals or Si and will not outgrow in the single crystal phase even on cubic substrates (except for growth on Ir [68]). There is no report on the SCD deposited on GaN, as an alternative, the nano-crystalline or polycrystalline diamond were fabricated [36, 69]. + +The first attempt to deposit diamond films onto hexagonal GaN was by the group of Oba and Sugino [46, 70], who deposited diamond on (0001)-oriented GaN films using microwave plasma CVD. To prevent the etching to the GaN surface, a carburiziation was conducted. The growth of oriented, heteroepitaxial isolated diamond crystals on the GaN surface was achieved, but the crystals did not coalesce into a continuous film due to the low nucleation density. In 2006, May et al. reported the growth of continuous layers of diamond on GaN using a hot filament CVD technique [67]. They found that there was a competition between the rate of diamond deposition and the rate of GaN decomposition, which determined whether net deposition or etching occurred. When the temperature was higher than $6 0 0 ^ { \circ } \mathrm { C } .$ , the GaN decomposed, evolving gaseous $\mathrm { N } _ { 2 }$ which created pinholes in the growing diamond layer or caused it to delaminate. Lowering the substrate temperature below $6 0 0 ^ { \circ } \mathrm { C }$ resulted in a prohibitively low growth rate and poor-quality diamond. + +The earliest pieces of diamond-on-GaN wafer were obtained by growing the poly-crystalline diamond on a dielectric-coated Ga-face GaN-on-Si wafer, then the Si was etched away, leaving behind an N-face GaN-ondiamond wafer [71]. The thickness of the CVD poly-crystalline diamond was 25 µm. By 2006, the Ga-face GaN-on-diamond HEMT was obtained. In this process, firstly the GaN-on-Si epitaxy was bonded onto a temporary Si carrier, then the host Si substrate was etched away, followed by the deposition of a 50 nm-thick + +dielectric (such as SiN ) onto the exposed rear of the GaN. The polycrystalline diamond was finally deposited onto the dielectric for the GaN-on-diamond HEMT devices by using a hot filament CVD process. The wafer size at that time was over 10 × 10 mm2 area, and the epi-surface exhibited countless particulates that obstructed device processing. The electrical current collapse for the first HEMT on diamond was observed. In 2012, the thickness of the polycrystalline diamond was increased to 100 µm [72], and the device performance was further improved [73]. Over 7 W/mm output power density at 10 GHz was reported, along with the peak PAE over 46% and power gain over 11 dB at 40 V. + +In the above process of diamond deposited on GaN, the residual stress was propagated into the GaN layer and caused local defects and uneven bonding of the wafer after the carrier wafer was removed, leading to the wafer bow and warp. The interaction between the intrinsic stress and the built-in stress in the GaN affected the electrical behavior of the device. The effects of the increased stress represent a significant reliability concern for the device, especially when considering the function and life time of the device [43, 74–78]. Jia et al. utilized a double-sided diamond deposition technique to reduce the stress problems [79]. A tensile stress of ∼0.5 GPa was obtained in the GaN layer of the GaN-ondiamond structure, and the crystal quality of the GaN + +was observed to not change significantly after the wafer transfer process. A seamless interface with a ∼10 nm SiC and a thin Al-Si–N intermediate interfacial layer were observed, which facilitated the adhesion between the GaN and the heat dissipation diamond layer. However, the interfacial layer may lead to a high $T B R _ { e f f }$ Zhou et al. investigated the interface thermal property of GaN and polycrystalline diamond with SiN and AlN barrier layers as well as without any barrier layer [80,81] (Figure 10). The $T B R _ { e f f }$ was estimated by TDTR method with a 100 nm-thick Au as the transducer layer, which had a value of ∼ 6.5 m2 K/GW when an ultrathin SiN barrier layers were utilized. The direct growth of diamond onto GaN results in one to two orders of magnitude higher $T B R _ { e f f }$ due to the formation of a rough interface. AlN barrier layers can produce a $T B R _ { e f f }$ as low as that with SiN barrier layers in some cases. However, the $T B R _ { e f f }$ is rather dependent on growth conditions. A decreasing diamond thermal resistance with increasing growth temperature was also observed. To compensate the thermal stress between GaN and diamond during the epitaxial growth, Cuenca and Smith proposed a membrane-based technology [35, 82]. From their analysis based on the analytical models, the bow for a membrane structure with small sizes were underestimated and the bow could be reduced if the membrane was pre-stressed to become flat at CVD temperatures. The + +![](images/4b32118c9f9b167244f6aca063613a6802495879bfc41ae6ccf2d54a18559ccd.jpg) + +![](images/ff85e4914f3de9918115ff7fbacab8070c363316c61848bda6cd578b22924183.jpg) + +![](images/a2675e55e3532c2f872af65c2c54f89dcbc6e813032aa600f6844ea1bd3c61e7.jpg) + +![](images/b5f7c08166913666cf43c532b75c9e43a3f250caf67df275f829dbd41099a2d8.jpg) +Figure 10. (a) sample structure and thermoreflectance measurement scheme. (b) Thermoreflectance signal as as function of time of GaN-on-diamond samples with different barrier layer. Lines represent the experimental value and dots represent an analytical model fitted to the experimental values. (c) Unnormalized and (b) normalized sensitivity curves for the GaN-SiN-diamond samples, with the sensitivity of △R/R corresponding to ±10% change in each input parameter in the model. The laser heating pulse stops at 10 ns [80]. + +sample was from the GaN-on-Si. The Si substrate was selectively removed, and the windows were formed for the polycrystalline diamond deposition using microwave plasma CVD. The diamond was grown on the + +![](images/284d83b9cdaa16546020766e07dc725462a4ed0fd9b05f55e71f85b7ea515c99.jpg) +(a) + +![](images/293d004b175b1dbeb430847f7ac199527a6eae708c040370eb2e36598192f515.jpg) +((b) +Figure 11. (a) The NC D growth was conformal across MESA edges and (b) no cracking or blistering of the growth NC D films even for films as thick as 6.2 µm. + +etched exposed N-polar AlN epitaxial nucleation layers. The high-quality diamond/voi-free AlN interfaces were confirmed from microstructure analysis. This approach is the important demonstration to solve the thermal stress issue during the epitaxial growth of diamond heat spreader on the GaN-based materials and devices. + +The nanocrystalline diamond (NCD) films deposited on the fabricated InAlN/GaN HEMTs devices were reported in 2011 by Alomari et al. [83] (Figure 11). The thermally stable contacts were prepared by depositing a Ta diffusion barrier on the Cu contacts. Then the HEMTs structure were passivated with a thin Si based interlayer (containing the passivated layer and Si-nucleation layer). The NCD nucleation and growth steps were conducted in a hot filament CVD chamber at $7 5 0 \mathrm { - } 7 7 0 ^ { \circ } \mathrm { C } .$ A nucleation density of about $3 \times 1 0 ^ { 1 0 }$ nuclei/cm2 was achieved. The average grain size on the top was around 120 nm in diameter. Raman spectroscopy indicated a dominant diamond peak as the main constituent while the presence of the graphitic phase at the grain boundaries. No degradation or change in the HEMT DC characteristics was observed despite the high temperature of the diamond overgrowth process. Tadjer et  al. reported the NCDcapped HEMTs exhibited approximately 20% lower device temperature from 0.5 to 9 W/mm dc power device operation [84]. NCD-capped HEMTs exhibited the improved carrier density and sheet resistance, but the on-state resistance and the breakdown behaviors were degraded, as shown in Figure 12. The selective growth of the NCD was reported by Ahmed et  al. [65] on the AlGaN/GaN wafer by hot filament CVD. A thin layer of $\mathrm { S i N _ { x } }$ by plasma enhanced CVD, deposited prior to seeding and diamond deposition, was found to be essential to protect the AlGaN/GaN wafer. A methane concentration of 3.0% was used to increase the diamond growth rate and faster surface coverage. Excellent selectively and minimal surface damage to AlGaN were achieved. The + +![](images/089a1bb14ee20bcfedd111e45b3c1d05c991429a1c3fc8bdefed9dcccfdc7455.jpg) + +![](images/b1dd73c8b83eaa3546761d3910ae1f594b319b1d18ed64febb1f287068afdc2b.jpg) +Figure 12. (a) Non-confocal Raman thermography profile of the device channel temperature of AlGaN/GaN HEMT s with and without NC D heat spreading. (b)Temperature-depth confocal Raman profile of the HEMT structures under dc bias. (c) Confocal Raman Si TO peak intensity depth profile showing the GaN/Si interface [84]. + +dual side deposition of the NCD was also reported with a AlN as the dielectric layer [85]. + +The interface thermal properties between GaN and NCD diamond were improved by Smith et al. using a two-step mixed-seeding method [86]. It was found that the mixture of microdiamond and nanodiamond seeding led to a low $T B R _ { e f f }$ While the diamond directly grown onto GaN was proved to be unsuccessful due to the poor adhesion. The two-step mixed-seeding method gave $T B R _ { e f f }$ values lower than 6 m2 K/GW, 30 times smaller than those of films using nanodiamond seeding alone. Such remarkably low thermal barriers obtained with the mixed-seeding process of microdiamond and nanodiamond offer a promising route for the fabrication of high-power GaN HEMTs using diamond as a heat spreader. + +The double-side diamond integration to the HEMTs devices were also developed to improve the heat dissipation. In 2019, Fujitsu Limited and Fujitsu Laboratories Ltd. deposited diamond films on the front surface of the HEMT device. They also bonded the SCD on the back side of HEMTs using SiC interlayer, as shown in Figure 13 [89]. The nanodiamond films were grown at a temperature of $6 5 0 ^ { \circ } \mathrm { C }$ without degrading the transistors’ performance. With the NCD layer on the front side, the amount of heat generated during HEMT operation was reduced by approximately 40% compared to that without the diamond film, and the temperature can be lowered by $1 0 0 ^ { \circ } \mathrm { C }$ or more. By combining the heat dissipation from the back side of the GaN HEMT with SCD substrate, the operating temperature is expected to be reduced by approximately 77%. + +# 5. Interface thermal property between GaN and diamond + +Thermal boundary resistance is an essential concern for GaN-on-diamond transistors especially when they are operating at high-power densities. According to the diffuse mismatch model (DMM), the theoretical limits of $T B R _ { \mathrm { e f f } }$ between GaN and diamond is 3 m2 K/GW [88]. However, the measured values are far from the ideality. The interface transition layer between GaN and + +diamond plays a key role for the contribution of the $T B R _ { e f f }$ no matter if the GaN is bonded to diamond, epitaxially grown on diamond or the diamond grown on GaN devices. The experimental methods that typically used to determine $T B R _ { e f f }$ include the TDTR technique, transient thermoreflectance (TTR) technique, and the three-dimensional (3D) Raman thermography mapping method. The principles of the TDTR and TTR methods are similar. In the TDTR measurement, a pump laser is utilized to periodically heat the sample surface and a probe laser monitor the thermal reflectivity signal of the transducer metal (Al or Au). The reflected intensity tracks the temperature change of the surface, and the normalized reflected intensity is equal to the normalized surface temperature change. Therefore, the change in the thermal reflectivity of the surface is directly proportional to the temperature change. The signal, which is picked up by a photodetector and a lock-in amplifier, is fitted with an analytical heat-transfer solution to infer the unknown parameters. The 3D Raman thermography mapping exploits the temperature induced phonon shift in a material, with respect to a reference phonon frequency measured at ambient temperature [89–92]. The stress can be simultaneously obtained by the simultaneously analyzing the multiple phonon modes [92]. By using a three-dimensional finite element thermal model, the temperature dependent thermal properties can be extracted [93]. The $T B R _ { e f f }$ can also be extracted by analyzing the HEMTs device performance using the steadystate heat conduction model [94]. + +For the GaN or HEMT devices grown on diamond substrate, the buffer layers with the lattice-mismatched induced high-density dislocations, are the main reason for the high $T B R _ { e f f }$ In the bonding process or the diamond deposition on GaN, since the gallium does not readily form a carbide, $\mathrm { S i N _ { x } }$ is typically used to form an interlayer. $\mathrm { S i N _ { x } }$ has a low thermal conductivity of 1-2 W/mK, which introduces an additional thermal resistance. For the polycrystalline diamond deposited on the GaN with a $\mathrm { S i N _ { x } }$ interlayer, the SiN contributes to most of the GaN-on-diamond interface thermal resistance, resulting in a $T B R _ { e f f }$ of more than 30 m2 K/GW, adding >20% to the total device resistance [95]. In + +![](images/8c38a5c316644ea0fae8ffc0cac42e16f94cbb3644ffcc0df419841af4156c08.jpg) +Figure 13. T he heat-spreading method using the double-side diamond and the heat dissipation efficiency [87]. + +Table 2. Summary on the $T B R _ { e f f }$ at the interface between GaN and diamond using different integration methods and different measurement methods. + +
Integration methodInterlayerthickness\(TBR_{eff}\)(m2K/GW)MeasurementRef.
Direct vdW bondingNo0220±70TTR[98]
RT surface-activated bondingSi10 nm~18TDTR[43]
RT surface-activated bondingSi4 nm~11
BondingThermosetting adhesion15-20 nm51Device finite-element analysis[45]
HT bondingAdhesion3-55 nm27-31TDTR[99]
Adhesion3-55 nm25-29DC joule heating
HT bondingSiNx22 nm17TDTR[100]
GaN grown on SCD using MBENot indicatedNot indicated<10TDTR[64]
Hot-filament CVD diamond on GaNdielectric25 nm273D Raman mapping[93]
MPCVD diamond on GaNdielectric50 nm36
CVD polycrystalline diamond grown on GaNSiN100 nm38.5±2.4TDTR[101]
CVD polycrystalline diamond grown on GaNAIN100 nm56.4±5.5
CVD diamond on GaNdielectric50 nm18Raman[102]
CVD diamond on GaNSiNx30 nm29TDTR[103]
CVD diamond on GaNSiNx28 nm12TTR[95]
CVD diamond on GaNSiNx~5 nm<10TDTR[105]
CVD diamond on GaNSiNx~5 nm6.5TTR[81]
CVD diamond on GaNSiC~5 nm30±5.5TTR[104]
CVD Mixed-size diamond seeding on GaNNo0<6TTR[86]
+ +addition, the diamond nucleation layer is usually composed of a few 10 s of nm poor-quality, small-grained diamond [96], which also contributed to the $T B R _ { e f f }$ The direct diamond growth on GaN or bonding to the GaN layer usually shows much weaker interfaces, although there are rare reports on $T B R _ { e f f }$ assessments. It is generally known that weaker interface results in higher $T B R _ { e f f }$ [97,98]. Table 2 summarizes the reported $T B R _ { e f f }$ at the GaN-on-diamond interface using different integration methods and different measurement methods. + +# 6. Summary and outlook + +In this paper, the fabrication of GaN-on-diamond wafers using different methods and the properties of the materials and the fabricated HEMT devices are systematically reviewed. The bonding of GaN or the well-fabricated AlGaN/GaN HEMTs device with diamond can maintain the quality of both the GaN devices and diamond substrate. However, the dielectric interlayers are necessary for either the high-temperature or room-temperature bonding along with an amorphous layer induced by the surface activation. This interlayer impedes the heat flow from the device channel, leading to a large $T B R _ { e f f }$ The growth technique of the GaN or AlGaN/GaN HEMTs structures on the SCD substrates were developed in recent years. A relatively lower $T B R _ { \mathrm { e f f } } ~ { < } 1 0 \mathrm { m } ^ { 2 } \mathrm { K } / \mathrm { G W }$ was reported at the interface between epitaxial GaN and diamond substrate. However, although the HEMTs devices are demonstrated, the crystalline quality of the material and the mobility of the HEMT devices are still far from those grown on the conventional substrates of sapphire, Si, SiC and free-standing GaN substrates due to the large lattice mismatch and thermal mismatch. The CVD + +growth of the polycrystalline or nanocrystalline diamond on the GaN or HEMTs devices has the uniqueness of the direct growth as close as possible to the Joule hot spots, which may be highly effective for the thermal dissipation. However, the damage to the GaN by hydrogen during CVD growth and the less nucleation of the diamond restrict the film quality, leading to a poor thermal conductivity of diamond and large $T B R _ { e f f }$ In addition, the large thermal expansion coefficient between diamond and GaN results in the stress problems in the GaN-on-diamond wafers, leading to the layer cracking and wafer bow and impact the electrical performance of the devices. Up to now, a variety of efforts have been performed to challenge the above problems with different methods. Without doubt, the heat dissipation using diamond as heat spreader for the AlGaN/GaN HEMTs devices is a highly effective way to reduce the operation junction temperature compared to the devices on sapphire, Si and SiC substrates. Although the large $T B R _ { e f f }$ exists between GaN and diamond, Fujitsu laboratory has reduced the device temperature during HEMT operation by more than 40%, and the temperature can be lowered by $1 0 0 ^ { \circ } \mathrm { C }$ or more due to the high thermal conductivity of diamond [87]. Nevertheless, to take the full advantage of the GaN technology for the high-power applications, the reduction of the $T B R _ { e f f }$ between GaN and diamond is still a challenge. Novel strategies and concepts are still required for the effective thermal management of GaN-based power devices using the diamond heat spreader. + +# Disclosure statement + +No potential conflict of interest was reported by the author. + +# Funding + +This work was supported by the JST-PRESTO Grant No. JPMJPR19I7, and World Premier International Research Center (WPI) initiative on Materials Nanoarchitectonics (MANA), Ministry of Education, Culture, Sports, Science & Technology (MEXT) in Japan, and National Key Research and Development Program of China (No.2018YFE0125700). + +# Notes on contributor + +Dr. Liwen Sang is the Independent Scientist at the International Center for Materials Nanoarchitectonics (MANA) in National Institute for Materials Science (NIMS), in Japan. She is also the JST-PRESTO researcher under the support of Precursory Research for Embryonic Science and Technology (PRESTO) program, Japan Science and Technology Agency. 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0000000..eb64523 --- /dev/null +++ b/金刚石散热项目/项目前期工作_MD/01_CVD金刚石薄膜与热扩散片/2_Heat-spreading diamond films for GaN-based high-power transistor devices/2_Heat-spreading diamond films for GaN-based high-power transistor devices.md @@ -0,0 +1,140 @@ +# Heat-spreading diamond films for GaN-based high-power transistor devices + +M. Seelmann-Eggebert a,, P. Meisena , F. Schaudela , P. Koidla , A. Vescanb , H. Leier b + +a Fraunhofer Institut fur Angewandte Festkorperphysik, Tullastr. 72, D-79108 Freiburg, Germany ¨ ¨ b DaimlerChrysler Hochfrequenzelektronik, D-89081 Ulm, Germany + +# Abstract + +We discuss the potential of heat-spreading films with respect to improving the performance of thermally limited high-power high-frequency GaN-FET devices and report on successful diamond deposition on GaN-FETs. Detailed conditions for process compatibility with GaN-FET technology are discussed and shown to be satisfied by the low-temperature deposition process developed.  2001 Elsevier Science B.V. All rights reserved. + +Keywords: Device modeling; Nucleation; Metal semiconductor field-effect transistor MESFET ; Processing Ž . + +# 1. Introduction + +Owing to the combination of high cut-off frequencies, breakdown voltages, saturation velocities and operating temperatures, GaN-based field effect transistors FETs have a great potential for high-frequency Ž . power amplifiers, offering an increase in power output of at least an order of magnitude in comparison with GaAs FETs of comparable device design 1,2 . A limita-  - tion of the performance of GaN is the rise in operating temperature caused by heat dissipation, which leads to large leakage currents 3 and reduced channel mobili-  - ties 2 .  - + +One approach to reduce the operating temperature is to spread the heat from the highly localized source over a larger part of the transistor area. Heat spreading can be achieved by deposition of an electrically insulat- + +ing but highly thermally conducting film. Distinguished by these film properties, diamond is the material of choice. Owing to its excellent heat conductivity -, which exceeds 20 W $\mathrm { c m } ^ { - 1 } ~ \mathrm { K } ^ { - 1 }$ at room temperature, chemical vapor-deposited CVD diamond is used forŽ . heat-spreading coolers. Although the thermal conductivity of CVD diamond reduces with decreasing thickness 4 , we found it to exceed that of gold, even at a  - thickness as low as 2 m. + +However, as for the deposition temperature, plasma exposure and non-invasive seeding of the diamond deposition process need to be fundamentally redesigned to satisfy the compatibility requirements of GaN devices. In this paper, we demonstrate experimentally that CVD diamond can be deposited on GaN-FETs without any degradation of the transistor characteristics by fabrication of the first heat-spread transistor prototypes. + +Before we report on the technical details of diamond deposition on GaN-FETs, we discuss some results of thermal simulations. The thermal efficiency of the dia- + +![](images/d2331ead1b4ea5bb3d8996406426e198fe5bebdb7fda1cd3c089671e044098a0.jpg) +Fig. 1. Sketch of the device geometry: a model used for the thermal simulations; and b enlarged cross-section of the transistor finger. Ž . Ž . + +mond spreading-layer is analyzed and compared with other approaches to the thermal management of power FETs. + +# 2. Thermal modeling of GaN-FETs + +Obtained by dicing a chip from the wafer, a GaN-FET is located on the surface of a cube-like slab. As sketched in Fig. 1, the FET dimensions are much smaller than those of the chip, which therefore have little influence on the thermal device performance. In the FET structure, the heat source is filament-like. Although the filament length is much greater than the thickness of the epilayer, the filament diameter is much smaller, and the heat is practically spread in the plane of the GaN layer. Also, there is already a significant temperature drop in the GaN layer and the device temperature can be less influenced by the choice of the substrate material than expected beforehand. For cost reasons, FET power devices would be preferably fabricated from GaN-structures grown on sapphire rather than SiC substrates, although the latter have a 10-fold higher heat conductivity. + +For our model simulations, a chip $5 0 0 \times 5 0 0 ~ \mu \mathrm { m } ^ { 2 }$ in area was chosen. The GaN layer $( \kappa = 1 . 3 \mathrm { W } \mathrm { c m } ^ { - 1 } \mathrm { K } ^ { - 1 } )$ is assumed to have a thickness of 1 m and to be grown on a sapphire substrate of 300 m thickness $\overline { { ( } } \ l _ { \mathsf { K } } = 0 . 4 6 \ \mathrm { ~ W ~ } \mathrm { c m } ^ { - 1 } \ \mathrm { ~ K } ^ { - 1 } )$ . In the center of the chip surface was a mesa $( 1 0 0 \times 1 2 5 \times 0 . 5 ~ { \mu \mathrm { m } } ^ { 3 } )$ with a single-finger FET of gate width 100 m and gate length 0.25 m. The gatesource and gatedrain distances were 1 m each. All contact pads were $9 0 \times 1 2 0 ~ \mu \mathrm { m } ^ { 2 }$ . Extending from the edge of the gate contact over a distance of 0.15 m towards the drain, the heat source was assumed to be 25 nm below the surface and to have a height of 15 nm. Heat generation was assumed to be uniform over the heat-source volume 5 . - + +The modeling results are discussed in terms of thermal impedance 5,6 rather then temperature. The local  - thermal impedance is the temperature rise as mea- Ž + +sured with reference to the heat sink normalized to . the heat power dissipated per gate-width unit as an Ž arbitrarily chosen characteristic device length . To esti- . mate the interdigital crosstalk of multifinger devices, it is convenient to analyze the thermal impedance profile Ž . TIP of a single FET-finger as a plot along the distance from the center of the finger in the plane of maximum temperature 6 i.e. a line just below the  - Ž GaN surface in the cross-section of Fig. 1b . Since the . TIPs of different fingers are very similar, the temperature profiles over a comb of fingers can be obtained from a single TIP by lateral translation and subsequent superposition. The introduction of the TIP is useful owing to two simple scaling rules for stationary heat flow: for a considered thermal-impedance profile the temperature is proportional a to the heat load and b Ž . Ž . to the reciprocal gate width, if the whole geometry is rescaled with no change in shape. The peak temperature in the device is obtained by multiplying the peak impedance with the heat dissipated per gate-width unit. + +The curves shown in Fig. 2 are thermal simulation + +![](images/80f1528fed435d61216e4ac7257f20d84aa7bfebc9bccaf16d90c82eed3366f1.jpg) +Fig. 2. Thermal impedance profiles of GaN-FETs for different thermal setups: a bare finger; b 3- Ž . Ž . Ž . m gold bridge as thermal shunt; c 3-m diamond spreading layer; d gold bridge in flip-chip geometry; Ž . and e spreading layer in flip-chip geometry. InŽ . $\mathrm { ( a ) \mathrm { - } ( c ) , }$ the heat sink is at the substrate bottom, in d and e at the top. Ž . Ž . + +results for the single-finger device sketched in Fig. 1a operating under different thermal conditions. Despite the asymmetric position of the heat source with respect to the gate contact, differences in the left and right branch of the curves in Fig. 2 are negligibly small. + +The uppermost curve represents the situation of the GaN-FET finger without any thermal precaution. The peak thermal impedance, $\dot { R _ { \mathrm { o } } } = 2 9 . 5 \ \dot { \mathrm { K } } \ \mathrm { m m } \ \mathrm { W } ^ { - 1 }$ , obtained by the simulation is only approximately half the value predicted by the Smith formula for the considered heat source geometry embedded in a sapphire matrix 5 . Hence, the GaN epilayer, with a compara-  - tively high thermal conductivity itself, results in a significant reduction in the operating temperature. In Fig. 2a, the heat-spreading effect of this layer is seen as a broad appearance of the TIP and a slow decrease at large distances x from the gate $( \Delta T \sim 1 / \sqrt { x } )$ . This behavior implies a comparatively strong thermal coupling between multiple FET fingers if placed at a typical distance of 30 m: the respective coupling impedance $R _ { \mathrm { c } }$ is approximately one third of $R _ { \mathrm { o } } .$ . + +Conventional thermal shunts are formed by thick goldair bridges $( \kappa = 3 \mathrm { ~ W ~ c m } ^ { - 1 } \mathrm { ~ K } ^ { - 1 } )$ attached to the gate contact and a large fraction of the chip area. In our model calculation, the gold bridge was 3 m thick and as sketched in Fig. 1b connected to the large Ž . drain and source contact pads via a 0.2-m thick insulating SiN layer $( \ l _ { \mathsf { K } } = 0 . 0 1 \hat { 6 } \mathrm { ~ W ~ c m } ^ { - 1 } \mathrm { ~ K } ^ { - 1 } )$ . The thermal design reduces the thermal peak impedance to 70% of its initial value, however, the coupling impedance is lowered only by 10% Fig. 2b . Hence, the situation Ž . regarding thermal crosstalk is not improved. + +A drastic temperature reduction is found if a closed spreading layer of diamond is deposited by use of Ž $\overset { \cdot \mathrm { ~ ~ \hat { ~ } { ~ \theta ~ } ~ } } { \left. \mathrm { ~ \bf { { K } } ~ } \right. } = 3 \mathrm { ~ \bf { W } ~ } \mathrm { { { c m } ^ { - 1 } ~ } } \mathrm { { \bf { K } } ^ { - 1 } }$ , fine-grained CVD diamond of poor thermal quality is assumed . Protecting the exposed . surface areas of the transistor, a 20-nm thick SiN layer was included in the calculation. The thermal impedance peak is reduced to 40% of the original value, but the $R _ { \mathrm { c } } / R _ { \mathrm { o } }$ ratio increases further. The large improvement in $R _ { \mathrm { o } }$ with respect to the thermal-shunt approach is a consequence of the lateral distance of the heat source and gate contact. Covering the heat source completely, the film bridges this gap. + +Appreciable thermal crosstalk is common to the three considered situations Fig. 2a,b,c and results from the Ž . position of the heat sink at the bottom of the substrate, i.e. in the thermal far field. Flip-chip designs bear an appreciable advantage over backside-cooled spreading or shunt designs, if the sink can be placed in the near field and, consequently, crosstalk is substantially reduced. Although technically impossible to realize, it is instructive to place the heat sink directly on top of the bridge or spreading layer. The gold bridge design with topside cooling brings the $R _ { \mathrm { c } } / R _ { \mathrm { o } }$ ratio down to 0.02, although the peak temperature reduction is approxi- + +mately the same as for the diamond spreading-layer with backside cooling. If a respective flip-chip set-up could be realized with the use of a closed spreadinglayer, the temperature rise during operation could be reduced by a factor of 5.2 with a crosstalk ratio $R _ { \mathrm { c } } / R _ { \mathrm { o } }$ of 0.01. + +In practice, the apparent advantages of the flip-chip approach are less pronounced. Since relatively thick thermal bumps are required to mediate the bonding of the heat sink to the chip surface, the device temperature is lowered by heat spreading in the thermal bumps rather than by near-field heat flow to the sink. Nevertheless, to achieve high cooling efficiencies, both thermal management approaches require coupling of the film to a large area over the heat source  a condition much easier to satisfy with a dielectric than with a metal. + +# 3. Process requirements + +In view of its superior thermal and dielectric properties, CVD diamond appears to be the material of choice for heat-spreading applications. However, hostile conditions for deposition impede this application of CVD diamond on delicate semiconductor devices. High quality diamond is obtained in a methanehydrogen plasma at elevated deposition temperatures $( 8 0 0 ^ { \circ } \mathrm { C } ) .$ . Deposition at lower temperatures is possible, but brings about losses of quality and low growth rates 7 . For - GaN devices, which are themselves grown at temperatures near $1 0 0 0 ^ { \circ } \mathrm { C } ,$ a temperature of $\mathrm { \bar { 5 } 0 0 ^ { \circ } C }$ appeared to be an appropriate upper limit for diamond deposition, mainly implied by stability considerations of the FET contacts. On silicon substrates, we realized and routinely grow diamond films at this temperature, with thermal conductivities exceeding that of gold 3 W Ž $\mathsf { c m } ^ { - 1 } \ \mathsf { K } ^ { - 1 } )$ , a thickness $> 3$ m and growth rates of $0 . 2 ~ \mu \mathrm { m } ~ \mathrm { h } ^ { - 1 }$ . However, on epitaxial GaN films, diamond deposition is impossible, due to rapid GaN etching $( > 1 \mu \mathrm { m } \mathrm { h } ^ { - 1 } )$ in the plasma, with the eventual formation of gallium droplets. A further problem may occur for p-doped material, by rapid diffusion of atomic hydrogen from the plasma into the GaN matrix. + +# 4. The protective layer + +In view of the obvious instability of GaN in the CVD plasma, the question arises as to how diamond films can be deposited on GaN-based FET devices. However, this apparent obstacle can be overcome by the use of a suitable protective layer satisfying the following requirements. The protective layer should be plasmaresistant and impervious to hydrogen. With regard to the high-temperature cycle employed, it should show + +![](images/c7fcaab5445039e3bab69353246f28b317063be499f99ce7487d6ce6c953def5.jpg) +Fig. 3. IR transmission spectra of: a plane 300- Ž . m sapphire substrate; b 500 nm SiN on sapphire after 30-min CVD plasma expo- Ž . sure; and c same layer after 2-h vacuum annealing at 550 Ž . C. + +good adhesion properties to GaN and diamond. Seeding properties of the protective layer are important for stability and seed formation. High heat conductivity is desirable, although not necessary, if plasma protection is warranted, even at a very small thickness of the interlayer. + +AlN, $\mathrm { S i O } _ { 2 }$ and SiN layers were investigated and found to satisfy these requirements, but SiN was finally chosen, mainly in view of process compatibility for the entire device fabrication. Crystalline $\mathrm { S i } _ { 3 } \mathrm { N } _ { 4 }$ has a heat conductivity of 0.3 W $\mathrm { c m } ^ { - 1 } \ \mathrm { K } ^ { - 1 }$ ; however, for thin microcrystalline films, much lower values are reported  - 8 . Measurements on CVD-SiN layers routinely used in the standard fabrication process yielded a conductivity of 0.016 W $\mathrm { c m } ^ { - 1 }$ $\mathbf { K } ^ { - 1 }$ 9 . Even upon long-term  - exposure 10 h , films of this CVD-SiN material were Ž . not etched by the deposition plasma. + +However, SiN films of 0.10.2 m thickness showed a trend of wrinkle formation after prolonged exposure, indicating local delamination 10 . IR transmission  - studies implied some migration of atomic hydrogen through the SiN films. After plasma exposure, absorption features corresponding to SiH and NH stretching vibrations are observed, indicating that hydrogen is trapped in the SiN layer Fig. 3 . The hydrogen features Ž . are appreciably diminished upon annealing in vacuum. Hydrogen diffusion is likely to be much less efficient in the active parts of the device, since these become quickly passivated by the growing diamond overlayer. + +# 5. Seeding + +Appropriate seeding is required to achieve a high nucleation density and to rapidly form a closed diamond film. Common seeding procedures, such as mechanically polishing with diamond powder, are not + +suited for structured micro-devices. We therefore attempted to seed by means of an ultrasonic treatment in a dispersion of diamond powder in water or methanol. This approach yielded generally high nucleation densi-Ž ties in the order of $1 0 ^ { \overline { { 1 0 } } }$ 2 cm with good uniformity, . even over edges of the device structure, and facilitated quick overgrowth of the GaN device. However, the protective SiN coating was found to be damaged by this treatment, even when the treatment was short and diamond powder with a sub-m particle size was used. After a few minutes of plasma exposure, the surface appearance of ultrasonically seeded samples became spotty and the SiN film was observed to peel off locally at these spots. + +Ultrasonic seeding was also found to deteriorate the transistor performance. At the typical pinch-off voltage of 6 V, large leakage currents of approximately 10% of the saturation current were found after the ultrasonic treatment. The observed change in the electrical characteristics indicated damage of the gate electrode caused by the impact of fast diamond particles. + +Owing to the failure of the conventional approaches, an alternative non-invasive seeding method was developed, based on the sedimentation of fine diamond particles from an agitated emulsion. An advantage of this technique is its independence from the surface properties of the substrate. With this seeding technique, very homogeneous and uniform diamond films have been grown on 2-inch wafers of silicon and glass. + +To achieve a selective area growth for diamond films, the sample surface was structured by selective seeding. To avoid diamond growth on specific areas, such as the contact pads, these were covered by photoresist. After the seeding procedure, the photoresist was dissolved in hot acetone or dimethylformamide. The seed layer was completely removed with the photoresist. Hence, diamond grew only on the exposed areas of the samples, whereas unseeded regions remained practically free of deposits. + +# 6. The deposition process + +Diamond deposition was performed in a plasma-CVD ellipsoid reactor 11 operating at 2.45 GHz with a  - microwave power up to 6 kW and a power density of 100 W cm3 . The sample was located at the bottom of the reactor on a water-cooled stage. The deposition temperature was adjusted by variation of the process pressure or heat-resistance of the stage. For low-temperature deposition, a pressure in the range 70110 mbar was chosen. At higher pressures, the heat flux density exceeded a critical value, causing thermal cracking 12 of the thin 2-inch sapphire substrates due  - to temperature inhomogeneities. + +Low-Temperature nucleation of the film proved to be a critical step. To form a closed film within a period of no longer than 10 min, a temperature of at least 500C as measured by a pyrometer operating at a Ž wavelength of 1000 nm and a largely increased con- . centration of methane 4% was required. After suc- Ž . cessful nucleation as monitored by interferometry , Ž . the methane content of the process gas was set to 1% and, optionally, the temperature was reduced. The following parameters are a typical example of the process conditions: at a pressure of 80 mbar, the heat flux density at the sample surface was 30 W cm2 . If the deposition temperature of 500C is adjusted by a suitable choice of the thermal resistance of the stage, a growth rate of $0 . 2 5 ~ \mu \mathrm { m } ~ \mathrm { h } ^ { - 1 }$ is obtained. + +Adhesion of the diamond films was problematic if a large-area deposition was attempted. If the thickness exceeded approximately 1 m, spontaneous flaw formation was observed and the diamond film delaminated as a result of compressive thermal stress built up during cooldown. Somewhat surprisingly, the diamond films turned out to adhere very well on selectively seeded FET samples. On the structured chips, adhesion was aided by separation grooves which were left unseeded and defined individual spreading layers for each FET. + +With the process developed, diamond spreading layers were routinely deposited on GaN-FETs. As demonstrated by the REM image of Fig. 4, the diamond layer is distinguished by very good selectivity and exhibits remarkably abrupt edges at the contacts. + +The devices investigated in this study were fabricated using a standard GaN-transistor process with gates defined by optical lithography. More details on the fabrication process are found elsewhere 13 .  - + +For checks on the compatibility of the CVD-diamond deposition with the overall FET fabrication process, a piece of a wafer was electrically DC characterized before diamond deposition and the protective layer was then put down. Next, areas to be excluded + +![](images/f0afa72faec37c6c1c8a978d85de32fee1bc1ae682f5f83add505ffe8c3e86a3.jpg) +Fig. 4. REM image of a two-finger GaN-FET with a 0.7-m diamond film deposition temperature 440 Ž . C . + +![](images/f80c5ad4e2c00ec3811e9c2b7bb89eb46bd94d9be67cd0650ace4c6ef5ea4cd5.jpg) + +![](images/9e33edcd0f5fa244166f8fbead9ad2a677ef299bf7bfef81c306f0a573847c4d.jpg) +Fig. 5. a Output; and b transfer characteristics of a two-finger Ž . Ž . GaN-FET before and after deposition of a 0.7-m thick diamond film. + +from deposition such as the contact pads were cov- Ž . ered with photoresist. After selective seeding, the diamond film was grown. The protective layer was then removed from the contacts by dry etching in a $\mathrm { C F } _ { 4 } / \mathrm { O } _ { 2 }$ plasma and the FET characteristics were recorded for a second time. Fig. 5 shows a comparison of the output and transfer characteristics before and after diamond deposition at a deposition temperature close to $4 0 0 ^ { \circ } \mathrm { C } .$ Fig. 5 demonstrates a well-working transistor, which was subject to only minor alterations in the process of diamond deposition. However, above $4 4 0 ^ { \circ } \bar { \mathrm { C } } ,$ serious degradation of the NiAu gate contact occurred, leading to failure of the transistors. EDX measurements revealed that this deterioration had to be attributed to alloying, which caused the gate contact to lose its Schottky character. With an improved GaN-FET processing scheme, stable device operation was accomplished, even at temperatures above $5 1 0 ^ { \circ } \mathrm { C } .$ In the absence of gate-metal alloying phenomena, no signs of transistor degradation after diamond deposition were found. However, for stable gate contacts, an increase in Schottky barrier height was observed after the plasma treatment, an effect typical for heat-treated GaN transistors, caused by an interfacial reaction between the gate metal and the semiconductor. + +In a recent run, transistors with the new stable Schottky contacts were investigated at a deposition + +temperature of $4 8 0 \mathrm { { } ^ { \circ } C }$ . Even after deposition of $2 { \cdot } { \mu } \mathrm { m }$ diamond, the FETs remained fully operational. + +# 7. Summary and outlook + +Heat-spreading diamond films are shown to be an useful approach to reducing the operating temperature of GaN-FETs and, hence, to have the potential for boosting power output of high-frequency GaN devices. Direct CVD deposition of diamond on GaN-FETs is demonstrated to be feasible and completely compatible with state-of-the-art GaN device-processing technology. To achieve this compatibility, a novel gentle seeding process and a low-temperature $( < 5 0 0 ^ { \circ } \mathrm { C } )$ deposition process was developed, involving a protective layer. Temperature-resistant gate contacts are required to allow the deposition temperature to exceed $5 0 0 ^ { \circ } \mathrm { C } .$ . To our knowledge, this is the first demonstration of successful CVD diamond deposition directly on an operational IIIV semiconductor circuit. Detailed experiments proving the cooling efficiency of the spreading diamond still remain to be performed. + +# Acknowledgements + +H. Guttler and M. Hirsch of DaimlerChrysler and E. ¨ + +Worner, C. Wild and R. Sah of Fraunhofer IAF are ¨ gratefully acknowledged for helpful discussions. This work was supported by the BMBF. + +# References + + - 1 Y.-F. Wu, B.P. Keller, S. Keller et al., IEICE Trans. Electron. E82-C 11 1999 1895.Ž . Ž . + - 2 C.E. Weitzel, Institue of Physics Conference Series, 142, 1996 Ž . 765 chapter 4 . Ž . + - 3 E. Kohn, W. Ebert, A. Vescan, Isr. J. Chem. 38 1998 105. Ž . + - 4 E. Worner, in: B. Dischler, C. Wild Eds. , Low-Pressure Syn- ¨ Ž . thetic Diamond, Springer-Verlag, Berlin, 1998, p. 137 chapter Ž 9 .. + - 5 R. Anholt, Electrical and Thermal Characterization of MES-FETs, HEMTs and HBTs, Artech, Norwood, 1994, pp. 5863. + - 6 R. Anholt, Solid State Electron. 42 1998 849.Ž . + - 7 Y. Muranaka, H. Yamashita, H. Miyadera, Diamond Films Technol. 5 1995 1. Ž . + - 8 S.R. Mirmira, L.S. Fletcher, J. Thermophys. Heat Transfer 12 Ž . Ž . 2 1998 121. + - 9 H. Guttler, personal communication. ¨ + - 10 G. Gille, B. Rau, Thin Solid Films 120 1984 109. Ž . + - 11 M. Funer, C. Wild, P. Koidl, Surf. Coat. Technol. 116 ¨ 119 Ž . 1999 853. + - 12 K.J. Gray, H. Windischmann, Diamond Relat. Mater. 8 1999 Ž . 903. + - 13 A. Vescan, R. Dietrich, A. Wieszt et al., Institute of Physics Conference Series, 166, 1999 503 chapter 7 . Ž . 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differ diff --git a/金刚石散热项目/项目前期工作_MD/01_CVD金刚石薄膜与热扩散片/3_State-of-the-art synthesis and post-deposition processing of large area CVD diamond substrates for thermal management/3_State-of-the-art synthesis and post-deposition processing of large area CVD diamond substrates for thermal management.md b/金刚石散热项目/项目前期工作_MD/01_CVD金刚石薄膜与热扩散片/3_State-of-the-art synthesis and post-deposition processing of large area CVD diamond substrates for thermal management/3_State-of-the-art synthesis and post-deposition processing of large area CVD diamond substrates for thermal management.md new file mode 100644 index 0000000..b1aed36 --- /dev/null +++ b/金刚石散热项目/项目前期工作_MD/01_CVD金刚石薄膜与热扩散片/3_State-of-the-art synthesis and post-deposition processing of large area CVD diamond substrates for thermal management/3_State-of-the-art synthesis and post-deposition processing of large area CVD diamond substrates for thermal management.md @@ -0,0 +1,217 @@ +# State-of-the-art synthesis and post-deposition processing of large area CVD diamond substrates for thermal management + +W.D. Brown a,b,\*, R.A. Beera a,b, H.A. Naseem a,b, A.P. Malshe a,b,c + +$^{a}$ High Density Electronics Center (HiDEC), University of Arkansas, Fayetteville, AR 72701, USA $^{b}$ Department of Electrical Engineering, University of Arkansas, Fayetteville, AR 72701, USA $^{c}$ Materials and Manufacturing Research Laboratory (MRL), Department of Mechanical Engineering, University of Arkansas, Fayetteville, AR 72701, USA + +# Abstract + +Primarily due to its outstanding thermal conductivity and high electrical resistivity, diamond is an ideal heat spreader for a variety of applications such as multichip modules (MCMs), laser diode arrays, power modules, etc. For these and similar applications, there is a requirement for the synthesis and post-deposition processing of high quality, stress-free, large area $(1 - 100\mathrm{cm}^2)$ , CVD diamond substrates. Although the thermal characteristics of CVD diamond are primarily determined by proper nucleation and growth conditions, post-deposition processing such as polishing, planarization, drilling/cutting, and metallization are required for the practical application of diamond as a heat spreader. Such processes should be fast, compatible with conventional electronic packaging, adaptable to large area fabrication, environmentally-safe, and economically-viable. In the past few years, significant technological advances have been made in all these areas. Furthermore, the cost of these processes has been decreasing steadily so that fabrication of diamond-based electronic products is becoming affordable. This paper reviews the present technological and economic status of CVD diamond for thermal management applications. + +Keywords: Synthesis; Post-deposition processing; Cvd diamond substrates; Thermal management + +# 1. Introduction + +Rapid changes in the electronics industry are driven by the attractiveness of smaller, faster, and lighter weight electronic systems. Consequently, multichip packaging technology is receiving widespread attention. The use of multichip modules (MCMs) promises to increase packaging density and system performance beyond what is otherwise possible with advances in VLSI and surface mount technology. As is the case for most advanced technologies, the drive towards high performance multichip modules has given rise to a number of engineering challenges. In advanced electronic packages, where high speed, high power chips are mounted shoulder-to-shoulder on large circuit boards, heat spreading and removal become major issues. Typically, thermal management substrate materials for electronic packages have been limited to silicon, alumina, and to a much lesser extent, AlN, SiC, and glass/ceramics. However, as the power dissipation of ICs continues to increase, + +heat dissipation and removal will be major issues in future MCMs. More than likely, diamond will play a significant role in the solution to these performance challenges. + +# 1.1. Why diamond for heat spreaders? + +For the single chip or multichip module, or any high density electronics packaging to be successful and reliable, efficient thermal management (TM) is essential. Diamond has an excellent thermal conductivity ( $k = 2000\mathrm{W / m}\cdot \mathrm{K}$ ) for natural diamond and about 700-1600 W/m·K for CVD diamond, a low thermal expansion coefficient ( $\approx 10^{-6}$ ), high electrical resistivity ( $>10^{12}\Omega \cdot \mathrm{cm}$ ), and high mechanical strength (Young's modulus $\approx 10^{11}\mathrm{Nm}^{-2}$ ). These properties of diamond surpass those of all the established substrate materials such as copper, silicon, alumina, aluminum nitride, etc., and make it an ideal heat spreader. Recent advances in CVD diamond deposition technology make it possible to fabricate large area ( $10\times 10\mathrm{cm}$ ) diamond substrates in the laboratory at a reasonable cost. The availability + +of such freestanding diamond makes possible the practical implementation of such a novel heat spreader. + +# 1.2. Thermal considerations + +In the investigation described here, a popular software package, ANSYS, (version 5.0SASI, Houston, PA), which employs the finite element method for thermal simulation, was used to observe the effects of substrate thickness parameters on TM [1]. It was assumed that the contact resistance between materials (for example, an IC and the substrate) is zero. Also, heat loss due to radiation is neglected. For all simulations, the substrate was assumed to be a square with each side being $10.16\mathrm{cm}$ in length. The power dissipated on the top surface of the substrate was $10\mathrm{Wcm}^{-2}$ for a total power of $1032.256\mathrm{W}$ dissipated uniformly on an area of $103.2256\mathrm{cm}^2$ . This is an extreme situation for most electronic packaging applications. The two opposite edges were kept at a constant temperature of $300\mathrm{K}$ with adiabatic boundary conditions applied to the other two edges and the bottom surface of the substrate. Thus, the maximum temperature rise is expected to be in the center with isotherms parallel to the two opposite edges held at $300\mathrm{K}$ . For calculation purposes, the thermal conductivity $(k)$ of the diamond substrate was assumed to be isotropic thus avoiding any thermal resistance caused by grain boundaries in the diamond. + +Fig. 1 is a plot of the maximum temperature as a function of diamond thickness for the simulation conditions noted above. The thermal conductivity was assumed to be $1500\mathrm{W / m}\cdot \mathrm{K}$ . As the thickness of the diamond increases, its contribution to the reduction of temperature becomes smaller. This trend has an important implication in determining an optimum value for the thickness of the substrate. Currently, a substrate thickness considered sufficient for MCM application is typically $1000~{\mu\mathrm{m}}$ . According to the simulation results, diamond substrates $20 - 30\%$ thinner can be used with + +similar TM performance. For example, the rise in temperature for a diamond substrate of $800\mu \mathrm{m}$ thickness under the conditions noted above is $107.54\mathrm{K}$ . This result suggests that the cost of growing CVD diamond can be reduced by limiting the thickness of a substrate to the maximum value required for a specific application. + +# 1.3. CVD diamond substrate fabrication-related problems + +# 1.3.1. Diamond synthesis + +In state-of-the-art synthesis, CVD diamond films are deposited at high temperatures $(>750^{\circ}\mathrm{C})$ either by localized or bulk heating in a mixture of a carbon precursor gas (methane, acetylene, etc.) and a graphite etchant gas (hydrogen) in the presence of an activation agent(s) such as IR radiation, microwave radiation, a d.c. arc, r.f. radiation, or a laser(s) [2]. Presently available CVD diamond equipment, such as d.c. arc jet, microwave plasma, and r.f. plasma systems, are capable of growing diamond films as large as $10\times 10\mathrm{cm}$ . However, films synthesized using various types of activation agents typically have chemical and morphological non-uniformities over the surface and throughout the volume of the material [3]. Also, large diamond substrates can exhibit high intrinsic stresses that result in warping unless deposition conditions are tightly controlled. In general, the higher the thermal conductivity, the higher the stress. This limits the range of thermal conductivities available for large area, thermal management substrates [3]. + +These techniques are used to deposit diamond over larger and larger areas in order to reduce the cost (per square) of the material. This approach does decrease the cost to some degree, but it also introduces the need for a laser cutting capability in a manufacturing process which presents an added expense. Another factor that must be addressed is the fact that electronic packages are available in various shapes, sizes, and materials + +![](images/50de69532d92b83fb069c14a2b6c503a398ff30cbcf1ae224fc432d29afae47a.jpg) +Fig. 1. Plot of maximum temperature vs. diamond substrate thickness for $k = 1500 \, \mathrm{W/m} \cdot \mathrm{K}$ . + +which requires a capability for depositing diamond onto different materials of both flat and complex geometries. + +# 1.3.2. Diamond post-synthesis processing + +Diamond growth in the initial stages of deposition progresses by nucleation at mechanically-induced and/or randomly-seeded and/or thermally-favored sites because of statistical thermal fluctuations at the substrate surface. Depending on the growth surface temperature, pressure, and reactive environment conditions, favored crystal orientations dominate the competitive growth process. Consequently, the resulting films have a random polycrystalline structure and a rough surface morphology. Large surface roughnesses $(R_{\mathrm{a}} = 2 - 15\mu \mathrm{m})$ often limit the TM efficiency of diamond substrates because of large and numerous surface pits which result from the presence of microcavities in the as-grown bulk material [4]. As a result, post-deposition processing is required if CVD diamond is to be used in many electronic packaging applications. Several novel approaches to post-deposition polishing of CVD diamond films have been investigated recently, including laser/ion beam/plasma treatment, hot metal treatment, and chemical-assisted mechanical polishing and planarization (CAMPP) [5]. + +Unfortunately, these polishing techniques do not eliminate "microcavities" in diamond substrates because diamond crystals grow faster in a direction normal to the substrate surface than parallel to it resulting in columnar growth. During growth, these crystals grow into each other creating microcavities throughout the material (Fig. 2). The presence of microcavities prevents any polishing process from yielding a pit-free surface. This, in turn, presents serious problems for photolithographic processes, such as the formation of metal traces, and die bonding in applications requiring uniform thermal contact between the die and substrate. These surface + +![](images/90109b2bbaff5f5252be072468e1291cf9fc334b39daddedfff6a445c8359ac6.jpg) +Fig. 2. SEM micrographs of microcavities in CVD diamond. + +pits also present a problem for attachment of large die by introducing stresses. Furthermore, the sharp edges of the microcavities lead to an increase in surface nonuniformities after electroplating of a metal layer [6]. The size and density of these microcavities can possibly be reduced, but not eliminated, by increasing the density of nucleation sites on the substrate and depositing at a lower rate. At the present time, the only practical way of eliminating these surface pits is by "filling-for-planarization" [7]. + +Because of the attractiveness of CVD-diamond for passive electronics applications, reliable metallization systems must be available. Furthermore, in order to produce thick interconnects, plating is the preferred method of applying a metal, such as copper or gold, to the diamond. Unfortunately, owing to its chemical inertness, diamond is expected to have poor adhesion to metals, especially to gold (Au) - a noble metal best suited for conductor traces because of its excellent electrical conductivity. The use of gold, then, requires a metal adhesion/seed layer. Fortunately, diamond is known to form adhering carbide layers with refractory metals such as Ti, W, Mo, etc., and transition metals such as Cr, Ni, Fe, etc. Cr has been used extensively in MCM technologies as an adhesion promoter and diffusion barrier metal and, since it lies close to the refractory metals in the periodic table, it is expected to have excellent carbide forming properties with diamond. Plated gold or copper can then be used to form the interconnections. + +# 2. Experimental + +Some of the major steps in the fabrication of diamond heat spreaders are: (1) seeding of substrates for the nucleation of diamond growth, (2) surface finishing of diamond substrates, (3) cutting and drilling of diamond substrates, and (4) metallization of diamond substrates. Subsequent sections address these subjects in some detail. + +Raman spectroscopy (RS) was used for the chemical analysis of diamond, scanning electron microscopy (SEM) and/or atomic force microscopy (AFM) were used to study the surface morphology of diamond substrates, and contact surface profilometry and AFM were used for surface roughness analysis of processed diamond heat spreaders. X-ray diffraction was used to determine residual stress in diamond films. + +The thermal conductivity of diamond is an extremely important parameter for thermal management substrates. Various techniques, such as the heated bar technique, the Sinku-Riko modified angstrom technique, the Omega technique, the laser flash technique, the transient grating technique, and the photothermal technique, have been and continue to be investigated for + +thermal conductivity analysis of heat spreaders. A detailed discussion of these techniques is beyond the scope of this paper, but can be found in a recent publication by Grabner et al. [8]. + +# 3. Results and discussion + +# 3.1. Seeding for large area diamond growth + +In state-of-the-art CVD diamond deposition, the deposition process is usually initiated by depositing on an iron mandrill or by providing defect sites on a substrate surface. In order to decrease the time necessary to initiate the growth, seeding is often used. This can be accomplished by mechanically abrading the surface using diamond particles to physically generate defect sites. Numerous groups have used this inexpensive approach to accelerate the initiation of diamond growth. Also, small diamond particles, mixed with photoresist, can be uniformly distributed on a substrate surface by spin coating [9]. Once the photoresist is burned off in a plasma, the diamond particles act as nucleation sites for diamond growth. Both of these techniques can be used for small or large area applications. Typically, $10^{7}$ to $10^{9}$ particles $\mathrm{cm}^{-2}$ density can be achieved using these techniques. However, a low nucleation density and/or clumping of diamond particles generally results in material with a significant number of large microcavities in the diamond film. Thus, a need exists for a technique which will yield a very high and uniform nucleation density over the entire surface of the substrate onto which diamond is to be deposited. Such a technique should only be limited by the minimum diamond particle size that can be uniformly deposited onto the substrate surface. + +# 3.2. Large area diamond synthesis + +Two of the most important requirements for any process for depositing large area diamond films are: (1) a high growth rate (at least $5\mu \mathrm{m / h}$ ) and (2) the capability of depositing uniform films over a large area (at least $6.5\mathrm{cm}^2$ ). In the published literature, there are a few techniques being used for the production of thermal management diamond substrates. Some of these techniques are d.c. arc jet CVD, microwave plasma CVD, and hot-filament CVD. Of the other techniques being studied, the most promising approach may be a laser-activated deposition technique. The reported growth rate is about $1\mu \mathrm{m}\mathrm{s}^{-1}$ over a $1\mathrm{cm}^2$ area [10]. + +In our laboratory, we have deposited diamond onto silicon using the microwave plasma CVD technique. The deposition system is made by Wavemat, Inc. and operates at the TM 012 mode at $2.45\mathrm{GHz}$ . In this system, uniform and continuous diamond films have + +been deposited over $75\mathrm{mm}$ silicon wafers using the photoresist seeding technique with $4\mathrm{nm}$ diamond particles. Nucleation densities of $>10^{8}\mathrm{cm}^{-2}$ are routinely obtained. + +# 3.3. Intelligent quality control using optical emission spectroscopy (OES) + +An optical emission spectrum of a microwave plasma showing the different reactive species present during diamond deposition is shown in Fig. 3. The atomic traces from hydrogen and the molecular structure from CH, $\mathbf{C}_2$ , and $\mathrm{H}_{2}$ are clearly visible. However, these transitions only provide direct information for the excited states. To infer information related to the chemically-important ground state densities, a knowledge of the electron energy distribution function (EEDF) is needed. Using argon as an actinometer, we can study the population of ground state atomic hydrogen. Because OES is relatively easy to implement, the determination of in situ correlations between OES and diamond growth parameters is of great interest. Numerical simulations which account for the non-Maxwellian EEDF will facilitate this determination. + +Our present actinometry data show that increasing methane concentrations depress the EEDF and negligibly affect the atomic hydrogen ground state. The effect on the EEDF presumably results from increased electron-molecular vibrational coupling. Our data also indicate that increasing power increases the atomic hydrogen ground state density but only marginally increases the energy of the EEDF. Numerical simulations have been combined with experimental optical emission measurements of free and bound-excited electron number densities to demonstrate the potential of OES as a quantitative diagnostic tool for microwave CVD reactors [11]. + +# 3.4. Cutting and drilling of diamond substrates + +Cutting of CVD diamond substrates is an essential step in the fabrication of thermal management substrates. Free-standing diamond substrates are usually manufactured by deposition over a large circular mandrill and often have a lip (or collar) around the outer edge caused by deposition on the edge of the mandrill. This creates a need for a method of cutting substrates into either circular wafers without the collar or substrates of various other sizes and shapes. However, the cutting of single-crystal gem diamond and polycrystalline CVD diamond is different in several respects. The hardness and chemical inertness of CVD diamond, along with its polycrystalline structure, present a serious challenge to conventional substrate cutting techniques. The problem is even more complicated if one desires to cut thick diamond films deposited on a substrate with + +![](images/366c07ae9c06c588e4a35142e370deeecc4e91c18009ea3493240a5933323d52.jpg) +Fig. 3. OES spectrum of a typical $1.6\mathrm{kW}$ microwave plasma with $\mathrm{H}_{2}$ and $\mathrm{CH}_4$ gases. + +different thermal properties. In fact, although diamond cutting technology is new to the CVD diamond industry, it was developed a few decades ago for cutting single-crystal diamond slabs (not gems). Typically, CVD diamond is cut using a pulsed Nd-YAG laser. The wavelength, power, and repetition rate of the laser, and the chemical quality of diamond, critically determine the definition of the cut, i.e., the cutting-induced burr, the contamination, melting, and refusing of the material, and micro-cracking of the material. Also, the definition of a corner can depend on the proper choice of laser wavelength. Furthermore, the cutting speed of a laser can be enhanced by providing proper photo-chemistry or certain gases at the workpiece. + +The formation of vias for substrate-to-substrate interconnects is another important technology in a series of processing steps necessary for the fabrication of diamond substrates, particularly 3-D multichip modules (MCMs). Vias, which are typically $100 - 150\mu \mathrm{m}$ in diameter, are filled with metal to define metal interconnects through the diamond plane. Depending on the size and design of an electronic system, the number of vias can vary from a few hundred to a few thousand on a single substrate and diamond, due to its hard and chemically-inert nature, presents a formidable challenge to their formation. + +Laser drilling is the only practical approach to the formation of vias in CVD diamond. The challenges are to form very precise diameter vias in a short time with a desirable edge angle, little, if any, wall contamination, and excellent definition of the top surface-hole wall junction. A further complication in state-of-the-art CVD diamond substrates is the existence of internal stresses. Drilling of a substrates changes its stress pattern, and + +thus, the flatness of the substrate. Although, a significant amount of research has been conducted in this area, very little has been published because the techniques used and results are considered proprietary. A most promising approach to this problem consists of drilling the film while it is submerged in a liquid using a focused Q-switched Nd-YAG laser. To date, drilling has been performed in air, de-ionized water, and a few other solutions (Fig. 4) [12-16]. Research on this approach to via drilling of diamond is continuing. + +# 3.5. Lapping and polishing + +The chemical inertness, hardness, polycrystalline nature, and non-uniform surface chemistry of CVD + +![](images/c91c123217ff918c818a985b08e5a82a8309715a94a1a53389010d02167609ba.jpg) +Fig. 4. SEM micrograph of a laser drilled hole in CVD diamond performed under water. + +diamond present serious challenges to the development of surface finishing processes (i.e., polishing and planarization). Techniques which have been tried include conventional mechanical lapping, hot metal polishing, chemical-assisted mechanical polishing, plasma etching, ion beam etching, and laser beam trimming. In general, these approaches can be divided into coarse lapping and fine polishing. For example, mechanical lapping and laser trimming can be used to produce a coarse finish (Fig. 5). Using these approaches, the surface roughness of diamond can be reduced from an $R_{\mathrm{a}}$ of $15\mu \mathrm{m}$ to about $1 - 1.5\mu \mathrm{m}$ as determined by contact surface profilometry. If an electronic packaging application requires a further reduction in the surface roughness, a technique such as chemical-assisted mechanical polishing can then be used for fine polishing. This technique is capable of reducing the average surface roughness from an $R_{\mathrm{a}}$ of $1.5\mu \mathrm{m}$ to about $100 - 150\mathrm{nm}$ as measured by contact surface profilometry [16]. Table 1 provides a comparison of the capabilities of these lapping and polishing techniques. + +As noted previously, a close examination of even finely polished diamond surfaces ( $R_{\mathrm{a}} = 50 \mathrm{~nm}$ by contact surface profilometer) using AFM reveals the presence of surface pits resulting from "microcavities" being intersected by the surface plane. Microcavities have been observed in all CVD diamond material independent of the deposition technique. Often, these surface pits are not detected by contact surface profilometry, but they exist and must be addressed by some post-deposition processing technique. + +# 3.6. Planarization-by-filling + +Basically, the "planarization-by-filling" process consists of filling diamond substrate surface pits (microcavi + +![](images/92a400279ff9b5ecc438a9a396514d79b0c1ef3b70ebacf17c5354d841101bcf.jpg) +Fig. 5. SEM micrograph showing as-deposited and laser trimmed (polished) CVD diamond. + +ties) with a polymer, glass, diamond-filled glass, or similar material [7]. The filler is applied as a thin overlayer on the diamond substrate which planarizes the surface. The overlayer can then be polished back to expose the previously polished diamond surface. However, thermal simulation studies of such planarized diamond surfaces have shown that backpolishing is not necessary from a thermal performance point of view. In fact, this planarization technique offers a simple, conventional approach to the reduction of surface roughness without polishing the diamond surface at all, thereby significantly reducing the time and cost of preparing CVD diamond films for electronic packaging applications. + +Fig. 6 shows an AFM plot of a diamond substrate which was coated with polyimide prior to evaporating and defining an aluminum pattern on the surface. The polyimide is about $5\mu \mathrm{m}$ thick with an average surface roughness of $25\mathrm{nm}$ . Testing of the polyimide overcoat yielded an adhesion strength value of about $2500\mathrm{psi}$ . Although the adhesion strength degraded with thermal shock testing, it was still sufficient for use of the planarized diamond substrate in thermal management applications. Fig. 7 shows a planarized and metallized diamond substrate. + +# 3.7. Metallization + +# 3.7.1. Metalization of large area diamond substrates + +Diamond MCMs require complex metallization systems for defining the ground and power planes, and the signal interconnects. In order to realize these required electrical functions, large area metallization, selective metallization, and metal via filling capabilities must be available. In particular, gold (Au) or copper (Cu), due to their low resistivity, are considered to be the ideal metals for MCM interconnects. Depending on the MCM performance requirements, the required metal thickness can range from a few microns to a few tens of microns, and the required interconnect width can be a few tens of microns. For such requirements, plating is preferred over thermal evaporation or sputtering because it is a simple and economical technique. Large area plating processes have been developed for both gold and copper at the University of Arkansas. + +The adhesion of metallization to diamond in MCMs is a reliability concern. Consequently, the adhesion of Au or Cu to diamond must be enhanced by the use of other metals as seed layers prior to plating. It is important to note that the surface pits discussed previously can cause serious blistering of the plated metal if the plating is not performed properly (Fig. 8). However, "planarization-by-filling" can be used to eliminate the blistering problem. + +Table 1 Cost analysis for post-synthesis finishing of CVD-diamond substrates: estimated costs for various processing technologies + +
Cost-related parameters (for 3" × 3")
LT(8 h day-1operation)CAMPP(8 h day-1operation)Cleaning(8 h day-1operation)FP(8 h day-1operation)
1. Capital cost plus % for installation(% capital cost increase for4" × 4" sample)$75 000:$47 000:$5000:$4800:
laser + table + optics (0%)machine + accessories (27%)glassware (50%)spinner + IR oven + soft bake oven (10%)
2. Type of consumable$120 year-1$4800 year-1$6000 year-1$5000 year-1
3. Maintenance$600 year-1$180 year-1$60 year-1$600 year-1
4. Energy usage100 W2.5 kW80 W60 kW (ovens)
5. Time for processing/sample40 min3 h30 min3.30 h
6. Rigidity of equipment(i.e., vacuum, and critically)Atmospheric process,critical optical alignmentAtmospheric process,critical sample mountingNoneNone
7. Down time to repeat this same process5 min30 min5 min2 min (final process)
+ +LT, Laser trimming; CAMPP, chemical-assisted mechanical polishing and planarization; and FP, filling for planarization. Diamond substrate material specifications: (1) substrate size: $3 \times 3 \mathrm{in}^2$ ; (2) initial $R_{\mathrm{a}} = 3 - 5 \mu \mathrm{m}$ (measured by mechanical surface profilometry); (3) final $R_{\mathrm{a}} = 0.25 \mu \mathrm{m}$ (measured by mechanical surface profilometry); (4) starting thickness, $90 \mu \mathrm{m}$ ; (5) final thickness, $700 \mu \mathrm{m}$ ; (6) bow in the substrate, $\leq 25 \mu \mathrm{m}$ ; (7) chemical quality (in terms of thermal conductivity), $12 \mathrm{W/cm} \cdot \mathrm{K}$ ; (8) primary film orientation, (110); (9) substrate without collar; and (10) grain size cariation, $20 - 25\%$ . + +# 3.7.2. Adhesion/seed layers and metal plating + +Owing to its chemical inertness, diamond is expected to have poor adhesion to most metals, especially to gold - a noble metal best suited for conductor traces because of its excellent electrical conductivity. Thus, there is a need for a metal adhesion/seed layer as noted previously. Fortunately, diamond is known to form strongly adhering carbide layers upon annealing at high temperatures $(700 - 900^{\circ}\mathrm{C})$ with refractory metals such as Ti, W, Mo, etc., and transition metals such as Cr, Ni, Fe, etc. [17]. + +Some of the metal systems examined for potential use as the adhesion/seed layer for diamond-based MCM metallization systems are Au/Ti, Au/Ti-W, Au/Cr, Au/Ni-Cr and Cu/Cr. These metal systems are generally deposited by evaporation or sputtering and then annealed at temperatures ranging from 150 to $500^{\circ}\mathrm{C}$ . Depending on deposition parameters, such as deposition rate, substrate temperature, pressure, etc., adhesion values ranging from about 3 kpsi to greater than 10 kpsi have been achieved [18]. These adhesion values are more than adequate for MCM metallization although they can be increased by sputter etching of the diamond surface prior to deposition of the metals. + +# 3.7.3. Electroplated gold + +Although gold and copper plating are mature technologies that have been used in the electronics industry for decades, bare, polished, MCM grade CVD diamond, with its rough surface due to the presence of surface pits, presents a plating challenge. If plating of diamond is attempted using the parameters specified by plating solution manufacturers, the roughness of the diamond surface is enhanced. For example, in initial attempts at gold plating of diamond at the University of Arkansas, + +the surface roughness increased from $330\mathrm{nm}$ to $2\mu \mathrm{m}$ for a $4\mu \mathrm{m}$ thick gold layer. This surface roughness is unacceptable for MCM substrates. A thorough investigation revealed that the increase in surface roughness had nothing to do with the gold thickness but, in fact, was related to the plating current density. The current density recommended by the plating solution manufacturer was identified as the problem. Upon reducing the plating current density by a factor of 10, the surface roughness actually decreased with increasing gold thickness. Doubling of this lower plating current density maintains the original surface roughness while yielding an acceptable deposition rate of $2\mu \mathrm{m}\mathrm{h}^{-1}$ . A $4\mu \mathrm{m}$ thick layer has a resistivity of about $4.0\mu \Omega \cdot \mathrm{cm}$ which compares favorably with the standard value for gold of $2.2\mu \Omega \cdot \mathrm{cm}$ . + +Large area plating of diamond substrates (10 cm diameter) for MCM applications can be accomplished by evaporating a $\mathrm{Au} / \mathrm{Cr}$ adhesion/seed layer followed by annealing at a temperature of $300^{\circ}\mathrm{C}$ (Fig. 9). The samples are then gold plated to a thickness of several microns and annealed at $300^{\circ}\mathrm{C}$ . This process yields adhesion values greater than 3 kpsi even after the samples are thermally cycled between 150 and $-65^{\circ}\mathrm{C}$ . + +# 3.7.4. Diamond substrate rework + +Because of the relatively high cost of diamond substrates, research has also been performed to establish a viable method for reclaiming them after they have been metallized. Two wet processes and a lapping technique were investigated. A wet process consisting of cerric ammonium nitrate, perchloric acid, and water was found to be the most effective method for reclaiming diamond. EDS showed no indication of chromium on the sample + +
Image Statistics
Z range2.209 μM
Mean684.54 nm
RMS (Rq)293.22 nm
Mean roughness (Ra)141.42 nm
Max height (Rmax)2.077 μM
Surface area
Surface area diff
+ +
Box Statistics
2 range277.36 nm
Mean934.20 nm
Rms (Rq)69.134 nm
Mean roughness (Ra)27.572 nm
Max height (Rmax)166.04 nm
Box x dimension20.706 μm
Box y dimension13.176 μm
+ +![](images/1f2cc7f79c536271065c17d016854c49b03a62d8d635281fdb856b405de748b9.jpg) +Fig. 6. AFM micrograph of a polyimide planarized CVD diamond substrate. + +after an etch time of $15\mathrm{min}$ in the solution. However, care must be taken to ensure that all the etchant is removed from the sample before it is re-used. + +# 3.7.5. Metallization of vias + +Metal filling of substrate vias is obviously important in MCM technologies because of the multi-level nature of the metal interconnection system. Proper via filling is also critically important to multi-substrate communication in a 3-D MCM architecture. For diamond, which has a graphitization temperature between 700 and $800^{\circ}\mathrm{C}$ , low temperature via filling materials such as metal-filled epoxy or plating can be used. Recently, another approach to via filling using tungsten-copper + +![](images/90706b91bc06658099394178fd1e858b6f1d3b509e0e8290512af0336429c66d.jpg) +Fig. 7. Optical photograph of photolithographically-defined aluminum patterns on a polyimide planarized 2.0 inch diamond substrate. + +![](images/396097ff1050b36319f54a1ec84eaf37e08f74337337eb82a1816e8c0335fa33.jpg) +Fig. 8. SEM micrograph of a metallized diamond substrate on which a blister has been deliberately cut open to reveal multiple microcavities. + +composite has been developed using a technique developed by Micro Substrates Corporation [19]. Although there are several other approaches to via filling, the method chosen to fill vias in diamond must be based on other construction aspects of the substrate, subsequent processing steps in the fabrication schedule, and the desired function of the vias. + +# 4. Economics of thermal management diamond substrates + +At the present time, the electronics industry is extremely cautious about using diamond in packaging because of its cost. As with most products, the introduction of diamond thermal spreaders into electronic products will come down to a trade-off between price and + +![](images/9223205169ec1677d7c73f3be9040944af2e4ad39b1dcb1eba538e8bf3cb4862.jpg) +Fig. 9. Optical photograph of a gold plated 4.0 inch diameter diamond wafer. + +product performance. This probably means that the initial approach to incorporating diamond into electronics will be applications where the use of diamond is critical for proper system operation and the cost is not a prime consideration. Although diamond heat spreader prices have dropped dramatically in the past few years, a further substantial decrease in price is needed before the widespread use of diamond will occur. For users of diamond, the price depends on the specific application. Questions which must be answered are; (1) can the diamond be used as grown or must it be polished; (2) must the sample be polished on one side or both sides; (3) what is the required surface finish; (4) what is the required thermal conductivity; (5) what is the substrate quantity required; and (6) what are the desired substrate dimensions, including thickness? These questions are very similar to those which must be answered when ordering substrates made from any other material. A $10\mathrm{cm}$ diameter diamond substrate $1\mathrm{mm}$ thick contains approximately 140 carats of diamond. As-grown diamond can cost anywhere from $100 to$ 3500 per carat, depending on its quality. Post-synthesis processing such as lapping or polishing add to this cost, and they are very expensive processes. The high cost of post-deposition processing is partially a result of the fact that most vendors are using conventional processing techniques. For example, it takes a long time to lap and polish polycrystalline CVD diamond substrates using techniques developed for polishing gem stones. Non-conventional approaches, such as a combination of laser trimming/quick lapping/filling for planarization are capable of reducing post-deposition processing costs by orders of magnitude as suggested by the data in Table 1. + +# 5. Summary and conclusions + +Although technologies for the deposition of thin film synthetic diamond have existed for over a decade, only + +in the past 5 years have research efforts been concentrated on economical deposition and post-deposition processing of large area diamond thin films for application in electronic packaging. Because of its unique properties, diamond is the ideal thermal management material if its cost is competitive with other materials and if technologies are available to process it. During the past 5 years, significant advancements in diamond synthesis have resulted in the deposition of large area substrates and a substantial reduction in cost. The availability of large area material prompted a concentrated effort to develop technologies, such as lapping, polishing, planarization, cutting, drilling, and metallization, which would permit diamond to be used in electronic packaging applications. Although not mature, these post-deposition processing technologies have evolved to the point that as-deposited diamond can realistically be utilized in thermal management applications. However, a further reduction in cost of both diamond synthesis and post-deposition processing will be necessary before diamond finds widespread use in commercial electronic packaging. + +# References + +[1] A.P. Malshe, S. Jamil, M.H. Gordon, H.A. Naseem, W.D. Brown and L.W. Schaper, Advanced Packaging, September/October 1995, pp. 50-53. +[2] J.E. Butler et al., Phil. Trans. R. Soc. Lond., A 342 (1993) 209-224. +[3] A.P. Malshe, unpublished work at University of Arkansas, Fayetteville. AR, 1995. +[4] A.P. Malshe, G.J. Glezen, H.A. Naseem and W.D. Brown, Proceedings of the Fourth International Symposium on Diamond Materials, The Electrochemical Society, Pennington, NJ, 95-4 (1995) 515-522. +[5] D.G. Bhat, D.G. Johnson, A.P. Malshe, H.A. Naseem, W.D. Brown, L.W. Schaper and C.-H. Shen, Diamond Relat. Mater., 4 (1995) 921-929. + +[6] G.J. Glezen, Masters thesis, University of Arkansas, Fayetteville, AR, 1995. +[7] A.P. Malshe, W.D. Brown, H.A. Naseem and L.W. Schaper, US patent No. 5,472,370, 1995. +[8] J. Grabner, in K. Azar (ed.), in Handbook of Experimental Methods in Electronic Cooling, CRC Press, in press, 1996. +[9] R.A. Beera, M.S. Haque, W.D. Brown, A.P. Malshe and H.A. Naseem, Proceedings of Topical Conference on Synthesis and Processing of Materials, AIChE 1994 Annual Meeting, San Francisco, CA, 1994, pp. 65-70. +[10] Pravin Mistry, QQC, Inc., Workshop on Characterizing Diamond Films IV, NIST, Gaithersberg, MD, March 4-5, 1996. +[11] U.M. Kelkar and M.H. Gordon, Electrochem. Soc. Proc., 96-5 (1996) 75-82. +[12] S. Jamil, M.H. Gordon, G.J. Salamo, H.A. Naseem, W.D. Brown and A.P. Malshe, Proceedings of Applications of Diamond Films and Related Materials: Third International Conference, NIST special publication 885, 1995, pp. 283-286. +[13] T. Chein, C. Cutshaw, C. Tanger and Y. Tzeng, Proceedings of + +Applications of Diamond Films and Related Materials: Third International Conference, NIST special publication 885, 1995, pp. 257-260. +[14] P. Tosin, W. Luthy and H.P. Weber, Proceedings of Applications of Diamond Films and Related Materials: Third International Conference, NIST special publication 885, 1995, pp. 271-274. +[15] V.G. Ralchenko, S.M. Pimenov, T.V. Kononenko, K.G. Korotoushenko, A.A. Smolin, E.D. Obraztsova and V.I. Konov, Proceedings of Applications of Diamond Films and Related Materials: Third International Conference, NIST special publication 885, 1995. pp. 225-232. +[16] A.P. Malshe, unpublished work at University of Arkansas, Fayetteville, AR, 1995. +[17] K.L. Moazed, J.R. Zeidler and M.J. Taylor, J. Appl. Phys., 689(5) (1990) 2246-2254. +[18] W.D. Brown, H.A. Naseem, A.P. Malshe, J.H. Glezen and W.D. Hinshaw, Mater. Res. Soc. Symp. Proc., 391 (1995) 59-70. +[19] W.E. Wesolowski, R.J. DeKcnipp and M.P. Ram Panicker, Proc. 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0000000..e4c2b34 --- /dev/null +++ b/金刚石散热项目/项目前期工作_MD/01_CVD金刚石薄膜与热扩散片/4_Fabrication of free-standing diamond membranes/4_Fabrication of free-standing diamond membranes.md @@ -0,0 +1,93 @@ +# Fabrication of free-standing diamond membranes + +M.C. Salvadori a, M. Cattani a, V. Mammana a, O.R. Montciro b, J.W. Ager III b, I.G. Brown b + +$^{\text{a}}$ Institute of Physics, University of São Paulo C.P. 66318, São Paulo, SP, CEP 05389-970, Brazil + +b Lawrence Berkeley National Laboratory, University of California Berkeley, CA 94720, USA + +# Abstract + +We describe here a method for fabricating free-standing diamond membranes. Diamond films were deposited on a silicon substrate by microwave plasma-assisted chemical vapor deposition and then part of the substrate chemically removed. The films described here were 15 mm in diameter with thickness of approximately $12\mu \mathrm{m}$ . A novel feature of our approach lies in the method used to obtain the selective dissolution of the substrate; a container with O-rings was used, instead of masks, allowing a fast and clean isotropic dissolution of part of the silicon substrate. The deposited diamond films as well as the free-standing membranes were characterized by scanning and transmission electron microscopy, electron diffraction and Raman spectroscopy. + +Keywords: Diamond; Chemical vapour deposition; Plasma processing and deposition + +# 1. Introduction + +Diamond thin films possess a number of unique chemical and physical properties that make them attractive for a broad range of applications in research and industry. The rapidly developing technology for the growth of polycrystalline diamond films by plasma-assisted chemical vapor deposition (CVD) on a range of substrates has increased interest in this material. Applications of diamond thin films include: wear-resistant coatings for cutting tools, by virtue of its hardness; as thermal management of electronic devices, because its room-temperature thermal conductivity is the highest of all materials; as free-standing windows with good transparency in the visible and infrared region; and as free-standing permeable membranes [1] for use as filters, by virtue of the chemically inert and mechanically resistant nature of diamond. Several CVD methods have been developed and successfully used to produce diamond films such as, for example, hot filament CVD [2], electron-assisted CVD [3], microwave plasma-assisted CVD [4], d.c. discharge CVD [5], and use of the oxygen-acetylene torch [6]. + +In this paper we describe a method for fabricating freestanding diamond membranes from a film deposited on silicon. Selective chemical etching of the substrate promotes dissolution of the silicon leaving the diamond membrane intact. The method is particularly attractive because of its simplicity, ease for producing large area membranes, and its low cost. + +# 2. Experimental details + +The diamond films were synthesized by microwave plasma-assisted CVD using a system that has been fully described elsewhere [7]. The growth parameters used here were: 300 sccm for the hydrogen flow rate, 1.5 sccm for the + +![](images/a60d42d2644182647aa1d9eb912e3e39b5d20adada89864218da620546ec9de3.jpg) +Fig. 1. Container with O-rings for dissolution of part of the silicon substrate. The sample was loaded on the lower O-ring with the diamond side down, then the upper part of the container was screwed tight so as to sandwich the sample between the two O-rings. + +![](images/0f0cc0fb2971d3cc3d7d47e2fc0368f099da108a5d052a12f0194a2390921ee5.jpg) +Fig. 2. The circular region inside the square $(18 \times 18 \mathrm{~mm}^2)$ silicon substrate is a free-standing diamond membrane. + +![](images/04d0d18a95ec6d2bd9d13bf0f48453738c6f72e065bf9b5502f7e527ad704a19.jpg) +Fig. 3. SEM of the top side of the diamond membrane showing morphology typical of CVD-grown polycrystalline diamond film. + +![](images/43b3c40feb9e1ea72c6916d9d24f8f879befbb60a492d767e2ba1f518bd629bf.jpg) + +![](images/09e8b91d711c6a105d5adbcaa0a7a12f9d55c32daf706c6b1f0e38252e3dabf5.jpg) + +![](images/97010ad919752006bbdf4c6f4773656206f83fcafa83f935d7233e57cec70924.jpg) +Fig. 4. (a) Bright-field transmission electron microscopy image of a diamond grain with normal parallel to [110] direction. Twin bands are indicated in the picture, and their signatures indicated in the diffraction pattern. (b) Electron diffraction pattern of the grain shown in (a). (c) High-resolution transmission electron micrograph of a CVD diamond membrane. + +methane flow rate (0.5 vol.% methane), $9.3 \times 10^{3} \mathrm{~Pa}$ for the chamber pressure, $850^{\circ} \mathrm{C}$ for the sample temperature, and a nominal $870 \mathrm{~W}$ for the microwave power. Film thickness was controlled by variation of the growth time between 10 and $48 \mathrm{~h}$ . + +A novel feature of our approach lies in the method used to obtain the selective dissolution of the substrate: a container made of PVC, with Buna-N O-rings, as shown in Fig. 1, was used instead of masks, allowing a fast and clean isotropic dissolution of part of the silicon substrate. The sample (diamond film on its silicon substrate) was loaded on the lower O-ring with the diamond film down. Then the upper part of the container was screwed tight so as to sandwich the sample between the two O-rings. A room-temperature mixture of hydrofluoric acid, nitric acid and acetic acid [8] was used to dissolve the silicon. The HF:HNO₃ volume ratio was 2:1, and the acetic acid was added to prevent violent reaction that can fracture the diamond membrane during the substrate etching. The time required to dissolve the substrate was between 10 and 30 min for a silicon substrate of thickness of about 200 μm. + +Scanning and transmission electron microscopy were used to characterize the deposited diamond films as well as the free-standing membranes. A Jeol JSM-840A scanning electron microscope (SEM) with an accelerating voltage of 25 kV was used to evaluate the surface morphology of the films, as well as the etch profile of the silicon substrate. Transmission electron microscopy (TEM) was conducted in a TOPCON 002B. At 200 kV the point resolution of this microscope is 0.19 nm, which is suitable for viewing [111] lattice fringes of diamond. Samples for TEM were prepared by mounting pieces of thin free-standing membranes (5 μm thick or less) on copper grids and ion milling the films with an Ar⁺ beam at 3.5 kV until perforation occurred. Raman spectroscopy was also carried out to provide additional information on the quality of the membranes produced. + +# 3. Results and discussion + +We have in this way made a number of thin diamond membranes with the central $15\mathrm{mm}$ diameter substrate free. Note that the diameter of the free-standing region is determined simply by the O-ring size, which can readily be increased so as to prepare free-standing films of arbitrary size. + +An example of one of our diamond membranes is shown in Fig. 2, where the diamond film lies within the circular region and the square silicon substrate is $18 \times 18 \mathrm{~mm}^2$ . Fig. 3 shows a SEM of the top side (growth side) of the diamond film, showing morphology that is typical of polycrystalline diamond films. + +The surface morphology of the deposited films consisted mainly of (111) and (100) facets. Transmission electron microscopy also indicated the presence of stacking faults, and micro- and macro-twins, which have been often observed in CVD diamond films. Fig. 4(a) shows a transmission electron + +micrograph of one of the membranes produced with one of the grains oriented with the electron beam parallel to the [011] direction. Observation of stacking faults and twin domains is most easily done under this conditions. In this orientation the stacking faults and twin domain boundaries, which coincide with {111} planes, are seen edge-on, and their signature on the diffraction pattern is easily identifiable. Fig. 4(b) shows the electron diffraction pattern of the grain shown in Fig. 4(a). Diffraction spots resulting from multiple twinning, and streaks resulting from stacking faults or microtwins are observed. The high resolution image of another crystal of the diamond m r rane under similar orientation is also shown in Fig. 4(c). Multiple twinning and stacking faults are easily identified. + +The Raman spectra of the two sides of one membrane are presented in Fig. 5. These spectra are consistent with that of diamond with very low amorphous carbon phase contamination. It is interesting to note that there is slightly more amorphous carbon at the silicon side than on the growth side. + +Fig. 6 shows the back side of the sample, where the diamond film was in contact with the silicon substrate before its chemical dissolution; one can see in this micrograph that the diamond grain size is about $4\mu \mathrm{m}$ . Fig. 7 shows the back side of the diamond film where it adjoins the thicker silicon substrate; the smooth region on the right hand side of the figure + +![](images/dced7f6e09c5325976707792cc4d3a7c2aec48f3e20a3742ea4f3b808bb9e735.jpg) + +![](images/20ccda04107868fa9fba05d43294f780a45a58c5d43848ac207ade5c5de7bc1f.jpg) +Fig. 5. Raman spectra of both sides of a free-standing diamond membrane. + +![](images/cebfed6b6ba6f5eb3cbc411f19f193ca984f80ed39d66affc4be2e02f7fb4944.jpg) +Fig. 6. SEM of the back side of the diamond membrane, where the diamond film was in contact with the silicon substrate before its chemical removal. The diamond grain structure is clearly shown. + +![](images/236a787687bf2010c78b5a903a8cd876252e161fa0aee4ac9dad36c6ca73f5af.jpg) +Fig. 7. SEM of the back side of the membrane near the edge. The smooth part (right) is the free-standing diamond film and the rough part (left) is the silicon substrate. + +is the free-standing diamond film and the rougher part on the left side is the silicon substrate. Note that the silicon profile near the free-standing diamond film is sharp, in spite of the fact that the chemical dissolution was isotropic. The film + +thickness of about $12\mu \mathrm{m}$ was determined by breaking a film grown under similar condition and measuring it under SEM. + +# 4. Conclusion + +We have fabricated a number of diamond membranes by selective dissolution of the silicon substrate using a container with O-rings, instead of masks, allowing a fast and clean isotropic removal of a chosen part of the substrate. The test diamond membranes, made in the work described here, were circular with a diameter $15\mathrm{mm}$ . The size of the free-standing region can readily be increased. + +# Acknowledgements + +This work was supported by the Fundação de Amparo à Pesquisa do Estado de São Paulo and the Instituto de Física da Universidade de São Paulo, Brazil. We would also like to acknowledge the support and use of the facilities of the National Center for Electron Microscopy at the Lawrence Berkeley National Laboratory; the NCEM is supported by the Director, Office of Basic Energy Sciences, Materials Science Division, US Department of Energy under Contract No. DE-AC03-76SF00098. + +# References + +[1] M.C. Salvadori, Y. Miyao and G. Moscati, Diamond Related Mater., 4 (1995) 1069. +[2] A.M. Bonnot, Thin Solid Films, 185 (1990) 111. +[3] A. Sawabe and T. Inuzuka, Thin Solid Films, 137 (1986) 89. +[4] M.A. Brewer, I.G. Brown, M.R. Dickinson, J.E. Calvin and M.C. Salvadori, Rev. Sci. Instrum., 63 (1992) 3389. +[5] K. Suzuki, A. Sawabe, H. Yasuda and T. Inuzuka, Appl. Phys. Lett., 50 (1987) 728. +[6] L.M. Hanssen, W.A. Carrington, J.E. Butler and K.A. Snail, Mater. Lett., 7 (1988) 289. +[7] M.C. Salvadori, V.P. Mummann, O.G. Martins and F.T. Degasperi, Plasma Sources: Sci. Technol., 4 (1995) 489. +[8] W.R. Runyan and K.B. 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JAGANNADHAM + +Materials Science and Engineering, North Carolina State University, Raleigh, NC 27695-7916, U.S.A. + +(Received 3 March 1998; accepted 4 June 1998) + +AbstractÐSingle layer diamond and multilayer diamond heat spreader substrates are prepared and bonded to device wafers of silicon and gallium arsenide. Metallization schemes for the diamond surface and the backside of the device wafers are described. Bonding of the device wafers to the diamond substrates using the high thermal conductivity gold±tin eutectic solder is carried out. Characterization of the bond for the distribution of di€erent elements in the metallization layers and the solder, for the presence of microscopic defects such as voids and cracks, for the adhesion strength and for the stability of the bond under thermal cycling is performed. The heat spreader characteristics of the substrates with single and multlayer diamond are determined using infrared imaging of the bonded device wafers and compared with that of wafers bonded to metal substrates. Modeling and analysis of the e€ective thermal conductivity showed that the multilayer diamond substrates are better heat spreaders and reduce the device temperature so that the life of the electronic devices is prolonged. # 1998 Published by Elsevier Science Ltd. All rights reserved + +# 1. INTRODUCTION + +Heat spreading applications in electronics packaging, hitherto, are handled by high thermal conductivity materials such as beryllium oxide. Recent progress towards microminiaturization and high packaging densities leading to higher density of devices and modules as well as introduction of high frequency and high power devices compounded the problem of heat dissipation. It is realized that more ecient and fast dissipation of heat generated by the devices can only be met by the use of diamond ®lms. + +Power semiconductor devices rely upon the e- cient removal of heat for high frequency switching[1±3]. Natural diamond with high thermal conductivity (20 W/cm K), high electrical resistivity $( 1 0 ^ { 1 6 } \Omega \mathrm { c m } ) ^ { \dag }$ , low dielectric constant (5.7), high dielectric strength (106 V/cm), and matching thermal expansion coecient has been used in bonding devices such as laser diodes[4,5] to dissipate the thermal load. The advent of low pressure synthesis of diamond has helped to replace the high cost natural diamond. It is increasingly realized that the high thermal conductivity of synthetic diamond (10 W/cm K) by itself is not enough to dissipate the heat as the heat capacity of diamond is not large. Therefore, the electronic packaging industry can use the diamond heat spreaders provided the heat is absorbed by a high heat capacity substrate such as molybdenum or silicon nitride which can in turn be cooled by solid, liquid or vapor coolants. The important requirements that should be met before diamond coated heat spreaders can be used + +successfully include good adhesion of diamond to the substrates, stability of the bonded structure to thermal cycling, good heat spreader characteristics of the multilayer bonded structure, reliable process and bonding technology. + +We have developed the multilayer synthetic diamond coating technology on two types of substrates: molybdenum metallic substrate or silicon nitride ceramic substrate. The metallization and bonding procedure that provides reliable bonding to the device wafers such as silicon and gallium arsenide are also established. Characterization of the bonded microstructure, thermal stability and heat spreader properties are presented in the following. + +# 2. EXPERIMENTAL PROCEDURE + +Single layer continuous diamond ®lm of thickness 6 mm or multilayer diamond ®lm of thickness 10 mm was deposited on molybdenum or silicon nitride substrates of thickness 3±4 mm. The diamond layer in these ®lms was deposited by hot ®lament chemical vapor deposition. The details of diamond deposition were described in detail in our earlier work[6,7]. Single layer diamond ®lm deposited on molybdenum for 8 h to achieve a thickness of 10- 14 mm was found to delaminate from the diamond substrate. The large thermal stresses resulting from di€erences in thermal expansion coecient between diamond $( \alpha = 3 . 1 \times 1 0 ^ { - 6 } / \mathrm { K } )$ and molybdenum carbide $( \alpha = 5 . 8 \times 1 0 ^ { - 6 } / \mathrm { K } )$ formed on molybdenum $( \alpha = 4 . 9 \times 1 0 ^ { - 6 } / \mathrm { K } )$ substrate were responsible for the delamination. The results of residual stress + +analysis using X-ray di€raction and Raman spectroscopy are described more completely in a recent work[7]. The diamond ®lms grown on silicon nitride $( \alpha = 2 . 3 \times 1 0 ^ { - 6 } / \mathrm { K } )$ were not susceptible to delamination as a result of matching thermal expansion coecients[8,9]. The multilayer diamond coated substrates consist of a ®rst discontinuous layer of diamond deposited for 4 h under the same conditions as in the single layer diamond. In order to improve the adhesion of this ®rst layer of discontinuous diamond, an intermediate or second layer consisting of aluminium nitride of 1±2 mm thickness was deposited. The choice of aluminium nitride was based on the reasonably large value of its thermal conductivity (3.7 W/cm K), matching thermal expansion coecient $( 4 . 1 \times 1 0 ^ { - 6 } / \mathrm { K } )$ and high dielectric constant (8.8). In addition, the ease of deposition of high quality aluminium nitride ®lms was a contributing factor. Deposition of aluminium nitride was carried out either by pulsed laser ablation (KrF at 248 nm) or reactive magnetron sputtering[10] with the substrate maintained between 600 and 6508C. The experimental details of deposition of AlN are described in a recent publication[10]. Silicon carbide with a high thermal conductivity $( 5 \mathrm { { W / c m } \mathrm { { K } ) } }$ , matching thermal expansion coecient $( 4 . 3 \times 1 0 ^ { - 6 } / \mathrm { K } )$ and a high dielectric constant (9.0), although a good candidate for the embedding intermediate layer in the composite, cannot be deposited with the same good crystalline quality at these low temperatures. A ®nal layer of continuous diamond was deposited on the top for 6±8 h to provide a total thickness of 10±12 mm of diamond ®lm. Figure 1 shows a schematic illustration of the multilayer diamond ®lm bonded to a device wafer. The multilayer diamond ®lms on both molybdenum and silicon nitride substrates were found to be strongly adherent. Raman spectroscopy was used to characterize single layer and multilayer diamond ®lm and the results showed the characteristic diamond peak with the absence of either graphite or diamondlike phase[6,7]. X-ray di€raction was used to determine + +![](images/b46036063af7ef5b9d0c6acfb9c00d608258902d8496a3e3659377b3a3e50735.jpg) +Fig. 1. Schematic illustration of the structure of multilayer diamond heat spreader bonded to a device wafer. + +the crystalline quality of the aluminium nitride phase. + +Ecient heat dissipation from the devices was made possible by forming a solid state bond with the diamond substrate. The thermal resistance o€ered by the bond should be minimized at the same time the strength and thermal stability of the bond should be maintained. In order to enable the bond formation, metallization of diamond ®lm surface was carried out by deposition of titanium, gold and copper layers by laser physical vapor deposition. The multitarget deposition chamber was evacuated to a predeposition vacuum of better than $1 0 ^ { - 7 }$ Torr to prevent titanium oxide formation. Titanium was deposited at a substrate temperature of 6008C initially to prevent diamond decomposition to graphite followed by quickly heating to 7008C. Deposition of gold and copper ®lms was continued while cooling the substrate to room temperature. The thickness of each layer was estimated to be 0.3±0.5 mm. The deposition of copper layer, because of the high re¯ectivity, was carried out for 40 000 pulses to achieve a thickness of 0.5 mm. The choice of titanium to form the carbide with diamond, that of gold to prevent reaction of tin in gold±tin eutectic solder with titanium and that of copper to form a low temperature solid state bond with the solder was also found suitable for relaxation of thermal stresses during thermal cycling of the bond[9,11]. Similarly, the metallization of the backside of silicon wafer of thickness of 500 mm was carried out by deposition of titanium, gold and copper each for 20 000 pulses at room temperature. On the other hand, a di€erent scheme for metallization of GaAs wafer was used which consisted of deposition of gold followed by that of Cu±Ge alloy by laser ablation with the wafer substrate maintained at room temperature. This metallization procedure in case of thin GaAs wafer was chosen so that the thermal stresses developed during cooling of the bond are relieved by plastic relaxation and secondly, wetting of the solder with the device wafers is accomplished[6]. The thickness of GaAs wafer was only 100 mm and yet the metallization and bonding procedure was carried out successfully without failure of the thin device wafer. + +The device wafer was bonded to the diamond substrate by heating with the low melting point (2808C) gold±tin eutectic alloy solder ®lm, placed in the middle, to a temperature of 3608C in a hydrogen atmosphere of 5 Torr followed by slow cooling to room temperature after 3 min. Spring loaded setup was used to apply small pressure so that a thin ®lm of the bond was formed. The high thermal conductivity and smaller thickness of the bond were important features that improve the e€ective thermal conductivity of the diamond heat spreader. Although gold±tin eutectic solder is not very ductile at lower temperatures, as will be seen in the follow- + +ing results, the ability to wet the copper layer helped the bond formation. + +In order to test the e€ective thermal conductivity and the heat spreader characteristics of the bonded single layer and multilayer diamond substrates, platinum resistance heater was deposited by laser physical vapor deposition on the front side of silicon wafer prior to bonding. Figure 2 is an infrared image of the platinum resistance heater after a constant power input. Calibration of the IR images enabled the temperature changes to be measured with time along the complete wafer for di€erent ®xed values of power input into the heater[12]. + +# 3. EXPERIMENTAL RESULTS + +The bonded specimens were characterized for adhesion strength, uniformity of composition of the bond, resistance to cracking or delamination under thermal cycling, and heat spreader characteristics or e€ective thermal conductivity. The strength of the bond was tested using a pull test[6] on a prototype silicon substrate on which diamond ®lm was deposited. The strength of the bond was found to be 5.0 MPa with the delamination taking place along titanium/diamond interface. Formation of titanium carbide is responsible for the adhesion strength of the metallization layers with diamond. Higher substrate temperature for deposition of titanium is expected to improve the formation of higher quality titanium carbide. The distribution of di€erent elements present in the bonded wafers was determined by X-ray analysis and mapping of the crosssection samples prepared after thermal cycling. Figure 3 shows the X-ray maps of distribution of Si, Au, Sn, C, Al, Ti and Cu and scanning electron microscopy (SEM) image of a cross-section specimen of silicon wafer bonded to diamond/AlN/diamond/Si3N4 substrate prepared after thermal cycling between 25 and 1508C. The absence of silicon and carbon in the solder region helps to show + +![](images/16a6d4a90ca6d5a517654aa610920d0bfd6d35990dbc325bdff3161e183bf719.jpg) +Fig. 2. Infrared image of the resistance heater and the device wafer bonded to the diamond heat spreader. + +that di€usion of these two elements is con®ned. Di€usion of copper into the device wafer is prevented by the presence of gold and titanium layers. Also, the irregular interface between titanium, diamond (carbon) and the substrate seen in the SEM image is an artefact of cutting and irregular features seen on the surface. The signal from diamond was not very strong since carbon is a low atomic number element and in addition, the cross-section specimen was not polished to expose the diamond but more importantly to reveal the presence of voids or cracks in the bonded region that might have formed during thermal cycling. We have not observed the presence of any microscopic defects indicative of failure by cracking or delamination. Figure 4shows the distribution of di€erent elements in a cross-section specimen of GaAs wafer bonded to multilayer diamond on molybdenum prepared by angle polishing at 118 to the plane of the bonded interface. As a result the interface region could be seen magni®ed with delineation of di€erent regions. Similar to the results shown in Fig. 3, the distribution of di€erent elements is con®ned with no interdi€usion into the device GaAs wafer. A more important result common to both Figs 3 and 4 is that the metallization layers of titanium, gold and copper are distributed uniformly in the solder region which is expected since the di€usion in the liquid phase $( 1 . 0 \times 1 0 ^ { - 4 } \mathrm { c m } ^ { 2 } / \mathrm { s } )$ is much faster and the solder layer is only a few micrometers in thickness. + +Thermal cycling was performed in two temperature ranges on di€erent bonded specimens. In the ®rst, each cycle consisted of slow heating to 1508C from room temperature and cooling back in an atmosphere of nitrogen. Another set of specimens were subjected to the second range of thermal cycling wherein the specimens were rapidly heated to 1508C and quenched to ÿ208C in air. The specimens were examined by optical and scanning electron microscopy for delaminated regions either along diamond/substrate interface or diamond/ solder interface and for the presence of cracks in the wafer. The bonded specimens tested by slow cycling were very good without cracking or debonding. However, a small fraction of specimens of silicon wafer bonded to diamond/silicon nitride and subjected to rapid cooling to 208C showed delamination after several cycles. Figure 5 is the cross-section SEM image with X-ray maps of the silicon wafer bonded to diamond on silicon nitride which exhibited upon quenching resistance to delamination. Figure 6, on the other hand, shows the SEM image and X-ray maps of the planar region of the bond that failed by delamination. The interface in Fig. 5 was very good without the presence of any cracks and voids[9]. The distribution of all the elements in the X-ray maps shown in Fig. 6 was uniform and, in particular, there was no segregation that could be responsible for delamination. The high strain rate to which the specimens are sub- + +![](images/3b9c312c6df3a1085b57bdc34b69524c33a9c4d78843e535f0a13216515b5646.jpg) + +![](images/274bb3a351818c19a2265cd5844ecb71db721e5350de6f0fe36d6b8391e1e3e7.jpg) + +![](images/0839709bfe8db58b613d7f024bc2e0ca48455317d8963235c0a6acbb4d8683d5.jpg) + +![](images/ec5e1b4cf5f48777f8525c39a358b3e58fdb2d1c6258994a59e2ec9362b7c289.jpg) + +![](images/657567d7d646dd194f536c569b190ff233d1f1ef8aa2d831e84715f2ee2b23d2.jpg) + +![](images/bf58ab86cfc2e3cbd25c283356df941b7d7e7c639eaa00b61c47e2d4d036ebaf.jpg) + +![](images/773ee72e9c9872bc775775904dc6fb4bc145c1275cc8195d0176db8e74897d5d.jpg) + +![](images/a59f5d64dbb3e1116e344eea223d800ae6a1ded1b871bd1a0b42f82df0b6bcaa.jpg) +Fig. 3. X-ray mapping of di€erent elements and SEM image of the cross-section of silicon wafer bonded to silicon nitride after thermal cycling between 25 and 1508C. + +![](images/1083a390430b1a30723f52fa3fe92653ac6cf246f9ab033effe2f786c88c0fd7.jpg) + +![](images/04c32dfb34949f3bd5379b646bf123f1dd1603f7e5bef5fc53179b2572e46f4c.jpg) + +![](images/fe097e54e002d0c87ab68006e88321d6f2fe7a3a350d58fbd9cc09746bc1cbda.jpg) + +![](images/c3728430226387317041b4b499fd44937d533609676344eca8e6990c18d479af.jpg) + +![](images/e456daefa4d5208f4969a8a860ffe51c8d4d64edd7914fce105d379b5d21aa3e.jpg) + +![](images/6d9d998973c45dfab94a52f6ef0d35637634d7454f6e12e27fb099c5f59eb6d6.jpg) +Fig. 4. X-ray mapping of di€erent elements and SEM image of the cross-section of GaAs wafer bonded to molybdenum after thermal cycling between 25 and 1508C. + +![](images/1b953b39d112b1f216627cf93455aff665bd5b6f2abd61e876c255d8ced6f2f7.jpg) + +![](images/ad8d8b50b0ac9fa7f4e8f4c8adf85e9e618d83280fb082705f8e3a4a4b8dfaf1.jpg) + +![](images/d00ef30835ce1013724f285e686427db0d1eaabf91b1c970659bf2dbcad00ce1.jpg) + +![](images/08dd74120392f09a729944179f842c8b202318a6fc861a53a01a5b3d27eab4ca.jpg) + +![](images/44c3ac0e2bf56498ef1e9e63f07ea0c98c340e4833b9aa98ac99319f84186a14.jpg) + +![](images/9bc436e135c118bbf79fa4ef5e13875d54ef7c1ba21e3cb00a12695f21db0057.jpg) +Fig. 5. X-ray mapping of di€erent elements and SEM image of the cross-section of silicon wafer bonded to silicon nitride after rapid thermal cycling between ÿ20 and 1508C. + +jected during quenching may be responsible for failure along voids or cracks that could be present in the solder region. The brittle nature of the solder region may also be a contributing factor. Introduction of excess gold in the eutectic solder is expected to improve the ductility and the toughness of the solder and thereby prevent failure during quenching. + +The temperature raise on the wafer was measured along a line perpendicular to the platinum resistance heater for di€erent values of power input. The rate of increase of temperature per unit power input was determined and shown in Fig. 7 for the three types of substrates, namely, bare molybdenum, molybdenum with single layer diamond, and molybdenum with diamond/AlN/diamond. These results clearly show that diamond heat spreaders perform much better than the bare molybdenum heat spreader and in addition, the multilayer diamond is better than the single layer diamond. Thus, AlN embedding layer has improved the e€ective thermal conductivity of multilayer diamond heat spreader[12]. + +# 4. DISCUSSION + +The adhesion strength of the bond is found to be high in the device wafers bonded to diamond on molybdenum or silicon nitride substrates. X-ray maps of di€erent elements showed the uniform distribution of di€erent elements in the metallization and solder regions. Thus, the wetting reaction of + +solder with the device wafer and the diamond substrate was good without interdi€usion of either carbon or silicon in case of silicon wafer and gallium or arsenic in case of GaAs wafer into the solder. Similarly, interdi€usion of copper or gold into the device wafers was also absent. Di€usion lengths of di€erent elements in liquid solder within the bond processing period of 3-5 min is estimated to be 1- 2 mm which is considerably larger than the thickness of solder region. Therefore, the metallization layers and the solder region have become a single region with small gradients in concentration. Titanium carbide present across the interface with diamond provides the adhesion strength of the metallization layer to the substrate. The low conductivity of the metallization layers and that of the carbide layer is partly responsible for reduction of e€ective thermal conductivity associated with diamond heat spreader. However, the high heat capacity of these layers compensates for the low value of diamond. Thermal stresses in the thin GaAs device wafers should be reduced to prevent failure of the wafer by cracking. The plasticity of the metallization layer of copper±germanium on the backside is found to improve the toughness of the bond and prevent cracking of the thin (100 mm) GaAs wafer. Failure of the silicon wafers bonded to diamond substrate by delamination upon quenching below room temperature may be associated with phase transformation of tin in the tin-rich solder regions around 138C and the accompanying large volume change (27%). Although no detection of + +![](images/00d1e9090bf759e9c9b18d65333fd5639cf46de290754086790fda10b46aa0f1.jpg) + +![](images/c8ee843d03fe9f24346782ba6e387272f7e027187e542880b91839662af60f20.jpg) + +![](images/99de9d132fa71ce9b3eadada3b7263c490133e8f840ef6302e4cd6a3cc8e36c8.jpg) + +![](images/3b9ffe52b88261b7700944b605f02e61ae078bed308587cbf8a7424695512614.jpg) + +![](images/b02119e0433923b875ce9d366f622d53acc55de982fa4cb9172b8c71e902563c.jpg) + +![](images/7c9a727df5e6a9305f5c2f094559b6dd5d109a8406ce5dd3efdafe0a021dc940.jpg) + +![](images/1618cde7ae816f5b30442db4f74231ce5673c4cf2873846e5311ef0b93d63646.jpg) + +![](images/6281350992d6e77868011f71c15f4c72d878bb3e3859a5e9e2cb381f3333aca3.jpg) + +![](images/c09c84b43d28785c5b6ab8c62db4faaf081b15c559aae4ea5f7d93faace5611b.jpg) +Fig. 7. Results of infrared imaging of silicon wafers bonded to bare molybdenum, single layer diamond coated molybdenum and multilayer diamond coated molybdenum substrates. Temperature at a ®xed point measured along the line perpendicular to the heater is plotted as a function of time. Average slope corresponds to that for three input power settings. For each substrate, the values of the slope differed only in the ®fth decimal place for the three power input settings[11]. + +free tin or segregation of tin is observed in the SEM micrographs of the planar specimens shown in Fig. 6, we believe that there may be small amounts of tin responsible for the large volume change. Also, tin-rich solder may be more sensitive to high strain rates obtained by rapid cooling. We will use solders that are richer in gold to prevent delamination[13]. + +The heat spreader characteristics of the multilayer diamond were found to be better than that of single layer diamond. Single layer polycrystalline diamond deposited on any substrate is known to contain microscopic voids and amorphous diamondlike carbon regions between the grain boundaries. These regions of discontinuity in heat transfer reduce the e€ective thermal conductivity. The e€ective thermal conductivity of diamond/AlN/diamond has been modeled using the multilayer arrangement shown in Fig. 8 with heat ¯ow in the direction normal to the interface[14]. Three types of interfaces between AlN and diamond are possible. These are a sharp interface, a graded interface and a di€used interface. A sharp interface is de®ned by a discontinuous change in thermal conductivity from that of diamond to that of AlN. A graded interface is obtained when the thermal conductivity across the interface changes linearly from that of diamond to that of AlN across a region of width d of the interface. A di€used interface, on the other hand, consists of a change in the thermal conductivity from that of diamond to that of AlN but not necessarily linearly and may also have a minimum or a maximum across the interface, as shown in Fig. 8. A graded or di€used interface forms when the two adjacent phases interact to form a region of varying compo- + +sition. The presence of a maximum can only be fortuitous and has never been realized. The e€ective thermal conductivity, $K _ { \mathrm { e f f } } ,$ of this composite region may be obtained from[15] + +$$ +1 / K _ {\mathrm {e f f}} = V _ {\mathrm {d}} / K _ {\mathrm {d}} + V _ {\mathrm {a}} / K _ {\mathrm {a}} + 2 V _ {\mathrm {i}} / K _ {\mathrm {i}}, \tag {1} +$$ + +where the subscript d indicates diamond, a represents aluminium nitride, and i represents the interface formed on either side, as shown in Fig. 8. In the above equation the V represents the volume fraction and K the thermal conductivity of each phase and for the model shown in Fig. 8, the volume fraction is proportional to the dimension of each phase in the direction of heat ¯ow. The last term in the above equation can be determined using any variation in the thermal conductivity across the + +![](images/63e0f10acb8a7f60194e818fac9614394bdf2c581e43a5d40594086c56b3a079.jpg) +Fig. 8. Schematic illustration of a di€used interface formed between diamond and aluminum nitride. A graded or smooth interface is formed when there is no minimum in the interface. A sharp interface is formed when the interface volume is zero. + +![](images/bc88e819a5f7aedf73900b37ee1a6f20d1c81f73edc30ef369c87946654e1777.jpg) +Fig. 9. E€ective thermal conductivity of the composite structure illustrated in Fig. 8 shown as a function of the volume fraction of the interface. Interfacial thermal resistance for a sharp interface is considered to be zero. Volume fraction of the interface region plus the AlN phase is kept constant at 0.0588. + +interface. Thus for a graded interface, $1 / K _ { \mathrm { i } } { = } 1 /$ $( \alpha d ) \mathrm { l n } ( K _ { \mathrm { d } } / K _ { \mathrm { a } } )$ where $\alpha = ( K _ { \mathrm { d } } { - } K _ { \mathrm { a } } ) / d .$ On the other hand, the equivalent expression for a di€used interface takes the form + +$$ +1 / K _ {\mathrm {i}} = (1 / d \alpha_ {1}) \ln \left(K _ {\mathrm {d}} / K _ {\mathrm {i}}\right) + (1 / d \alpha_ {2}) \ln \left(K _ {\mathrm {i}} / K _ {\mathrm {a}}\right) ], \tag {2} +$$ + +where $\alpha _ { 1 } = ( K _ { \mathrm { d } } { - } K _ { \mathrm { a } } ) / d _ { 1 } , \alpha _ { 1 } = ( K _ { \mathrm { i } } { - } K _ { \mathrm { a } } ) / d _ { 2 } .$ d1 and d2 are the dimensions of the interfacial regions $( d = d _ { 1 } + d _ { 2 } )$ , with changes in thermal conductivity, as shown in Fig. 8. For a sharp interface, $V _ { \mathrm { i } } { = } 0$ but an interfacial thermal barrier resistance can be de®ned in the form $2 V _ { \mathrm { i } } / K _ { \mathrm { i } } { = } 2 V _ { \mathrm { a } } / d h _ { \mathrm { c } }$ . The e€ective thermal conductivity in the presence of a sharp interface is controlled by the non-dimensionless parameter, $K _ { \mathrm { a } } / d h _ { \mathrm { c } }$ . When $K _ { \mathrm { a } } / d h _ { \mathrm { c } }$ becomes in®nitely large $( h _ { \mathrm { c } } { = } 0 )$ , the e€ective thermal conductivity corresponds to the value for polycrystalline diamond with a dispersed void space. + +A detailed numerical analysis of the e€ective thermal conductivity of the diamond/AlN/diamond composite is presented in Ref.[14] for three di€erent situations, namely, when the interfacial layer is formed in the AlN region or in the diamond region or in both the AlN and diamond regions equally. The results of analysis when the di€used interface is formed in the AlN region are presented in Fig. 9. The e€ective thermal conductivity of the composite is seen to increase with increase in volume fraction of the interfacial region although a minimum is present in the thermal conductivity across the interface. The diamond layer of thickness 4 mm on either side surrounds the AlN region of 0.5 mm with the thickness of interface region increase and that of AlN decrease for larger volume fraction of the interface. The thermal conductivity of diamond was assumed to be 20 W/cm K and that of AlN at 3.7 W/cm K with the minimum value of the thermal conductivity across the interface chosen to be 1.0 W/cm K. It is seen that the composite with the graded interface has the highest thermal conductivity with that of di€used interface at the intermediate value and that + +with sharp interface lowest when the interfacial thermal resistance was zero. If the interfacial thermal resistance were to be very high as in the case of diamond ®lm with voids across the interface, the e€ective thermal conductivity of the composite with sharp interface will be very small. We have not assumed this situation since percolation paths are also available for heat ¯ow. These results shown in Fig. 9 illustrate that the e€ective thermal conductivity of diamond/AlN/diamond composites will be higher than that of polycrystalline synthetic diamond with diamondlike carbon or voids. + +# 5. CONCLUSIONS + +Diamond heat spreaders were successfully bonded to device wafers of either silicon or GaAs using gold±tin eutectic solder with proper metallization procedures. AlN embedding layer improves the adhesion strength of diamond to molybdenum or silicon nitride, thermal stability of the bond and the heat spreader characteristics. In the present study, we have shown that multilayer diamond ®lms exhibit higher e€ective thermal conductivity than either single layer diamond or molybdenum. Modeling and analysis of the interface region suggests that AlN ®lm replaces the diamondlike carbon phase and void space between large grains of diamond and thereby volume fraction of higher conductivity region is improved in multilayer diamond composites. Further, we have shown that the multilayer diamond heat spreaders are helpful to lower the temperature of the electronic devices and thereby prolong their life. + +AcknowledgementsÐThis research is supported by the Division of DMII, NSF. Also, this research is sponsored by SURA and the Assistant Secretary for Energy Eciency and Renewable Energy, Oce of Transportation Technologies, as part of the HTML User Program, ORNL, managed by Lockheed Martin Energy Research Corp. for the U.S. DOE under contract number DE-AC05-96OR22464. + +# REFERENCES + +1. Kuttel, O. 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b/金刚石散热项目/项目前期工作_MD/02_金属化金刚石载板与器件封装/6_Au-(Ti-W) and Au-Cr metallization of chemically vapor-deposited diamond substrates for multichip module applications/6_Au-(Ti-W) and Au-Cr metallization of chemically vapor-deposited diamond substrates for multichip module applications.md @@ -0,0 +1,142 @@ +# $\mathrm{Au} / (\mathrm{Ti}-\mathrm{W})$ and $\mathrm{Au} / \mathrm{Cr}$ metallization of chemically vapor-deposited diamond substrates for multichip module applications + +Ilango Meyyappan, A. P. Malshe, H. A. Naseem, W. D. Brown + +High Density Electronics Center, University of Arkansas, Fayetteville, AR 72701, USA + +# Abstract + +Since diamond obtained by chemical vapor deposition (CVD) has an extremely high thermal conductivity, it holds great promise in solving thermal management problems in high performance multichip modules (MCMs). Consequently, there is a need to develop a reliable metallization system for CVD diamond. Refractory metals such as Ti, Mo, Ta and W are known to form adhering carbide layers at high temperatures. Also, transition metals such as Cr, Ni and Ni-Cr are widely used in other MCM technologies involving Si, AlN, SiC and alumina substrates. In the work reported here, adherent $\mathrm{Au} / \mathrm{Cr}$ and $\mathrm{Au} / (\mathrm{Ti} - \mathrm{W})$ metallization systems were produced at low temperatures using d.c. magnetron sputtering and electron beam evaporation techniques. Adhesion at low temperature is essential since CVD diamond could lose its thermal and electrical properties at high temperatures. Furthermore, interaction between metal layers may cause an increase in conductor trace resistivity and delamination. Adhesion was measured using a Sebastian V-A thin film stud pull tester. The deposition parameters were optimized to give maximum adhesion using a statistical design software package, ECHIP. In the case of the sputtered metallization, pre-sputter cleaning of diamond surface improved adhesion significantly. Values above $9\mathrm{klbf}$ in $-^2$ were obtained in the case of $\mathrm{Au} / (\mathrm{Ti} - \mathrm{W})$ and $11.8\mathrm{klbf}$ in $-^2$ in the case of $\mathrm{Au} / \mathrm{Cr}$ . Post-deposition annealing was performed in nitrogen ambient to investigate the effect of post-metallization processing on adhesion and also to test for any possible interaction between the metals at high temperatures. Annealing temperatures were limited to $450^{\circ}\mathrm{C}$ since MCM substrates are seldom exposed to temperatures higher than these. Energy-dispersive spectroscopy (EDS) analysis indicated outdiffusion of W through Au at $400^{\circ}\mathrm{C}$ . No interdiffusion was observed in the case of $\mathrm{Au} / \mathrm{Cr}$ as per optical microscopy and EDS analysis. Auger electron spectroscopy results indicate interaction between the metals in both $\mathrm{Au} / (\mathrm{Ti} - \mathrm{W})$ and $\mathrm{Au} / \mathrm{Cr}$ metallizations at $450^{\circ}\mathrm{C}$ . + +Keywords: Chemical vapour deposition; Diamond; Gold; Metallization + +# 1. Introduction + +Multichip packaging technology has received widespread attention in the electronics industry. The use of multichip modules (MCMs) promises to increase packaging density and system performance beyond what is otherwise possible through the use of very-large-scale integration and surface mount technology [1]. As in the case of most advanced technologies, the drive towards high performance MCMs has given rise to a number of engineering challenges. In advanced MCMs where high speed, high-power chips are mounted shoulder to shoulder on large circuit boards, heat spreading and removal become major issues. Diamond has the highest known thermal conductivity and at the same time is an electrical insulator. It holds great promise in solving thermal management problems in MCMs. The use of diamond substrates may also make the use of three-dimensional MCMs possible. + +# 1.1. Diamond as a substrate + +Substrate materials are chosen on the basis of dielectric strength, dielectric constant, dissipation factor, thermal conductivity, thermal expansion and surface finish. Substrate material choices have essentially been limited to silicon, aluminum oxide and to a much lesser extent AlN, SiC and glass-ceramics. Today, AlN substrates and packages are serious contenders in many new thermal packaging designs. However, power dissipation in some next-generation equipment will require materials capable of removing heat more efficiently than AlN or even copper. + +Recently, there has been considerable progress in the preparation of diamond and diamond-like films and even of free-standing diamond substrates. High purity diamond exhibits thermal conductivities five times higher than copper and is being developed as a sub + +strate or heat sink for high power radar, microwave and super computer applications [2]. + +Diamond has an additional advantage. Its thermal expansion coefficient is similar to that of silicon and silicon dioxide. Other attractive properties of diamond as a substrate include good mechanical strength, low dielectric constant, excellent dielectric properties, non-toxicity and stability under working conditions. It can also withstand high thermal shock [3]. + +# 1.2. Metallization systems + +The desirable characteristics of a metallization system are good adhesion to the substrate, low stress, good electrical conductivity and minimal reactions at subsequent processing temperatures (up to $400^{\circ}\mathrm{C}$ [4]. Gold was chosen as the metal for the conductive layer for its superior electrical conductivity and ease of deposition. Gold protects the surface from environmental attack and helps to ensure bondability. It has good solderability and is an extremely inert metal which protects it from oxidation. However, as gold is a noble metal, is expected to have poor adhesion to diamond, which itself is chemically inert. + +Two classes of metals have been reported to form carbides with diamond at high temperatures [5]. These are refractory metals such as Ti, W and Mo and transition metals such as Cr, Ni and Fe. Roser [6] reported that refractory metals from carbide precipitates at the metal-diamond interface. These carbide precipitates were reported to result in stable, adhering, ohmic contact to diamond. Moazed et al. [7] and Venkatesan and Das [8] also reported carbide formation for $\mathrm{Mo / N_i}$ and Ti metallizations at $950^{\circ}\mathrm{C}$ and $850^{\circ}\mathrm{C}$ respectively. However, temperatures as low as $580^{\circ}\mathrm{C}$ were reported by Gildenblat et al. [9] for carbide formation with Au/Ti on diamond films obtained by chemical vapor deposition (CVD) [9]. Cr has been used extensively in many MCM technologies as adhesion promoter and diffusion barrier metal. In this study, we investigated $\mathrm{Au} / (\mathrm{Ti} - \mathrm{W})$ and $\mathrm{Au / Cr}$ metallizations for potential use as high density interconnects on diamond-based MCMs. + +Since it is expected that CVD diamond undergoes some change in its properties at elevated temperatures, and since interdiffusion of metals may cause increase in the resistivity of Au conductor traces and even delamination, it was proposed to study the adhesion properties of these metals at low deposition and annealing temperatures. In the case of sputtered metallization, improvement in adhesin was taught by activating the surface carbon atoms of diamond using a sputter-etching treatment. + +# 2. Experimental details + +A Perkin-Elmer Randex 2400 diode sputtering system was modified to do d.c. magnetron sputtering and + +a CVC electron beam evaporater was used to deposit films by the evaporation technique. In order to investigate several metallization systems to achieve adherent conducting films on diamond substrates (that would meet stringent quality control specifications for MCM applications), the effects of various deposition variables on adhesion were studied. Since sputtering and evaporation involve many control variables such as deposition rate, power, pressure, substrate etch, deposition time and sputter etch time, a one-variable-at-a-time approach to optimize adhesion becomes prohibitive. + +Thus a statistical design that minimizes the number of experiments to be performed was followed. In this approach, any interaction between the control variables is easily detected. Furthermore, regression models can be used to give an optimum set of control variables for the desired response values. Of the several software packages available, the response surface methodology supported by ECHIP was used in this study. + +The free-standing diamond substrates used for this work were supplied by Norton Diamond Films, Northboro, MA, and were grown using a magnetically enhanced d.c. arc jet technique. This system gives large-area (up to $15\mathrm{cm}$ diameter), high quality, freestanding substrates at very high deposition rates. The substrates used for this work had a surface roughness of less than $1\mu \mathrm{m}$ (after polishing) and a dielectric constant of 5.57. + +The diamond substrates were laser cut to either $9\mathrm{mm}\times 6\mathrm{mm}$ or $5\mathrm{mm}\times 5\mathrm{mm}$ in size. These samples were lapped-polished and plasma etched to remove graphite from their surfaces. They were cleaned in trichloroethylene, acetone and methanol solutions (in that order) in an ultrasonic cleaner just before metallization. In that case of sputtering, the films were deposited at a chamber pressure of $10\mathrm{mT}$ (base pressure, $(2-4)\times 10^{-6}\mathrm{T})$ in argon to a total thickness of $5000\AA$ $(2000\AA$ of Ti-W or Cr and $3000\AA$ of Au). These films were furnace annealed in a nitrogen ambient for $30\mathrm{min}$ . + +Testing for adhesion was performed using a Sebastian V-A thin film adhesion tester. Epoxy-coated aluminum studs, $2.7\mathrm{mm}$ in diameter, were attached to the samples by curing at $150^{\circ}\mathrm{C}$ for $1\mathrm{h}$ . Analytical techniques such as sheet resistance measurements, optical microscopy, energy-dispersive spectroscopy (EDS) and Auger electron spectroscopy (AES) were used to study the extent of interdiffusion between the films. + +# 3. Results and discussion + +# 3.1. $Au / (Ti - W)$ metallization system + +A tabulation of adhesion values obtained for the 17 experiments as suggested by ECHIP is shown in Table 1. The response surfaces plotted using ECHIP are shown in + +Table 1 Adhesion values of $\mathrm{Au / (Ti - W)}$ diamond samples + +
d.c. power (W)Sputter etch power (W)Sputter etch time (min)Adhesion (klbf in-2)
115015509.09
250101500.73
35015507.93
41505504.51
5150151509.09
6100151000.941
710010503.61
8150101008.54
9505504.71
1010051506.47
11100151509.09
1250151009.09
135051009.09
14150101503.21
155010504.23
115015508.39
250101501.029
+ +![](images/6366cf90166cb3e5dc81e592301f1b0c9b38bcee00b582332fe99c1fc558dca8.jpg) +Fig. 1. Response surfaces for $\mathrm{Au} / (\mathrm{Ti - W})$ /diamond samples. + +Fig. 1. As the table shows, adhesion values up to 9 klf in $^{-2}$ were obtained. The pull test was not performed at higher values because the substrates tended to break above this stress value. Experiments without a substrate sputter etch gave adhesion values around 0.2 klf in $^{-2}$ as opposed to values as high as 9 klf in $^{-2}$ for samples exposed to a sputter etch prior to metal deposition. + +The response surfaces indicate that a sputter etch power of $100\mathrm{W}$ gives maximum adhesion. A sputter etch power of $150\mathrm{W}$ also gives good adhesion if the etch time is 5 or $15\mathrm{min}$ . An etch time of $10\mathrm{min}$ gives minimum adhesion. A possible explanation for this is that etching for a short time activates the surface sufficiently for good adhesion, but continued etching may result in the formation of a graphitic layer that is not a good adhering layer. Further etching (15 min) + +may then remove the graphitic layer and expose the diamond surface for the deposition of metal layers. + +No change was observed in the adhesion values when the samples were annealed at 150 and $300^{\circ}\mathrm{C}$ . However, when they were annealed at $450^{\circ}\mathrm{C}$ , samples that initially gave an adhesion value of $9\mathrm{klbfin}^{-2}$ (good adhering films) and $1.8\mathrm{klbf}^{-2}$ (poorly adhering films) degraded to $5.75\mathrm{klbfin}^{-2}$ and $0.3\mathrm{klbfin}^{-2}$ respectively. The surface was observed under an optical microscope and was seen to turn slightly bluish green along the grain boundaries (Fig. 2(a)). The adhesion degradation can be attributed to the outdiffusion of Ti-W to the surface of gold where it oxidized owing to the presence of trace amounts of oxygen and/or moisture in the purge gas. The presence of oxygen in the annealing ambient may have exacerbated this outdiffusion owing to what is known as "ambient effect" [10]. Similar studies were performed using silicon substrates that also showed a significant number of defects appearing on the surface at $450^{\circ}\mathrm{C}$ (Fig. 2(b)). + +The energy-dispersive spectra of samples annealed at 150 and $300^{\circ}\mathrm{C}$ showed only the $\mathbf{M}\alpha$ characteristic peak of Au at $2.14\mathrm{keV}$ . The Ti K $\alpha$ and W M $\alpha$ peaks did not appear, but, for the sample annealed at $450^{\circ}\mathrm{C}$ , the W M $\alpha$ peak appeared at $1.79\mathrm{keV}$ . These results strongly support the explanation given for the discoloration of Au at $450^{\circ}\mathrm{C}$ . + +The compositional Auger depth profiles of the samples are shown in Fig. 3. No changes were observed for the 150 and $300^{\circ}\mathrm{C}$ sample, but, at $450^{\circ}\mathrm{C}$ , very interesting observations can be made. There are two regions where the W concentration exhibits a maximum. A scanning electron photomicrograph taken after the depth profile also confirmed this observation (see inset of Fig. 3). It can be seen that there is complete interdiffusion between Ti-W and Au. Ti and W outdiffuse through the Au layer first. About $60\mathrm{at.}\% \mathrm{W}$ and $15\mathrm{at.}\%$ Ti were present at this point. The Au concentration was only $25\mathrm{at.}\%$ in this region, but, Au outdiffused through Ti and W, resulting in $80\mathrm{at.}\%$ Au at the surface with only $10\mathrm{at.}\%$ W and $5\mathrm{at.}\%$ Ti. Clearly, this interdiffusion can only be attributed to grain boundary diffusion since the bulk solubility of either metal (Ti-W or Au) in the other is very small. The metal layers are therefore not stable at this temperature. It can also be seen that the slope of the C profile decreases considerably in the $450^{\circ}\mathrm{C}$ sample compared with the $300^{\circ}\mathrm{C}$ sample. This means that C atoms are slowly diffusing into the Ti and W regions, possibly forming TiC and WC. + +# 3.2. Sputtered $Au / Cr$ metallization system + +The $\mathrm{Au} / \mathrm{Cr}$ diamond samples deposited using the sputtering system gave adhesion values above 3 klfb in $^{-2}$ . Adhesion testing was stopped at this point be + +![](images/4e8e6a44fe3b5c931f5210b9a93b976f173db8cb99d5071bb9a31d02555320a6.jpg) + +![](images/5dc740520102336ce57e8a4ee6c004bdeb8a5bcb7861c6f11327bd7fd75c6940.jpg) +Fig. 2. Optical micrographs of (a) $\mathrm{Au} / (\mathrm{Ti} - \mathrm{W})/$ diamond samples annealed at 300 and $450^{\circ}\mathrm{C}$ and (b) $\mathrm{Au} / (\mathrm{Ti} - \mathrm{W})/$ silicon samples annealed at 300 and $450^{\circ}\mathrm{C}$ . + +![](images/46cdecb2bd08031d65c2a69fb324dd7e42da11294522cc071f55074dcdf565ee.jpg) +Fig. 3. Compositional Auger depth profiles of $\mathrm{Au} / (\mathrm{Ti - W}) / \mathrm{diamond}$ samples annealed at 150, 300 and $450^{\circ}\mathrm{C}$ . + +cause the diamond substrates were breaking beyond this value. Annealing studies on these samples showed some change in the surface of the gold at $450^{\circ}\mathrm{C}$ when observed under the optical microscope. The films deposited on silicon also exhibited a change in the color of the surface of the gold layer following annealing at this temperature (Fig. 4). The energy-dispersive spectra showed only the Au $\mathbf{M}\alpha$ characteristic peak and no Cr peak for all the samples. Therefore, although there is + +![](images/6e8181dcb13a49c7bc310c4395b66535eee82994f9cf1bb8dfd31a8c33680d8a.jpg) +Fig. 4. Optical micrographs of sputtered $\mathrm{Au / Cr / }$ diamond samples annealed at 300 and $450^{\circ}\mathrm{C}$ . + +some interdiffusion between Cr and Au at $450^{\circ}C$ , it is not extensive enough when compared with the $\mathrm{Au / (Ti - }$ W)/diamond samples. + +The compositional Auger depth profiles of the samples annealed at 300 and $450^{\circ}\mathrm{C}$ are shown in Fig. 5. It can be seen that the Au concentration remains at 85 at.% for both the samples, but the Cr concentration in the Cr region decreased from 65 at.% in the $300^{\circ}\mathrm{C}$ sample to 45 at.% in the $450^{\circ}\mathrm{C}$ sample. Also, the Au concentration increased from 23 at.% in the $300^{\circ}\mathrm{C}$ sample to 45 at.% in the $450^{\circ}\mathrm{C}$ sample. Therefore Au has diffused into the Cr layer. The Cr concentration in the surface increased from 5 at.% in the $300^{\circ}\mathrm{C}$ sample + +![](images/b0f4a951f8d032d18b876691265ff359f3aa83212c03feaea8ff7fcaa9813ee1.jpg) + +![](images/71f9640b889886d6288e84ddde15a281f68338fddf03d0d890899f078b7a3d25.jpg) +Fig. 5. Compositional Auger depth profiles of sputtered $\mathrm{Au} / \mathrm{Cr} / \mathrm{diam}$ diamond samples annealed at (a) $300^{\circ}\mathrm{C}$ and (b) $450^{\circ}\mathrm{C}$ . + +to $19\mathrm{at.}\%$ in the $450^{\circ}\mathrm{C}$ sample. The $\mathrm{Cr}$ concentration in the Au layer also increased from 2 at. $\%$ in the $300^{\circ}\mathrm{C}$ to 5 at. $\%$ in the $450^{\circ}\mathrm{C}$ sample. This may be due to the increased amount of Cr in the Au grain boundaries or higher solubility in the bulk Au at higher temperatures. Also, the slope of the C profile decreased considerably in the $450^{\circ}\mathrm{C}$ sample when compared with the $300^{\circ}\mathrm{C}$ sample. This means that C and Cr interdiffuse at this temperature, possibly forming CrC. + +# 3.3. Evaporated $Au / Cr$ metallization system + +Adhesion values obtained for the eight experiments as suggested by EChip are shown in Table 2. The deposition rate and total film thickness were varied. Adhesion values up to $11.8\mathrm{klbfin}^{-2}$ were obtained. The response surface obtained by plotting this data is + +Table 2 Adhesion values of evaporated $\mathrm{Au / Cr}$ /diamond samples + +
Deposition rate (Å s-1)Thickness (Å)Adhesion (klbf in-2)
550004.96
5300011.82
1050009.45
1550009.09
15100010.28
510009.85
550005.85
5300011.17
+ +![](images/d3d171499d33bf1fdf4bcd5869fda9653956f94cd4b96a2d4cbf2c22a4b572cf.jpg) +Fig. 6. Response surface of evaporated $\mathrm{Au / Cr}$ /diamond samples. + +![](images/3b53b5230c311b97dcccd8ab1c063b293418fc40d031cb93298a0d7081f9eb72.jpg) + +![](images/31e8935899ceb229ec4c2d350cb3179b4d6c353ebfad8318c50a952332930477.jpg) +Fig. 7. Optical micrographs of evaporated $\mathrm{Au / Cr / }$ diamond samples annealed at 300 and $450^{\circ}\mathrm{C}$ . + +shown in Fig. 6. As can be seen from the figure, maximum adhesion is obtained for a film thickness around $3000\mathring{\mathrm{A}}$ deposited at a rate of $10\mathring{\mathrm{A}}$ s. This corresponds to $2000\mathring{\mathrm{A}}$ of Au and $1000\mathring{\mathrm{A}}$ of Cr. + +Annealing studies did not reveal any change in the surface of the Au layer even for the sample annealed at $450^{\circ}\mathrm{C}$ . Similar studies on silicon also did not indicate any change in the surface morphology of the Au layer at this temperature (Fig. 7). This was surprising since sputtered Cr had shown some degradation at $450^{\circ}\mathrm{C}$ on diamond as well as Si substrates. This could be due to partial oxidation of evaporated Cr which was deposited at rates of about $10\AA \mathrm{s}^{-1}$ in a vacuum of $1\mu \mathrm{Torr}$ (base pressure). As in the case of the sputtered $\mathrm{Au / Cr}$ system, energy-dispersive spectra did not reveal any Cr peak even for the $450^{\circ}\mathrm{C}$ annealed sample. + +![](images/469b85d938f46766f349ed38398a8061fb40725ffe5f27e83403638f68632878.jpg) +Fig. 8. Compositional Auger depth profile of evaporated $\mathrm{Au / Cr / }$ diamond sample annealed at $450^{\circ}\mathrm{C}$ . + +The compositional Auger depth profile of $450^{\circ}\mathrm{C}$ annealed sample is shown in Fig. 8. There was no detectable interdiffusion between Au and Cr at $300^{\circ}\mathrm{C}$ . As observed for the sputtered sample, the slope of the C profile at $450^{\circ}\mathrm{C}$ indicates dislodging of C atoms into the Cr region. There is only about 10 at. $\%$ Cr on the surface of the sample, which explains why both optical microscopy and EDS techniques did not reveal outdiffusion of Cr to the surface. Since the concentrations of Cr and Au are about the same in the Cr region, Cr could be dissolving into the Au and, as a result, very little Cr outdiffuses to the surface. It was observed that below $400^{\circ}\mathrm{C}$ there is minimal interdiffusion between Au and Cr whereas, at $450^{\circ}\mathrm{C}$ , the interdiffusion occurs but is much less than that observed for either sputtered $\mathrm{Au / Cr}$ or $\mathrm{Au / (Ti - W)}$ metallization schemes. + +# 4. Conclusions + +Adhesion of sputtered $\mathrm{Au} / (\mathrm{Ti}-\mathrm{W})$ to diamond was found to improve dramatically when the surface was activated by the sputter etch technique. Extensive interdiffusion between Au and Ti-W was found to occur at $450^{\circ}\mathrm{C}$ as revealed by optical microscopy, EDS and AES studies. Interdiffusion between Cr and Au, although it occurs at these temperatures, was found to be more subdued compared with $\mathrm{Au} / (\mathrm{Ti}-\mathrm{W})$ . Thus Cr offers better adhesion-promoting and diffusion barrier properties (when compared with Ti-W) for Au-based conductor metallizations on diamond substrates in applications such as MCMs where processing temperatures are limited to $350^{\circ}\mathrm{C}$ . + +# References + +[1] K. M. Kearney, Semicond. Int., 12(2) (February 1989) 44. +[2] J. K. Hagge, IEEE Trans. Compon., Hybrids Manuf. Technol., 15(1) (1992) 29-39. +[3] H. A. Naseem, I. Meyyappan, C. S. Prasad and W. D. Brown, Int. J. Microcircuits Electron. Packag. 16(4) (1993) 257-269. +[4] F. A. Clay, N. T. Panaousis and R. W. Pierce, Jr., IEEE Trans. Compon., Hybrids Manuf. Technol., 10(4) (1974) 258-262. +[5] De Beers Diamond Research Laboratory, Intern. Doc., No. 3000/7/85/KP, 1985. +[6] M. Roser, Proc. Electrochemical Soc., 91-8 (1991) 551. +[7] K. L. Moazed, J. R. Zeidler and M. J. Taylor, J. Appl. Phys., 68(5) (1990) 2246-2254. +[8] V. Venkatesan and K. Das, IEEE Electron Devices Lett., 13(2) (1992) 126-128. +[9] G. S. Gildenblat, S. A. Grot and A. Badzian, Proc. IEEE, 79(5) (1991) 647-667. +[10] C. A. 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b/金刚石散热项目/项目前期工作_MD/02_金属化金刚石载板与器件封装/7_Advanced metallization schemes for bonding of InP-based laser devices to CVD-diamond heatsinks/7_Advanced metallization schemes for bonding of InP-based laser devices to CVD-diamond heatsinks.md @@ -0,0 +1,570 @@ +# Advanced metallization schemes for bonding of InP-based laser devices to CVD-diamond heatsinks + +A. Katz, C.H. Lee and K.L. Tai + +AT&T Bell Laboratories, Murray Hill, NJ 07974 (USA) + +(Received June 17, 1993) + +# Abstract + +Semiconductor devices, and in particular InP-based laser devices, are usually bonded on a mounting plate (called a submount or heatsink) or directly onto a package. This bonding assembly, which comprises the die-bonding metallic layers, the joint solder and the submount, serves the purpose of heat dissipator, mechanical support and electrical conductor. As such, the quality of both the bonding metallization and the submount are as important as the device die itself to assure the short- and long-term reliable operation of the electronic assembly. Due to its high thermal conductivity, diamond has been used as a very efficient thermal conductor and dissipator under electronic and photonic devices. This application has become even more attractive since the commercialization of the chemically vapor-deposited (CVD)-diamond, due to its lower cost and larger available surface area compared to natural diamond. Therefore it is only natural to use CVD-diamond as the material of choice for mass production of high-power InP-based lascr diode submounts. The device is attached to the submount by a metallic bonding medium that contains a few layers, such as an adhesion layer (typically of titanium (Ti) adjacent to the submount), a barrier layer (typically of platinum (Pt)) and capped with hard solder (such as gold-tin (AuSn) eutectic alloy, which has a melting temperature of $278^{\circ}\mathrm{C}$ ). While offering optimum bonding conditions, the Ti/Pt/AuSn bonding system provides a high quality bond of the laser diode to the CVD-diamond submount. However, the extensive reaction of the AuSn solder with the Pt and Ti layers, even after short bonding periods of 5–10 s, may lead to mechanical deterioration of the bonded assembly, resulting in delamination of the metal and failure of the bond. This failure occurs mainly due to the thermodynamically reactive nature of the Pt-Sn couple, which reacts almost spontaneously even at room temperature. The reaction consumes the Pt layer and an appreciable amount of Sn, leading to the disappearance of the barrier layer and to dilution of Sn from the solder. The variation in the molten solder stoichiometry results in premature solidification of the solder through the bonding cycle after consumption of the entire Pt barrier layer, and dissolution of the Ti adhesion layer, as well. In order to maintain the good wetting performance of the AuSn solder to the Ti/Pt metals, but yet to improve the thermodynamic stability of the metallurgical bonding system, various other metals such as Ni, W, Cr and some of their alloys were evaluated as advanced alternatives to the Pt barrier layer. The quality of the evaluated metallurgical systems was judged upon wetting performance, thermodynamic stability and lack of premature freezing of the molten solder. The Ti/W/W( $\mathrm{Ni}_3\mathrm{Sn}_4$ )/ $\mathrm{Ni}_3\mathrm{Sn}_4/\mathrm{Au}$ multilayered structure was finally defined as the optimal and superior metallurgical scheme for the purpose of bonding laser chip to CVD-diamond submount. The time required until the first surface local freezing phenomenon was observed at the AuSn solder (with eutectic composition of 80 wt.% Au) on Ti/W/W( $\mathrm{Ni}_3\mathrm{Sn}_4$ )/ $\mathrm{Ni}_3\mathrm{Sn}_4/\mathrm{Au}$ structure while melted at $320^{\circ}\mathrm{C}$ was 200 s, compared to 5–10 s at the Ti/Pt/Au/AuSn system. At the former system the solder was maintained melted at $320^{\circ}\mathrm{C}$ for more than 1 h before it was completely frozen, which is much longer than the 30 s measured for the latter system. It was also observed that not only did the AuSn solder on the Ti/W/W( $\mathrm{Ni}_3\mathrm{Sn}_4$ )/ $\mathrm{Ni}_3\mathrm{Sn}_4$ multilayer maintain its accurate stoichiometric composition through the entire bonding cycle, but it also provided excellent adhesion and integrity for the entire bonded assembly. + +# Introduction + +An InP laser diode is typically bonded to a submount that provides mechanical stability, heat dissipation and + +electrical connections to one or more electrical sources. A reliable device bonding is essential to provide short- and long-term operating reliability. The quality of the bonding medium depends on the appropriate selection + +of the submount material, the solder metallurgical system, the barrier metallization scheme and the bonding conditions. + +Submounts for laser devices have traditionally been manufactured from Si or ceramic materials such as BeO. Recently, due to its desirable high thermal conductivity and high electrical resistivity, diamond has been revealed as an attractive alternative for the submount material of choice, and in particular for operation of devices under extreme conditions [1, 2]. Thus, a natural application of diamond is likely to be one that takes advantage of its potentially high thermal conductivity, which can be as high as $25\mathrm{W}(\mathrm{cm}^{\circ}\mathrm{C})^{-1}$ (at room temperature). This is about five times better than any other excellent conductors such as copper [3, 4]. Because of its high thermal conductivity, natural diamond has been used for many years as a 'heat sink' or thermally conducting substrate for dissipating heat from electronic and photonic devices. The demand for such substrates formed an obvious market for new chemically vapor-deposited (CVD) diamond material, which potentially has the advantages of lower price and larger available processing surface areas. This potential as a passive electronic component, has been featured in all the reviews that assessed the future of CVD diamond. American Telephone & Telegraph (AT&T) Co. has recently announced [5, 6] the first commercial application of CVD diamond as submounts for InP-based laser devices, and it is expected that this concept will be applied in a wide range of electronic and photonic devices. + +Traditionally, solder preforms, paste or deposited solder have been used to bond microelectronic and photonic devices onto submounts. The deposited solder, however, provides significant advantages over other techniques, such as decreased oxide formation prior to the initiation of the bonding cycle, and more accurate control of the bonding media thickness and volume [7]. The few commonly-used solders are the lower-melting elements such as In $(183^{\circ}\mathrm{C})$ , medium-melting systems, such as binary Pb-Sn alloy $(223^{\circ}\mathrm{C})$ , and higher-melting alloys, such as AuSn $(278^{\circ}\mathrm{C})$ , AuGe $(361^{\circ}\mathrm{C})$ or AuSi $(364^{\circ}\mathrm{C})$ . Generally, the lower-melting solders exhibit better thermal conductivity but are mechanically weaker, compared to the higher-melting solders. Thus, for better thermal stability and improved long-term reliability of performance, use of the higher-melting solders is preferred [8-11]. Au-Sn solder has been widely used in optoelectronic applications due to its high, but practical, melting temperature, which allows for efficient bonding of devices that also contain temperature-sensitive elements such as Au-based contacts. + +The Au-Sn phase diagram is a eutectic binary phase diagram (as shown in Fig. 1(a)), with two eutectic melting points, at 80 wt.% Au $(278^{\circ}\mathrm{C})$ and at 90 wt.% + +Sn $(232^{\circ}\mathrm{C})$ . The former composition was used for the above-discussed soldering applications and is defined by very steep liquidus lines on both sides of the eutectic melting point. This means that, for example, enriching the eutectic composition by $1\mathrm{wt}.\%$ of Au leads to an approximately $30^{\circ}\mathrm{C}$ increase in the melting temperature. Thus, it suggests that the AuSn solder system requires very accurate process control of both the solder composition and the bonding temperature to permit successful bonding. + +The third important factor involved in the bonding assembly includes the adhesion-barrier metals that allow for efficient introduction of the AuSn solder onto the submount. The commonly-used adhesion-barrier metals are the Ti/Pt metals, which have been widely studied and reported [10, 12]. + +It was reported that introducing AuSn solder onto the Ti/Pt adhesion-barrier metal structure leads to a rapid and extensive metallurgical interaction between the Sn and Pt to form certain intermetallics, such as PtSn [13]. The binary phase diagrams for the elements comprised in the Ti/Pt/AuSn metallurgical scheme are provided in Fig. 1, and further discussion on the thermodynamics of these systems is available in the literature [7, 10, 14, 15]. The reactivity of these systems results in intensive reactions that take place at the Ti/Pt/AuSn interface. The most disturbing one is the consumption of the Sn from the solder overlayer, which leads to the formation of an Au-rich solder layer, and results in elevation of the solder melting temperature and its premature freezing through the bonding cycle. + +Thus, the complex nature of the AuSn solder system combined with the fact that SnPt is such a reactive system, suggest that the overall Ti/Pt/AuSn bonding metallurgical system is relatively unstable, and although being used, only a narrow bonding processing window is available. As a result, studies have been initiated to explore some potential metal alternatives for the solder and for the barrier metal systems, in order to achieve a more robust and stable bonding process. The initial, and simplest approach is to replace the reactive Pt barrier layer with a more stable metal, as is shown schematically in Fig. 2. + +Ni was found to be one of the more attractive candidates to replace the Pt barrier metal, for the following reasons. The Au-Ni phase diagram [16] (see Fig. 3(a)) suggests that the solubility of Ni in Au at temperatures below $350^{\circ}\mathrm{C}$ is very limited. Au and Ni do not dissolve in each other at temperatures lower than $100^{\circ}\mathrm{C}$ , and in addition Ni, as opposed to Pt, dissolves in Sn-based solders at very slow rates [17, 18]. However, some Ni-based intermetallics such as $\mathrm{Ni}_3\mathrm{Sn}_4$ and $\mathrm{Ni}_3\mathrm{Sn}_2$ [19-22] may still be formed during the bonding cycles at temperatures higher than $250^{\circ}\mathrm{C}$ . Both intermetallics are stable and do not consume much + +![](images/87ed39ec531c03cd8aa0e18f317ed5dc664927e73dd0c5b2078c6ee5b57443f2.jpg) + +![](images/a452b1e902bc59a59c703c67d51ea8361b29871423e72992ffba2cc76b38043c.jpg) + +![](images/43e6cf3d3bc4660a78316db5dfabf4a41d4dfa079b85b8d7298866237f719bb8.jpg) + +![](images/61a38737dd750b6bdc7b560fd5214edf89b9db8fd85103c65dfe8a4143abcaa8.jpg) +Fig. 1. Binary phase diagrams of the (a) AuSn, (b) TiPt, (c) AuPt, (d) PtSn systems. + +![](images/d2e6cfcebd4cc67b75757993c6e0643454ccd47efadc143e1fbf3b77db10e181.jpg) +Fig. 2. Schematic presentation of conceptual design of simple bonded device assembly. + +Sn (see Fig. 3(b)). In addition, pure Ni possesses superior thermal and electrical conductivities [23] compared to other potential barrier metals, such as Co or Pt. For + +these reasons Ni has, indeed, already been suggested as a barrier metal under the PbSn (95% Pb) solder [24]. + +Other metals that seem attractive as a replacement to the Pt barrier layer, are W and Cr. The binary phase diagrams of Au-W, Au-Cr and Cr-Sn (shown in Figs. 4(a), 4(b) and 4(c)) are given for reference, to be compared to the earlier-discussed Pt-Au and Pt-Sn systems, demonstrating the superior thermodynamic stability of the W- and Cr-containing systems. Unfortunately, no phase diagram is available for the Sn-W system, but the solubility of W in liquid Sn was reported to be very low ( $\sim 0.001$ at.%) even at $2000^{\circ}\mathrm{C}$ , and no intermetallics were found in this system. + +By replacing the Pt barrier metal layer with any of the above-mentioned alternative metals, better thermal stability of the bonding metallurgical system was, indeed, achieved. Cr was found to be less reactive than Pt. + +![](images/2e89e381d72d32f95bc29fca3ebf1f8d87a9c7a7e57032921c0f91ac2ab04fa0.jpg) +(a) +WEIGHT PERCENT NICKEL + +![](images/cad7925b047ff18b5c113dba6aa3c4f8da0b2649ee529232371d23b393c6ffe4.jpg) +(b) +WEIGHT PERCENT TIN +Fig. 3. Binary phase diagrams of (a) AuNi, (b) NiSn. + +When heated in contact with AuSn, Ni exhibited limited mutual solubility with the Au-based solder, and W provided a completely inert barrier under AuSn solder. The latter guarantees excellent thermodynamic stability of the system (see Fig. 2), but results in very poor adhesion between the AuSn solder and the W layer. Generally, the more inert the metal/metal interfaces, the better the thermal stability, which allowed for a wider bonding processing window and eliminated the premature freezing of the solder as a result of deviation from the eutectic stoichiometry composition. However, the most stable interfaces did not allow the formation of intermetallics, and thus resulted in poor adhesion of the metals as a result of lack of chemical bonding. + +In an attempt to improve, for example, the W/AuSn interface integrity, a thin Ti layer (10 nm) was introduced between the W layer and the AuSn solder [25] (see Fig. 2). The adhesion in this case was improved, but the inert nature of the W/AuSn interface was degraded, and again led to premature freezing of local solder spots on heating at $350^{\circ}\mathrm{C}$ , even after 15 s. + +Thus it was concluded that adhesion of AuSn solder to W in the Ti/W/AuSn system can be further improved without causing wider interfacial reactions, simply by introducing a better-designed wetting layer between the W and AuSn layers. Thus a wide range of Ti/ + +![](images/2e5f0d08b8fccfc445054d75c17a38a9504d49d292f28665aba679460192b301.jpg) +(a) +ATOMIC PERCENTUNGSTEN + +![](images/bf7a42292d2122aefe6ec0eb07960c11b251cacded0070f291773abf80912864.jpg) +(b) + +![](images/cfa21850f3a8629920ef3b41e604328e98d1cddbc440c6db4051dd05085c3c11.jpg) +ATOMIC PERCENTIN +Fig. 4. Binary phase diagrams of (a) AuW, (b) AuCr, (c) CrSn. + +$\mathrm{W}_x\mathrm{M}_y / \mathrm{AuSn}$ systems was studied, in which M represents a metal that can potentially be alloyed with W to provide a sufficient wetting layer. In particular Ti/ $\mathrm{W}_z(\mathrm{Au}_x\mathrm{Sn}_y) / \mathrm{AuSn}$ and $\mathrm{Ti} / \mathrm{W}_z(\mathrm{Ni}_x\mathrm{Sn}_y) / \mathrm{AuSn}$ systems were examined. The adhesion of the solder to the barrier metal improved, and the most stable performance was observed with a W/W $(\mathrm{Ni}_3\mathrm{Sn}_4) / \mathrm{Ni}_3\mathrm{Sn}_4$ multilayer structure under the AuSn solder. Initial solder local freezing on heating at $320^{\circ}\mathrm{C}$ required more than 3 min, and more than 1 h at this temperature was needed to completely freeze the entire $3\mu \mathrm{m}$ AuSn layer. In addition, an excellent solder-to-metal interfacial integrity was observed during and after the bonding cycle. + +This paper reviews the procedure for developing a superior bonding metallization scheme which allows stable and reliable bonding of laser diodes to any type of supporting submount, and in particular to a CVD-diamond submount, using a high-melting AuSn solder. + +# Experimental + +The diamond films used in this study were nucleated and grown on metal substrates by microwave-enhanced plasma chemical vapor deposition at $1.45\mathrm{GHz}$ . Depositions were carried out using a gas mixture of hydrogen, hydrocarbon and oxygen-bearing precursor, which resulted in an optically transparent CVD film. Except for some modifications in growth times, which led to changes in the CVD-diamond plate thickness, the deposition conditions were held constant from sample to sample. The self-supporting CVD-diamond plate was then chemically removed from the substrate. The plate surfaces were subsequently polished to mirror finish, metallized and laser-diced to heat-sink dimensions of about $0.65\times 1.25\mathrm{mm}$ with thickness of about $0.25\mathrm{mm}$ . These CVD-diamond films had a thermal conductivity in the range of $11 - 25\mathrm{W}(\mathrm{cm}^{\circ}\mathrm{C})^{-1}$ . + +An adhesion layer of $100\mathrm{nm}$ Ti was DC-magnetron sputtered onto the mirror-finished CVD-diamond plates, followed by deposition of $200 - 500\mathrm{nm}$ of Pt, Ni, W or W-based alloy. AuSn solder was electron-beam evaporated onto these base metal layers in a multilayer structure to provide an AuSn stoichiometric eutectic (Au 80 wt.%) composition $(T_{\mathrm{m}} = 278^{\circ}\mathrm{C})$ [26], at a total thickness of about $2.5 - 3\mu \mathrm{m}$ . The structure that was found to give the best results included an Au layer $(0.5\mu \mathrm{m})$ deposited onto the barrier layer, topped with four pairs of alternating $\mathrm{Sn}(0.35\mu \mathrm{m}) / \mathrm{Au}(0.2\mathrm{mm})$ layers, and covered with a final layer of $\mathrm{Au}(0.5\mu \mathrm{m})$ . + +InP-based laser devices were then bonded to the metallized CVD-diamond submounts during thermal cycles of $320 - 350^{\circ}\mathrm{C}$ for 5-10 s. The strength of the bond was assessed by shearing the device from the submount and analyzing the exact ruptured interface. + +In other samples the bonding cycles were simulated by carrying out reflow heating cycles at conditions similar to the bonding cycles. The metallurgical systems were reflowed by introducing the sample onto an electrical heating plate in a covered quartz vacuum chamber, in order to allow for visual tracking of the melting and freezing phenomena of the metals. A stereomicroscope with $30\times$ magnification was used to visually track any solder surface changes of the samples during the reflow process. The reflow process was carried out as follows: After evacuating the chamber to 5 mtorr for $2\mathrm{min}$ forming gas $(85\% \mathrm{N}_2,15\% \mathrm{H}_2)$ was flowed into it at a constant flow rate. After complete purge of the chamber (about 2 min of gas flow), the heating plate was turned on, and the metallized samples were heated to temperatures in the range of $300 - 350^{\circ}\mathrm{C}$ at a rate of $4.5^{\circ}\mathrm{C}s^{-1}$ as shown in Fig. 5. When the sample temperature reached $100^{\circ}\mathrm{C}$ timing was begun $(t = 0$ s). Once an initial surface local freezing phenomena was observed, the elapsed time was recorded and defined + +![](images/1844f8f9d0dd52af5cf8d66c377b37f681ddd59e8f3290334a80956b2218f0e3.jpg) +Fig. 5. Typical temperature-time profile of a reflow cycle under forming gas. + +as the 'surface local freezing time' (SLFT). After six more minutes of reflow, the power to the heating plate was turned off and the sample was cooled to room temperature. Some samples were maintained at the soaking temperature till the entire solder solidified. This time was also recorded, and defined as the solder 'global freezing time' (GFT). + +In addition to the reflow under forming gas, the samples were treated under flux (PEG400-5% abietic acid) in a simple bath furnace in order to evaluate the relative wettability of the different barrier metals to the AuSn solder in the temperature range $320 - 350^{\circ}\mathrm{C}$ , for durations up to 5 min. + +Composition and microstructural analysis of the different metallurgical samples were performed by a variety of means. Auger electron spectroscopy (AES) was used to obtain information on surface impurities (such as tin oxides) and, in conjunction with sputter depth profiling, to study the reaction of the adhesion and barrier metal layers with the AuSn solder through various reflow and bonding cycles. Rutherford backscattering spectrometry (RBS) was used to measure the solder composition and thickness in the as-deposited and annealed layer structures. Scanning electron microscopy (SEM) was used to obtain microstructural information, both of the solder pad surface morphology (in plan view), and of the quality of the layers and interfaces, obtained by cleaving the sample under the solder pad and examining the cross-section. Optical microscopy provided a quick way to screen surface morphology. Differential scanning calorimetry (DSC) experiments were conducted using a Perkin-Elmer DSC-4 apparatus. Samples of $2 - 10\mathrm{mg}$ were analyzed in aluminum pans previously coated with carbon. The DSC samples and chamber were flushed with ultrahigh purity nitrogen gas, and subsequently the samples were heated at a rate of $10^{\circ}\mathrm{C}\min^{-1}$ . + +Finally electro-optical analysis of bonded laser diodes was carried out to ascertain the final laser performance and the influence of the various metallurgical schemes on its performance. + +# Results and discussion + +# Ti/Pt/AuSn bonding scheme + +Optimized AuSn solder structure + +Currently, the AuSn soldering system is the most commonly used for high temperature bonding of high-power electronic and photonic devices. Bonding of the InP-based laser diode to the CVD-diamond submount is an excellent example of the use of this system. + +It was concluded that sufficient bonding of the laser device to the diamond submount is achieved with a eutectic AuSn layer (Au 80 wt.%) with total thickness of $3\mu \mathrm{m}$ or less, in order to guarantee a sufficient thermal conductivity through the solder joint, and eliminate technical problems associated with the presence of excess solder on the device. Given the unstable nature of the Pt/AuSn metallurgical system, a number of AuSn solder multilayer structures with the required stoichiometric average composition and thickness were investigated in order to determine the system that gave the best bonding results [27]. + +Bonding is carried out through a rapid thermal cycle, in which the temperature is quickly ramped up to $320 - 350^{\circ}\mathrm{C}$ and turned off in a period of 5 s. Clearly, differences in solder layer structure will affect the kinetics of the melting and bonding process, so that over the short time of the bonding process, different layer structures will lead to different results. For longer melting times (greater than 15 s), thermodynamic considerations become important. Then, no matter what the initial as-deposited solder structure may have been, the behavior of the melt over time will depend on the initial Au, Sn and Pt average ratio, since at this point, being the main interacting species, they have probably reacted completely. Presumably, the Pt barrier layer will consume an appreciable amount of Sn from the AuSn solder to form PtSn species. This has the detrimental effect of destroying the Pt barrier layer and altering the AuSn composition so that the melting behavior is adversely affected. + +The optimal metallization structure should allow a significant portion of the Pt barrier to remain intact during the bonding step, with the AuSn solder remaining molten long enough to allow for the bonding of the laser device. Note that the optimal metallization structure depends critically on the time scale of the bonding process. + +The effect of the AuSn solder layer structure on bonding properties is demonstrated through the evaluation of two samples. In the first sample (referred to hereafter as Sample I), a simple three-layered $\mathrm{Au}(0.8\mu \mathrm{m}) / \mathrm{Sn}(1.14\mu \mathrm{m}) / \mathrm{Au}(0.8\mu \mathrm{m})$ structure was deposited on top of the Pt/Ti barrier metals. This represents a simple processing sequence, with only two metal interfaces within the solder. The second sample + +(referred to hereafter as sample II) contains nine Au/ Sn alternating layers, produced through a relatively complicated processing sequence. Both samples were prepared with the same overall thickness and Au:Sn ratio. In both cases $0.5\mu \mathrm{m}$ of gold was deposited directly on top of the Pt layer in an attempt to separate the Sn and the Pt, and thus minimize the Sn-Pt interreaction. Each as-deposited sample was cut into several pieces to be processed at different temperatures. Two annealing conditions, $320^{\circ}\mathrm{C}$ for 5 s and $320^{\circ}\mathrm{C}$ for $10~s$ were applied, and the results were analyzed and compared to the as-deposited sample properties. + +RBS data of the two samples were used to determine the solder composition and thickness. The data for the as-deposited sample and for the sample annealed at $320^{\circ}\mathrm{C}$ for $10\mathrm{s}$ are summarized in Fig. 6. Interpretation of the as-deposited samples data shows that the overall thickness and Au:Sn composition of samples I and II were quite similar. The total thickness in both cases was about $2.7\mu \mathrm{m}$ , and the average Au:Sn ratio was close to $80:20$ wt.% Differences in the Au:Sn ratio with depth were observed between the two as-deposited samples, which was expected due to the different methods of preparation. For both as-deposited samples, the RBS data shows that no layers of pure Sn existed even after the deposition process. The general shape of the spectra confirms the existence of a predominantly pure gold layer at the bottom of the solder next to the Pt, and another Au layer on top of the solder. In the middle, the Sn intermixed readily with the gold to form an approximately 50:50 at.% alloy. The middle layer composition varies between samples I and II. In general, the picture that emerges is that during deposition, the first gold layer goes down on top of Pt and forms a sharp interface. The next layer deposited, which is Sn, is known to form a fairly porous layer. When the next Au layer is deposited onto it, it mixes readily with the Sn to form an AuSn alloy. The process continues in this manner until the final gold deposition, which is a fairly thick layer. Some portion of that may react with the previously deposited Sn layer, leaving a portion of mainly pure Au on top of the structure. + +After annealing, the RBS spectra from samples I and II are qualitatively different. Note that the analysis beam spot size used here is $\sim 250\mu \mathrm{m}$ in diameter, which covers most of the solder pad. Any lateral variations in thickness and composition will be averaged in the final spectrum. The data from sample II (Fig. $6(b_{2}))$ show a final solder composition of 80:20 wt.% and a fairly uniform thickness of solder over the pad. In contrast, the data from sample I (Fig. $6(a_{2}))$ show a similar solder composition but a greatly smeared low energy edge. This indicates very nonuniform solder thickness across the pad, and hence a nonuniform melting behavior across different areas of the pad. + +![](images/5367646517d6bc09fd7c9084721f87ecaccc1551ff171a0eb046d9ad804d7292.jpg) +Fig. 6. RBS spectra of metallized submounts; $(\mathbf{a}_1)$ and $(\mathbf{a}_2)$ are as-deposited and annealed spectra from sample I (simple 3-layer structure); $(\mathbf{b}_1)$ and $(\mathbf{b}_2)$ are as-deposited and annealed spectra from sample II (9-layer structure). See text for deposition details. + +Figure 7 gives SEM cross-sectional micrographs from as-deposited and annealed $(320^{\circ}\mathrm{C}5\mathrm{s})$ pieces from samples I and II. The microstructural information obtained from the SEM photos clearly shows the differences between samples I and II and the reason for the different melting behavior. The cross-sectional micrograph from sample I, which was the simple structure, shows the individual Ti, Pt, Au and Sn layers quite clearly. Note that within the solder layer the original $\mathrm{Au / Sn / Au}$ interfaces are separated in certain regions. From the RBS data, we know that, on average, the middle portion consists of 50:50 at.% Au-Sn with pure Au layers on top and bottom. Perhaps due to the greater strain present in the much thicker layers used in sample I, or to a volume change in the layer during formation of the $\cdot \mathrm{AuSn}$ alloy, empty spaces and gaps occurred in several places within the solder layers of sample I. This is in contrast to the more complex layering used in sample II. In that case, the SEM micrographs show closely packed layers of Au and Sn with no gaps or separations at any of the interfaces. On the basis of this information, the difference in melting behavior of the two samples is easily understood. For sample II, the uniform layering of Au and Sn leads to uniform melting of the solder which, during the 5 s of melting, forms a final layer of Au:Sn 80:20 wt.% solder which has uniform thickness throughout the area of the pad. On the other hand, in sample I, the gaps observed in the solder mean that during a short 5 s annealing, + +there is not enough time for uniform mixing of the AuSn solder, and this leaves areas that are either gold-rich or tin-rich. Since the melting is nonuniform across the pad, areas of different thickness are formed, which greatly differ in their local composition. + +This nonuniformity in melting for sample I is most easily observed in the SEM plan view photos shown in Fig. 8. The dendritic structures found after the melting of sample I are thick regions which are shown to be gold-rich by EDS analysis taken in situ by SEM. A notable difference in the feature size is observed depending on the reflow duration, as is demonstrated in the comparison of the 5-s and 10-s melted samples in Fig. 8. The premature resolidification of these localized features was observed in this case even after only 2-3 s of melting at temperatures above the eutectic melting point, and is consistent with the gold-rich composition as observed by SEM. In contrast to this sample, the plan view photo of sample II after melting shows a uniform surface morphology across the pad. + +Thus, it is concluded that a multiple-layered structure consisting of thinner Au and Sn layers is superior to a simple thick three-layered structure, with respect to melting behavior during bonding. The main reason is that a higher degree of strain and volume change occurs in the simple structure which produces gaps in the interfaces of the solder layer. During the bonding cycle, these gaps interfere with uniform melting behavior and + +![](images/550032dd7f6be7526a4e951de93cfdc980fb2c4afd7379b9a32e54c4d8fa5457.jpg) +(a) + +![](images/35ced04927dd0bb2a130f3b5e0b63110ff790e9a877a3afc50d335b6efc6bdb3.jpg) +(b) + +![](images/9e7659fc8b0ddbd413bc00b42181c54a934bfe71e8a84a6b1eba04ee7f3683af.jpg) +(c) + +![](images/8786b3669d4d12677315d2530b23a74460f45ef10fa5c9965e9e027587d08976.jpg) +(c) +Fig. 7. SEM cross-sectional micrographs of a metallized $\mathrm{Au}(0.8\mu \mathrm{m}) / \mathrm{Sn}(1.4\mu \mathrm{m}) / \mathrm{Au}(0.8\mu \mathrm{m}) / \mathrm{Pt}(0.2\mu \mathrm{m}) / \mathrm{Ti}(0.1\mu \mathrm{m})$ submount, (a) as-deposited, and (b) after a bonding cycle at $320^{\circ}\mathrm{C}$ for 5 s, and of a metallized $\mathrm{Au}(0.5\mu \mathrm{m}) / (3\times) / \mathrm{Sn}(0.35\mu \mathrm{m}) / \mathrm{Au}(0.2\mu \mathrm{m}) / \mathrm{Au}(0.5\mu \mathrm{m}) / \mathrm{Pt}(0.2\mu \mathrm{m}) / \mathrm{Ti}(0.1\mu \mathrm{m})$ submount, (c) as-deposited, and (d) after a bonding cycle at $320^{\circ}\mathrm{C}$ for 5 s. + +create localized areas of Au-rich solder which freeze prematurely. + +# AuSn-Pt interactions + +During a 5 s annealing, interaction of Sn with the Pt layer does not adversely affect the bonding process, especially if the as-deposited structure has a layer 0.5 $\mu \mathrm{m}$ or greater of pure gold directly in contact with the Pt. AES spectra on specially-deposited test structures + +were obtained to determine how much of the Pt barrier layer would be consumed during the typical bonding times of 5-10 s. These data are shown in Fig. 9. Basically, it shows that for short anneal times, Sn does tend to segregate towards the Pt interface, presumably to form Pt-Sn compounds, but that a significant amount of the original Pt layer is left intact during the bonding process. This does not mean that the Pt-Sn interaction is unimportant. As mentioned above, understanding of + +![](images/290ee8978b68e9036de126200d86f6c6414c4d2dc8e2d01c577dcc8d512b6404.jpg) +(a) + +![](images/d3d38711509adf28e246ba2145cb325c40517f07f926811533bf1c8a223b3c1c.jpg) +Fig. 8. SEM plan-view micrographs of a metallized $\mathrm{Au}(0.8\mu \mathrm{m})/$ $\mathrm{Sn}(1.4\mu \mathrm{m}) / \mathrm{Au}(0.8\mu \mathrm{m}) / \mathrm{Pt}(0.2\mu \mathrm{m}) / \mathrm{Ti}(0.1\mu \mathrm{m})$ submount after bonding cycle at $320^{\circ}\mathrm{C}$ for (a) 5 s, (b) 10 s. + +this interaction is very important for establishing and understanding the behavior of the metallization systems under more severe bonding conditions, and the long-term shelf life of the submounts and long-term reliability of the assembly under operating conditions. + +Previously it was reported [14] that the AuSn-Pt reactions take place both in the solid state and at the solid-liquid interfaces. The activation energy for the mutual diffusion process in the solid state was determined to be $1.35 \, \mathrm{eV}$ , for samples that had equal thicknesses of both AuSn solder and Pt film (about $300 \, \mathrm{nm}$ each). In the standard bonding procedure, however, $3 \, \mu \mathrm{m}$ of AuSn solder are deposited onto + +![](images/7f1ac21f83678d16b928a173c726f4ca6ab2da468e71b9632aea159e752a7f1a.jpg) + +![](images/bf4dee588cfe8491036343f3949a879fef7d03c1e269572c5e9dd201386a4cc6.jpg) +Fig. 9. Auger spectra of structure with $\mathrm{Au}(0.2\mu \mathrm{m}) / \mathrm{Sn}(0.25\mu \mathrm{m}) /$ $\mathrm{Au}(0.1\mu \mathrm{m})$ solder deposited on $\mathrm{Ti}(0.1\mu \mathrm{m}) / \mathrm{Pt}(0.2\mu \mathrm{m})$ barrier metals. (a) As-deposited; (b) annealed at $320^{\circ}\mathrm{C}$ for $5\mathrm{s}$ . + +![](images/ffb4eab82cdd18564a9b1f47deb9208784d564cda829f6bc01881a1539bf6121.jpg) +Fig. 10. Arrhenius plot of Sn-Pt reacted layer thickness squared divided by the reflow time for solid-state reactions vs. inverse temperature, at temperatures below eutectic melting point. + +0.1–0.2 $\mu$ m of Pt, and in this structure, almost complete consumption of the Pt by the Pt-Sn interaction was observed after 10–15 s of AuSn melting throughout the bonding. The Pt-Sn reaction was found to be thermally activated and not self-limiting as long as the supply of Sn is not exhausted. According to recent studies [28], the Sn-Pt solid state interaction activation energy was found to be $1.7\mathrm{eV}$ (see Fig. 10). It was also found that the entire Pt layer (up to $0.4\mu \mathrm{m}$ thick) may be consumed at the processing temperature, during storage and through solder deposition steps. Once Pt reacts with AuSn to form the PtSn intermetallics at the Pt/AuSn interface, as well as dissolving into the molten AuSn solder to form a ternary $\mathrm{Au_xSn_yPt_z}$ compound, it no longer acts as an effective diffusion barrier. Thus special attention must be paid to the thermal energy that is introduced into the Ti/Pt/AuSn metallization + +system before, during and after the bonding cycle to eliminate these reactions. As a consequence of the complete consumption of Pt, Ti will have intimate contact with Au, which is well known to motivate extensive formation of intermetallics as a result of the reactive nature of the Au-Ti system. This may jeopardize the short- and long-term electrical, mechanical and thermal performance of the bonded assembly. + +Further studies on the stability of the Ti/Pt/AuSn system were obtained by calorimetric (DSC) measurements of samples containing the above-described AuSn multilayer eutectic structures $(3\mu \mathrm{m}$ thick), which had been deposited onto $\mathrm{SiO}_2$ inert substrates and onto various thicknesses of Pt layers. Figure 11 provides a summary of this analysis. During heating of the AuSn/ $\mathrm{SiO}_2$ sample, a sharp peak is observed at $278^{\circ}\mathrm{C}$ , which is the exact AuSn eutectic melting temperature. Having the correct eutectic composition results in a heat of transition of $8.7\mathrm{calg}^{-1}$ through the melting reaction at $278^{\circ}\mathrm{C}$ . Samples containing an AuSn layer on both the $50~\mathrm{nm}$ and $100~\mathrm{nm}$ Pt layers exhibited the eutectic melting endotherm superimposed on a sharp exotherm as well as a broad endotherm centered near $290^{\circ}\mathrm{C}$ . There was also evidence of further endotherms in the temperature range $370 - 390^{\circ}\mathrm{C}$ . When depositing the solder onto a $200~\mathrm{nm}$ layer of Pt no endotherm corresponding to the eutectic transition is observed, but only a broad exotherm beginning at about $250^{\circ}\mathrm{C}$ and a well-defined endotherm at $370^{\circ}\mathrm{C}$ . + +The exotherm that occurred at around the eutectic temperature was attributed to Pt-Sn reactions, which were found to be rapidly accelerated when the solder was molten. When sufficient Pt is present, this mixing takes place and can be observed even below the solder melting point, implying that a rapid solid state diffusion has occurred. + +![](images/f372ea08330df0f464b5c076decb08479e2c59ce4db62d05a01fa5fa233cdbb9.jpg) +Fig. 11. DSC scan of AuSn solder film (3 $\mu$ m thick) deposited onto $\mathrm{SiO}_2$ and on Pt layers with various thicknesses. + +Once the DSC was completed, the samples were cooled to $200^{\circ}\mathrm{C}$ and annealed at this temperature for $40\mathrm{min}$ , followed by secondary DSC cycles, as shown in Fig. 12. The samples that contained a relatively thin Pt layer ( $50\mathrm{nm}$ , $100\mathrm{nm}$ ) still exhibited the eutectic transition at $278^{\circ}\mathrm{C}$ , although with diminished intensity, reflecting the reduction in the total amount of AuSn eutectic compound as a result of the partial consumption of the Sn by the Pt-Sn interaction. The AuSn heat of transition has changed from an unperturbed value of $8.7\mathrm{cal~g}^{-1}$ to $4.2\mathrm{cal~g}^{-1}$ in the presence of a $50\mathrm{nm}$ Pt layer, and to $1.1\mathrm{cal~g}^{-1}$ in the presence of a $100\mathrm{nm}$ Pt layer. Once a $200\mathrm{nm}$ layer of Pt was introduced under the AuSn solder, no melting transition was evidenced through the second DSC cycle at the solder eutectic melting temperature. The exotherms observed in the initial scans are not shown at any of the subsequent DSC studies, suggesting that the AuSn has been saturated by complete consumption of all the available Pt, and no eutectic melting reaction took place in this stage. + +The DSC results further emphasize the major deficiencies of the Ti/Pt/AuSn bonding system. No matter how thick the Pt barrier layer, it will always be partially or fully consumed through both solid state Pt-Sn interfacial reactions, and by dissolution into the AuSn eutectic liquid, once heated to $278^{\circ}\mathrm{C}$ . As a result, a two-fold problem arises. A complete loss of the Pt layer takes place, allowing for further intermixing of the Ti adhesion layer, initially adjacent to the submount, with the liquid Au-Sn-Pt, which eventually leads to delamination of the entire metallurgical system and mechanical failure of the bonded joint. In addition, the AuSn eutectic solder stoichiometry deteriorates as a result of dissolution of the Pt, and subsequently the Ti into it, and the solder freezes prematurely and cannot be + +![](images/e6421d9868cdc3a2a3caef33217f435fe3e54729808b3a94793b636368fa07b5.jpg) +Fig. 12. DSC secondary scan of samples of AuSn solder films (3 $\mu$ m thick) deposited onto Pt layers with various thicknesses, after annealing at $200^{\circ}\mathrm{C}$ for $40\mathrm{min}$ . + +remelted at the bonding temperature to permit a successful bonding process. + +Based on the above, by executing the 5 s bonding cycle, we have established the multilayer solder structure as our commonly used solder for the purpose of bonding InP-based laser devices to CVD-diamond submounts. Figure 13 presents an optical micrograph of a metallized CVD-diamond submount (Fig. 13(a)), and of the same submount after bonding the InP-based laser device onto it, and shearing it off (Fig. 13(b)). The center metal geometry represents the Au-Sn multilayered solder bonding pad which is deposited onto the barrier Pt-Ti larger area of metallization. The Au-Pt-Ti areas on the right-hand side of this submount are used as a ribbon or wire bonding pad to connect the back side of the laser device to the submount, as well as the submount to the laser package. The InP-based laser device was bonded to the CVD-diamond submount in a junction-down configuration, using a thermal cycle of $320^{\circ}\mathrm{C}$ for about 5 s. Figure 13(b) shows the almost complete dielectric and metal pattern of the front side of the InP-based laser device. One can conclude that the quality of wetting of the device to the CVD-diamond submount is excellent, leading to adhesion of almost $100\%$ of the device front area to the submount. Thus, the shearing results in loss of adhesion in the semiconductor/metal interface comprised in the laser structure, rather than in the device metal/dielectric-solder interface, leaving the bonded metal, and dielectric films bonded to the submount, as observed in the micrograph. + +Figure 14 presents an optical micrograph of an InP-based laser device bonded to a CVD-diamond submount which is then bonded to a BeO testing stud, as well as a schematic description of the structure. This bonded device configuration represents the state-of-the-art bonding technology in which an AuSn multilayered + +solder $(3\mu \mathrm{m}$ thick) is used on top of $\mathrm{Pt}(200~\mathrm{nm})/$ $\mathrm{Ti}(100\mathrm{nm})$ barrier metals adjacent to the CVD-diamond submount. This bonding configuration gave the InP-based laser device a superior bonding performance compared to the traditionally used bonding configurations, such as In solder on BeO submount. Improvements of up to $30\%$ in the device rollover current, power dissipation and optical radiated light were observed, as partially illustrated in Fig. 15. + +In summary, it is obvious that the combination of the complex nature of the AuSn solder system, and the fact that Sn-Pt is such a reactive system suggests that the overall Ti/Pt/AuSn bonding metallurgical system is thermodynamically unstable, and while being used, only a narrow bond-processing window can be explored to provide satisfactory results. Thus some studies have been initiated to evaluate potential alternatives for the solder and the barrier metal systems, in order to provide a more reliable and robust bonding process. The remainder of this review introduces a few of these alternative metals and metallurgical systems. + +# Ti/Ni/AuSn bonding scheme + +Figures 3(a) and 3(b) present the binary phase diagrams of the Ni-Sn and Ni-Au systems, showing the potential for formation of few intermetallics through solid-state reaction in the first system, and the very inert nature of the second system. In addition the Ni, in contrast to Pt, dissolves into Sn-base solders at very low rates [13]. Thus, it appears that Ni would behave as a more suitable barrier metal than Pt between the Ti and the AuSn solder. + +In order to evaluate the Ti/Ni/AuSn bonding system performance in comparison to the Ti/Pt/AuSn system two sets of similar samples from the two systems were + +![](images/a5dd0317892a7d322fea6370ee9b976c748c125a11a42fa6e0a7303013dfe267.jpg) + +![](images/4d30ded04f552856d27ca6186a4c29045f7a1d67e5c9d18444b145c14bbc8ec2.jpg) +Fig. 13. Metallized CVD-diamond heatsink submount (a) as processed, and (b) after bonding InP-based laser device onto it and shearing it off (note the device's front side metal/dielectric film which was sheared off the device and left bonded to the submount). + +![](images/1abce780ee39fc10d3c8393250fa3c905338faed7cd934ab998fa0477c1935dd.jpg) +(a) + +![](images/044a1afcd9b34230abe368c3adc27082a81db43ae7fdfd9255219f7dd03f10b4.jpg) +Fig. 14. Optical micrograph of an InP-based laser device bonded onto a AuSn/Pt/Ti metallized CVD-diamond submount on a BeO testing stud, and a schematic description of the configuration. + +![](images/880e2d5c1e2947f6e6fa990c7e3c755a36f9d0248cbb23e33fe7eeb1540e54e2.jpg) +Fig. 15. Comparison of laser diode performance while bonded by AuSn solder to CVD-diamond submount, and while bonded by In solder to BeO submount. + +prepared, containing similar Ti layers (100 nm thick) and layered structures of AuSn solder (3 $\mu$ m thick). The Ti/Ni/AuSn samples showed longer SLFT and GFT than the Ti/Pt/AuSn samples. The Ti/Ni sample that was heated to 320 °C did not freeze completely before the power supply was turned off (a period longer than 60 s), while the Ti/Pt sample resolidified completely within 27 s after melting. Another set of comparable Ti/Ni and Ti/Pt samples was prepared, with structures of Ti/Ni(300 nm) and Ti/Pt(400 nm). These two samples possessed the same solder arrangement as shown in Fig. 7. The Ti/Ni sample had a nearly five times longer SLFT (45 s) than the Ti/Pt sample (9 s), showing that the molten state of the Ti/Ni/AuSn system was more stable than that of Ti/Pt/AuSn. SEM micrographs verified that the melted/solidified Ti/Ni/AuSn surface was much smoother than that of the Ti/Pt/AuSn system, as shown in Figs. 16 and 17 for the Ti/Ni and Ti/Pt samples, respectively. This suggests that the Ti/Ni/AuSn system should provide a better bonding platform than the Ti/Pt/AuSn system. + +Figure 18(a) shows the AES depth profile of the Ti/Ni/AuSn as-deposited sample, which presents a limited interdiffusion between Au and Ni in the Au/Ni interface, whereas the Au and Sn in the solder have mixed completely even prior to reflow, as was previously reported [13]. After $1\mathrm{min}$ of reflow, a significant amount + +![](images/e6a7a0582382fe28b89aea057d9dc1f2463b243978939f745cc1dcb4bcb1ffb1.jpg) +Fig. 16. SEM micrograph of the surface morphology of a Ti/Ni/AuSn sample after reflow at $320^{\circ}\mathrm{C}$ for $60~s$ . + +![](images/0cf3d704ef8623ce448223b74d1082404b22de77be53a45475989720b93272d2.jpg) +Fig. 17. SEM micrograph of the surface morphology of a Ti/Pt/AuSn sample after reflow at $320^{\circ}\mathrm{C}$ for 60 s. + +![](images/8369d4a1911bb5c3ea3b7927c71f9631ea44ceda8edba26aeef76727c4709e9e.jpg) + +![](images/5c03e3f71918d3a0326f4021202aa9ebb13d64e32bfc50db96386314a90e884b.jpg) + +![](images/757311e686418ed6ff92290de4cc76301511494ae95012bddd9597428ea3dda8.jpg) + +![](images/313ca542278e1966e5d43563e9ac3e483d6b5e679a65bd787f99f78a2810ebcd.jpg) +Fig. 18. AES depth profiles of Ti/Ni/AuSn structure (a) as-deposited and after reflow at $320^{\circ}\mathrm{C}$ for (b) 1 min, (c) 3 min, (d) 5 min. + +of Ni was dissolved into the molten solder. Only 100 nm out of the $300\mathrm{nm}$ of the original Ni layer remained after the reflow, as shown in Fig. 18(b). It is interesting to note that the Au did not diffuse as far as the Sn into the Ni layer. After 3 min of reflow, as shown in Fig. 18(c), about $50~\mathrm{nm}$ of the pure Ni layer was still unreacted, however, with an observable limited migration of Au into the Ni. Au did not diffuse toward the Ti layer, as shown in Fig. 18(d). After two more minutes of solid state diffusion, Sn diffused almost all the way to the Ni/Ti interface, and only $25~\mathrm{nm}$ of pure Ni remained unreacted after this duration of reflow. The penetration of Au into Ni layer was still suppressed. Thus, it is clear that Ni is a better barrier metal than Pt for prohibiting the migration of Au in this system. + +From the above-discussed reactions that took place in the Ti/Ni/AuSn system, and from the earlier discussion about the reactions in the Ti/Pt/AuSn system, one can see that approximately the same thickness of Ni and Pt layers were dissolved by the molten Au-Sn solder during the first minute of reflow. A more severe premature freezing problem, however, was observed in the + +Ti/Pt/AuSn samples. It is thus clear that more Sn was consumed by the Pt than by the Ni, to form the related intermetallic compounds in the Ti/Ni/AuSn and Ti/Pt/AuSn systems, suggesting that the Ti/Ni/AuSn system is more thermally stable. + +A mechanism to describe the reflow reaction kinetics, which comprises five major stages, is suggested: + +(i) Melting: when the temperature of the solder exceeds $282^{\circ}\mathrm{C}$ , the Sn starts to melt and Au is dissolved into the liquid Sn through the formation of the eutectic composition. Complete melting usually occurred before terminating the reflow heating cycle. +(ii) Formation of intermetallic compound: when the molten solder contacts the barrier metal, Ni or Pt, ternary related intermetallic phases (containing Ni or Pt, Sn and Au) are formed. Parameters such as the relative dissolution rates of Ni or Pt, Sn and Au, related intermetallic phases in the liquid Au-Sn solder, and the formation rates of the intermetallic compounds affect the residual amount of unreacted Ni. According to the AES results, once the ternary related intermetallic phases are formed, they might effectively retard the + +A dissolution rate of Ni into the melt due to their high stability at the reflow temperature. + +(iii) Solid state interdiffusion: through the reflow process, intermetallic compounds keep forming and coalescing through intergrain diffusion mechanisms. In spite of the existence of stable intermetallic compounds, Sn and Au in the solder will still react with pure Ni after diffusing through the grain boundaries of the intermetallic compounds. At this stage, the diffusion rate of Sn and Au in these intermetallic compounds dictates the Ni consumption rate. +(iv) Solidification of solder: once the Sn has reacted with Ni, starting from stage 2, a local depletion of the Sn exists. According to the Au-Sn phase diagram, the freezing temperature of the AuSn solder increases by about $30^{\circ}\mathrm{C}$ for every deviation of $1\mathrm{wt.\%}$ of Sn from the eutectic Au-Sn (80:20) solder composition. Therefore, a local partial freezing always occurs as the Sn reacts with the Ni. When the concentration of the Sn in the solder drops below the eutectic composition, its melting temperature is higher than the pre-set reflow temperature, and solidification will occur at local sites, leading to subsequent total freezing of the solder. +(v) Solid state interdiffusion of the various elements: occurs with very low interdiffusion rates of the Sn from the solid solder to the base metal layers. + +In conclusion, several major advantages were discovered from using a Ni barrier layer instead of Pt at the Ti/Metal/AuSn bonding metallization scheme. In particular, it was observed that Ni has better thermal stability during the soldering step compared to Pt, and therefore provided a longer freezing time (both local and global), and a final better reflow solder surface morphology. + +It must, however, be emphasized that the Ni-Sn system, similar to the Pt-Sn system, is a very thermodynamically reactive system, and therefore the improvement in the thermal stability performance is not as significant as was expected. This conclusion motivated the search for more stable alternatives to the barrier metals of the bonding scheme. + +# Ti/Cr/AuSn and Ti/W/AuSn bonding schemes + +Based on the binary phase diagrams of Au-W, Au-Cr and Sn-Cr (unfortunately the Sn-W phase diagram does not exist), as shown in Fig. 4, both Cr and W seem to be very stable barrier metals and thus suitable to replace the Pt at a simple metallurgical bonding scheme. + +The compositional stability of the CVDD/Ti/Cr/AuSn and CVDD/Ti/W/AuSn systems was studied using AES and RBS analyses. Representative data on the stability of the different barrier layers with respect to heat treatment is summarized in Figs. 19 and 20. The AES + +![](images/00c45288bc55bd5c909c9fc5195c209447a2d2fc9ba00ed8a574b16765d0cf30.jpg) + +![](images/cd3121d97addc62c7b7e274b2284e3787b57b7125e523cb6ae6d28a470276f9f.jpg) + +![](images/68ffc71fa8c07cb30fcdd176e93a62950522d9f453be44d132c4c147281c06eb.jpg) + +![](images/33d4acd0c2e8b8525fbf15bb6afbf058b221aced70b2e565fa597adaaeee390e.jpg) +SPUTTER TIME (min) +Fig. 19. Comparison of AES depth profiles from annealed submounts having four different metallization barrier schemes. + +![](images/54822fe5235e3671ea16670be13f8daa528906fe442af9604c7b92bed3b47eb3.jpg) + +![](images/2599ea3c3ab0f74961bafcd4dc2abafd6911b2f2e56bb6e80d2a0ea0da2da417.jpg) + +![](images/dfbaba29cf1a1aea7d48872e863c72ed5e383bfee851807225cbda5d62daf3c3.jpg) + +![](images/336a80ffa068ca1e24c0255df3ca7a2c5b318ab77fb4cbddeba3099769ff4de7.jpg) +Fig. 20. Comparison of RBS spectra from submounts having four different metallization barrier schemes, as-deposited (—) and annealed $(\cdots)$ . + +data shown in Fig. 19 gives concentration depth profiles for the various elements present in four annealed samples. These data provide information on the intermetallic reactions that occur between the solder and the various barrier metals. The RBS data is collected in Fig. 20 and shows both as-deposited and annealed spectra from four samples. The RBS data can be interpreted to give the average AuSn ratio and overall thickness of the solder on a particular submount. For reference, Figs. 19(a) and 20(a) show the AES and RBS data from a sample containing the standard Ti/ Pt barrier metals, which are known to be reactive with the AuSn solder. Figures 19(b) and 20(b) show spectra from the specimen with the Ti/Cr barrier, and Figs. 19(c), (d) and 20(c), (d) show the data from the W/ + +Ti barrier (two different thicknesses of Ti wetting layer were investigated). + +The AES data show a very clear difference between the behavior of the standard Pt/Ti barrier and the other two barrier metals after annealing. There are two obvious differences. First, the ratio of Au:Sn concentration changes dramatically with depth for the sample containing the standard Pt/Ti barrier, while in the samples containing either Cr or W barrier, the Au:Sn ratio remains constant throughout the depth of the material. Second, interdiffusion of Pt and AuSn is observed in the standard sample, whereas no interaction is detected between the barrier and the solder layers in the samples containing either the Cr or W barrier. In those samples, the Auger spectra show that the Cr and W layers remain inert to the solder at temperatures as high as $350^{\circ}\mathrm{C}$ for melting times greater than 1 min. + +These observations are, indeed, consistent with predictions based on the binary phase diagrams for these systems. Pt and Sn are known to form intermetallics, even at low temperatures. The behavior of the barrier in the standard sample reflects the propensity of Sn to react with Pt at the AuSn/Pt interface. This leads to a depletion of Sn from the top of the solder layer, while it tends to accumulate at the AuSn/Pt interface. At the same time, the integrity of the Pt barrier is degraded as PtSn compounds are formed in the interface region. On the other hand, neither Cr nor W are expected to form alloys with Sn or Au at the temperatures used in this study, and as observed, the interface between solder and barrier layers indeed remains sharp. + +In the case of the W barrier layer, a very thin Ti layer of either 5 or $10\mathrm{nm}$ was introduced between the solder and the W layer to improve the adhesion between the solder and the W layers. After the anneal, the AES data shows that the Ti is still present at the solder/ W interface, although it appears to have interacted mostly with the W layer. + +In contrast to the W system, in the case of the Cr barrier, a Ti layer (100 nm thick) was introduced only between the Cr and CVD-diamond submount to improve the adhesion of the entire metal system to the submount. Some interactions were observed at the $\mathrm{Cr / Ti}$ interface resulting in the formation of Cr-Ti intermetallics. + +The AES data for both Cr and W barrier schemes reflect stable metallization systems during heat treatments of $350^{\circ}\mathrm{C}$ for as long as 5 min. This may suggest the advantage of using them to replace the Pt/Ti scheme in bonding InP laser diodes. It should be mentioned, however, that the typical bonding process is usually performed at lower temperature (about $320^{\circ}\mathrm{C}$ ) and for much shorter times ( $< 30$ s). + +The main difference observed in the RBS data, shown in Fig. 20, between the standard sample containing the + +Pt/Ti barrier and the other samples containing Cr or W barriers is in the surface roughness of the solder layers after being annealed. For the standard Ti/Pt sample, the solder layer remains fairly uniform in thickness after the annealing. For that sample, the overall solder composition is about 77:23 wt.% Au:Sn, and the total thickness remains 2.25 μm after the anneal, with an estimated roughness of 0.1 μm. In contrast to this result, the samples with Cr or W barrier show greater solder roughness after annealing. For the Cr sample, the overall Au:Sn ratio is 79:21 wt., and the total thickness is about 1.7 μm after annealing with a roughness of 0.4 μm. The W samples show even greater roughness after annealing, and the solder layer appears to become thicker. For the W sample, the overall Au:Sn ratio was found to be 80:20 wt., and the total thickness increased from 1,8 to 2,5 μm after annealing, with a surface roughness of about 0.7 μm. + +These RBS data indicate a potential problem with the wetting of the AuSn solder to more inert barriers, especially to W. The increased thickness and roughness of the solder layer surface measured after annealing is probably due to the tendency of the AuSn solder to 'ball up' on the W barrier layer as a result of surface tension and insufficient wetting. In the standard Ti/Pt system, the interaction of the AuSn solder with the Pt layer to form PtSn compounds at the interface is detrimental to the performance of the barrier layer, but actually improves the wetting of it, and thus the integrity of the system. + +Figure 21 presents optical $(100\times)$ micrographs of the metallized CVD-diamond submounts that were used as heatsinks for the bonded InP laser diodes. The various metallization areas are detailed in the attached schematic for both the Ti/W/Ti system (Fig. 21(a)) and Ti/Cr system (Fig. 21(b)). These two submounts were reflowed at $350^{\circ}\mathrm{C}$ for $5\mathrm{min}$ under $\mathrm{H}_{2}$ $(15\%)$ $\mathrm{N}_2$ $(85\%)$ ambient to study the stability of the bonding pad metallurgy during a long heat treatment at the bonding temperature. The same heat treatment applied to the standard Ti/Pt/AuSn system led to resolidification of the AuSn solder after $5 - 10\mathrm{s}$ as a result of the significant change in stoichiometry due to the reaction with Pt. In the Ti/W/Ti/AuSn and Ti/Cr/AuSn systems, the AuSn solder remained liquid throughout the $5\mathrm{min}$ reflow, providing clear evidence of the inert nature of AuSn on Cr or W barrier metals, as observed by AES. This long reflow process, however, led to the formation of a rough surface morphology containing pits and holes, in particular for the W/Ti system. It should be emphasized that when these submounts are maintained under more realistic bonding conditions, namely at $320^{\circ}\mathrm{C}$ for up to $20\mathrm{s}$ , pitting phenomena are hardly observed and the surface of the solder pad is much smoother. + +![](images/38262865e4bfb9717517b9fe75db721c5bd6f6e52616651ffdbda01ba0c0879c.jpg) +Fig. 21. Optical micrographs and schematic description of (a) W/Ti- and (b) Ti/Cr-based metallized CVD-diamond submounts after annealing at $320^{\circ}\mathrm{C}$ for $1\mathrm{min}$ . + +![](images/e818d8acaa6eb556eccec4db082716f38870251ac07392dc175524a62abafb9a.jpg) +Fig. 22. SEM micrographs of (a) W/Ti/AuSn and (b) Ti/Cr/AuSn solder pads on CVD-diamond submounts after annealing at 320 °C for 1 min. + +Figure 22 shows SEM micrographs of the Ti/W/Ti/AuSn (Fig. 22(a)) and Ti/Cr/AuSn (Fig. 22(b)) solder pad surfaces to demonstrate the density and size of the pits after reflow at $350^{\circ}\mathrm{C}$ for $5\mathrm{min}$ . For the Ti/W/Ti system, the phenomenon appears to be more severe, with dense pitting and hole sizes of $2 - 20~\mu \mathrm{m}$ . The small distance between the holes led to the for + +mation of a crack network on the AuSn solder surface. For the Ti/Cr system the pit size is much more uniform $(\sim 2\mu \mathrm{m})$ and the overall density of the pits is less than in the Ti/W/Ti system. + +In order to define the interface in which the dewetting took place, an SEM/EDAX analysis was used and the results are shown in Fig. 23. For both the Ti/W/Ti/AuSn and Ti/Cr/AuSn systems, EDAX profiles were taken at a point on the solder surfaces $(\mathbf{a}_1$ and $\mathbf{b}_1$ respectively), and inside holes $(\mathbf{a}_2$ and $\mathbf{b}_2$ respectively). The EDAX spectrum of point $a_1$ shows only Au and Sn, while that of $a_2$ shows mainly W and a small amount of Ti. The spectrum of $b_1$ also shows only Au and Sn, while $b_2$ shows mainly Cr. From these data one can conclude that in both systems the formation of thermal stresses throughout the thermal cycle led to the 'ball up' and delamination of the AuSn solder from the Ti/W/Ti and Ti/Cr layers. + +These phenomena can be disastrous when attempting to bond laser devices onto a submount, such as CVD-diamond. In Fig. 24 an example is given of a delamination of this type, which took place on a CVD-diamond submount subsequent to the bonding of an InP-based laser device onto it. The SEM micrographs clearly show the propagation of the delamination from mechanical damage that was introduced unintentionally into the solder pad during the bonding cycle. In this case, the mechanical delamination may lead to failure of the assembly as a result of the device losing adhesion and being sheared off the submount. Furthermore, and probably more disturbing, the delaminated metal flap may easily block the passage of light from the laser + +![](images/f93f8c63dc3609fe353bbe1c0d1f45ad21507dcdcc0fc380b11fcce38e8423d8.jpg) + +![](images/d4220d5ea7a0de3afaa0f0f3e3b97fde70e4f6cc619ceec01f8e0854d09329b5.jpg) + +![](images/50219d70f3289bf578e0eb192d71f6eb323f5af6aba8c1c05cc556bad2b4fded.jpg) + +![](images/ab84d06490d77f60cfac784ee832246fdc844d568f20c5b7fbb4f5d39bd38135.jpg) +Fig. 23. SEM micrographs of the immediate area of a pit at the (a) W/Ti/AuSn, and (b) Ti/Cr/AuSn samples (Fig. 22), and EDAX analysis of the surface around the holes and inside the holes. + +device, which is supposed to be incorporated into the attached fiber optics. This leads to optical failure of the entire packaged assembly, which is usually very costly. + +In summary, both Cr and W metals provided excellent inert barrier metal layers under the AuSn solder when applying the Ti/Metal/AuSn scheme, which is the scheme typically used for bonding laser devices to CVD-diamond submounts. The eutectic composition and the liquid state of the solder were maintained at $320 - 350^{\circ}\mathrm{C}$ throughout the entire bonding cycle and for a much longer thermal cycle (5 min) than were possible when bonding with Pt as the barrier layer (no longer than 15 s). This permits a major reduction in the AuSn solder pad thickness to about $0.5\mu \mathrm{m}$ , and thus improves the heat dissipation from high-power laser devices. + +# Ti/W metaly/AuSn bonding schemes Introduction + +In the previous section, we discussed the major disadvantage of Ti/Cr/AuSn and Ti/W/Ti/AuSn, namely the poor wettability of the AuSn solder onto the barrier metals, as a result of the almost completely inert nature of both W and Cr to AuSn. In order to improve the wetting performance of the solder to the submount while maintaining the extremely long AuSn freezing time as a result of the inert interfacial nature, we extended our investigation of adhesion and wetting layers to be introduced between the barrier metal layer + +![](images/25b042cee8b5e1e5187849e5877f77be15eabb668780a42b1abc3988f5a2781e.jpg) + +![](images/59e37953ffd73147512d17405fea12d1a5bbf22e4225be94eb8a3def44bfd0b6.jpg) + +![](images/a1606f80738d7e37878d317b406a1c12873c0023156e78f29c1d8dc8c7a8c011.jpg) +Fig. 24. SEM micrographs of examples of AuSn solder delaminated from metallized CVD-diamond submount. + +and the solder. This concept is illustrated schematically in Fig. 25, and its nature and performance are discussed below. + +# Conceptual approach + +Samples of Ti/W/AuSn, Ti/WAu/AuSn, Ti/W(AuSn)/ AuSn and Ti/W(Ni3Sn2)/AuSn systems were studied; the actual structures are given in Table 1. A summary of the AuSn solder surface local freezing times (SLFT) + +![](images/71ae10cf1b4d76faf0c662f7e7e594a0f2281fa4e6b08e5da6599f32fe6d4b4a.jpg) +Fig. 25. Schematic presentation of conceptual design of advanced bonded device assembly. + +and global freezing times (GFT), which were measured while heating the samples under forming gas at ambient and in liquid flux, is given in Table 2, and will be discussed in detail later. The quality of the adhesion between the solder and the metal layer structure is also given in this table, in order to provide a qualitative measure of the integrity of the metallization schemes. These results and the SEM-EDS, RBS and AES data will be discussed. + +# Ti/W/AuSn + +The $\mathrm{Ti}(100\mathrm{nm}) / \mathrm{W}(200\mathrm{nm}) / \mathrm{AuSn}(3000\mathrm{nm})$ system, which was studied as the basic metallization scheme used for bonding of the InP laser diode to the heatsink, was discussed earlier in this review. The SLFT and GFT of this sample (see Table 2) were longer than 500 and $4000\mathrm{s}$ , respectively, which were the best freezing durations measured in the entire study. The lack of interfacial reaction between AuSn and W eliminated the degradation of the eutectic solder composition, and thus the premature solder freezing. However, this metallization provided insufficient adhesion of the solidified + +solder to the metallization under it, as a result of poor wetting performance during the melting, while reflowing the solder. + +# Ti/W/WAu/Au/AuSn [30] + +The Ti/W/WAuSn/AuSn system was proposed in order to improve the adhesion of the AuSn solder to the W barrier layer. The system studied comprised the following multilayer scheme: Ti(100 nm)/W(100 nm)/ WAu(200 nm)/Au(100 nm)/AuSn(300 nm). SLFT and GFT were measured to be only about 20 s and 500 s, respectively. + +A generic mechanism to describe this melting, interactions and freezing of the solder, containing five stages, was assumed. First the Sn of the solder diffuses into the WAu codeposited layer. Subsequently, it reacts with the Au to form AuSn intermetallics which causes local Sn depletion zones at the solder and, as a result, premature local freezing in the liquid solder. After the formation of the AuSn intermetallics in the WAu codeposited layer and in the interface of the molten solder and the solid WAu codeposited layer, the indiffusion rate of the Sn into the WAu layer is reduced and thus the GFT is extended significantly. Therefore, even after a long reflow $(>10\mathrm{min})$ at 320 or $350^{\circ}\mathrm{C}$ , the AuSn solder did not solidify completely prior to turning off the plate and cooling the samples. + +Thus, it is clear that by adding the codeposited WAu layer between the W and the AuSn, the adhesion of the metals was improved, however, at the price of significantly shortening the SLFT and GFT. + +In addition, it was found that while heating the Ti/ W/WAu/AuSn samples to a higher temperature of $350^{\circ}\mathrm{C}$ , in both liquid flux and under forming gas, a limited dewetting was still observed. Figure 26(a) shows an optical micrograph of the sample that was reflowed under forming gas at $350^{\circ}\mathrm{C}$ for $5\mathrm{min}$ . Figure 26(b) + +TABLE 1. Metallization schemes studied to optimize bonding + +
Layer12345678
AuSn solder (F)AuSn(80% Au)AuSn(77.5% Au)AuSn(80% Au)AuSn(77.5% Au)AuSn(80% Au)AuSn(77.5% Au)AuSn(77.5% Au)AuSn(77.5% Au)
3 μm3 μm3 μm3 μm3 μm3 μm3 μm3 μm
Wetting layer (E)Ni3Sn2(200 nm)Ni3Sn4(100 nm)Ni3Sn4(200 nm)Ni3Sn4(400 nm)
Second adhesion layer (D)W(Ni3Sn4)2(100 nm)
First adhesion layer (C)(WAu2)(200 nm)W(AuSn)2(200 nm)W(Ni3Sn2)2(200 nm)W(Ni3Sn2)2(200 nm)W(Ni3Sn2)2(100 nm)
Barrier layer (B)W(100 nm)W(100 nm)W(50 nm)W(50 nm)W(50 nm)W(50 nm)W(50 nm)W(50 nm)
Adhesion layer (A)Ti(100 nm)Ti(100 nm)Ti(100 nm)Ti(100 nm)Ti(100 nm)Ti(100 nm)Ti(100 nm)Ti(100 nm)
+ +TABLE 2. Thermal stability and wetting performance of the various metallization structures studied + +
No.Metallization schemeFreezing timeW/AuSn Adhesiona
SLFT (s)GFT (s)Forming gasPEG flux
320 °C350 °C320 °C350 °C320 °C350 °C320 °C350 °C
1Ti(100 nm)/W(100 nm)/AuSn(3 μ)>500>500>4000>4000----
2Ti(100 nm)/W(100 nm)/WAg2(200 nm)/AuSn(3 μ)2030>500>500+---
3Ti(100 nm)/W(50 nm)/W(AuSn)2-(200 nm)/AuSn(3 μ)2520>500>500+-+-
4Ti(100 nm)/W(50 nm)/W(Ni3Sn2)2-(200 nm)/AuSn(3 μ)1515>500>500----
5Ti(100 nm)/W(50 nm)/W(Ni3Sn2)2-(200 nm)/Ni3Sn2(200 nm)/AuSn(3 μ)7565>500>500++++
6Ti(100 nm)/W(50 nm)/W(Ni3Sn2)2-(200 nm)/W(Ni3Sn4)(100 nm)/Ni3Sn4(100 nm)/AuSn(3 μ)180180>2000>500++++
7Ti(100 nm)/W(50 nm)/Ni3Sn4(200 nm)/AuSn(3 μ)300300>500>500+---
8Ti(100 nm)/W(50 nm)/Ni3Sn4(400 nm)/AuSn(3 μ)1515>500>500++--
+ +${}^{\mathrm{a}}\left( -\right)$ poor wetting between AuSn and metal; (+) excellent wetting between AuSn and metal. + +![](images/2e69e7d877a36b0f4787eff3eded9d99aeb012d6dddd366d7481b59bec83f265.jpg) + +![](images/2eddfcea95f8ccdd2d82497e854497d529396b1dd51f245892f1c181f0ff4f3f.jpg) +Fig. 26. SEM micrograph of Ti/W/WAu/AuSn sample, after reflow at $350^{\circ}\mathrm{C}$ for $5\mathrm{min}$ , (a) under forming gas ambient and (b) in liquid flux-PEG400 bath. + +shows the dewetting area of a similar sample that was reflowed at liquid flux. + +As opposed to the EDS spectra that were collected on the dewetted area of the Ti/W/AuSn samples discussed earlier, the EDS spectrum of this sample exhibited detectable Au and Sn peaks, in addition to the observed Ti and W peaks. This result revealed that, in spite of the fact that dewetting occurred during the reflow, some degree of solder-barrier metal interaction indeed took place at the interface. In summary, it is clear that the addition of the AuW intermediate layer did improve the adhesion of the AuSn solder to the W barrier layer, by promoting interfacial mixing and interactions. + +# Ti/W/W(AuSn)/Au/AuSn + +The Ti/W/WAuSn/Au/AuSn system was evaluated in order to further improve adhesion of the AuSn solder to the W barrier layer. The sample comprised the following layer structure: Ti(100 nm)/W(50 nm)/W(AuSn)₂(200 nm)/Au(100 nm)/AuSn(3000 nm). A small amount of Au was added to the solder in order to compensate for the extra amount of Sn at the W(AuSn)₂ interfacial layer. + +SLFT and GFT of 25 and $>500\mathrm{s}$ , respectively, were measured in this sample after reflow. The composition of the as-deposited AuSn solder was 81 wt.% Au-19 wt.% Sn, and the composition of the reflowed AuSn + +![](images/1a2d02aadab2877d616fdb91a21b54affd580571685e69331537f5689fd0cf37.jpg) + +![](images/64c2eb8e33624f3cc282edee5638325c6626503a7987f478ca145e27e0929375.jpg) +Fig. 27. Optical micrograph of Ti/W/W(AuSn)/AuSn sample after reflow at $350^{\circ}\mathrm{C}$ for $5\mathrm{min}$ , (a) under forming gas ambient and (b) in liquid flux-PEG400 bath. + +![](images/5b99c6c99b47727cb5aa2295019f1e65306e673b716d853550dcf07aeb3b4dde.jpg) +Fig. 28. EDS spectrum of a dewetted area at the Ti/W/W(AuSn)/AuSn sample, after reflow at $320^{\circ}\mathrm{C}$ for 5 min in liquid flux-PEG400 bath. + +solder was found to be 80 wt.% Au-20 wt.% Sn, which was almost the same as that of the as-deposited solder. + +In this sample the solder did not ball up on heating to $320^{\circ}\mathrm{C}$ , either at liquid flux or under forming gas. But solder balls were observed when the sample was heated to $350^{\circ}\mathrm{C}$ . Figure 27 shows optical micrographs of the dewetted area of the forming gas-reflow sample (Fig. 27(a)) and of the flux-reflow sample (Fig. 27(b)). + +EDS profiles were taken at dewetted areas, and an example is shown in Fig. 28. Au and Sn signals were also detected by the EDS in addition to Ti and W + +peaks. This suggested that some limited level of interaction between the metals and the solder took place also in this system prior to dewetting. + +# Ti/W/W(Ni3Sn2)2/Au/AuSn + +The Ti/W/W( $\mathrm{Ni}_3\mathrm{Sn}_2$ )₂/Au/AuSn metallization scheme was the next system that was evaluated, introducing the W( $\mathrm{Ni}_3\mathrm{Sn}_2$ )₂ codeposited film as a transition layer to improve the adhesion between the AuSn solder and the W barrier layer. The structure that was studied comprised six layers in the following order: Ti(100 nm)/W(50 nm)/W( $\mathrm{Ni}_3\mathrm{Sn}_2$ )₂(200 nm)/ $\mathrm{Ni}_3\mathrm{Sn}_2$ (200 nm)/Au(100 nm)/AuSn(3000 nm). The motivation for introducing a NiSn intermetallic layer was to improve the adhesion of the AuSn to the inert underlying layer. This role of the NiSn compound, discussed earlier in the literature, suggests that the mechanism responsible for the adhesion improvement induced by these intermetallics is attributed to its performance as an excellent wetting layer due to the enhanced formation of the Sn-rich alloys, such as $\mathrm{Ni}_3\mathrm{Sn}_2$ and $\mathrm{Ni}_3\mathrm{Sn}_4$ [19-22]. As a result of these reactions the Sn at the AuSn solder layer depletes relatively rapidly and thus in order to compensate for this depletion, a Sn-rich AuSn solder (24 wt.% Sn) was applied into this system. + +Figure 29 provides the AES profiles of the as-deposited Ti/W/W( $\mathrm{Ni}_3\mathrm{Sn}_2$ ) $_2/\mathrm{Ni}_3\mathrm{Sn}_2/\mathrm{Au}100/\mathrm{AuSn}$ sample (Fig. 29(a)), and of the same sample after reflow at $350~^\circ \mathrm{C}$ for $5\mathrm{min}$ (Fig. 29(b)). It was observed that the initial solder layer ( $3\mu \mathrm{m}$ thick) having a Sn-rich solution of 76:24 (wt.%) Au:Sn, underwent a significant intermetallic reaction during the heating cycle, which resulted in an increase in thickness to about $3.9\mu \mathrm{m}$ and in a depletion of the Sn concentration at the solder to a final stoichiometry of Sn:Au 20:80 wt.%. + +The presence of the NiSn layer was found to improve the integrity of the overall multilayer structure, as was reflected in the excellent adhesion of the AuSn solder to the W layer. The SLFT was measured to be longer than 50 s and GFT longer than 500 s. + +Figure 30 shows a SEM micrograph of the Ti/W/ $\mathrm{W}(\mathrm{Ni}_3\mathrm{Sn}_2)_2 / \mathrm{Ni}_3\mathrm{Sn}_2 / \mathrm{Au100} / \mathrm{AuSn}$ system after reflow at liquid flux at $350^{\circ}\mathrm{C}$ for $5\mathrm{min}$ . The typical phase-separation morphology of a eutectic reaction is clearly observed. As anticipated from the evaluation of the AuSn phase diagram, pure AuSn compound and other Au and Sn solid intermetallics were present in the solidified structure. The brighter phase was identified by EDS as the $\mathrm{Au}_4\mathrm{Sn}$ compound, as shown in Fig. 31(a), and the darker area was identified to be AuSn, as shown in Fig. 31(b). The excellent stability of this system during the reflow heating cycles, at temperatures as high as $350^{\circ}\mathrm{C}$ for a duration of $5\mathrm{min}$ , is clearly shown as well at the AES depth profile (Fig. 29). One can observe the extent of the WNiSn interfacial reaction, + +![](images/8db0683d3c3e2aefcdfe988a1299d544e7aac1dafc0f25c2d349063a81573c63.jpg) + +![](images/591c79dab7f2b55c7dd99a2113ae4808f500a94714bc9c22389d6b7ede99d290.jpg) +Fig. 29. AES depth profiles of the (a) as-deposited Ti/W/ $\mathrm{W}(\mathrm{Ni}_3\mathrm{Sn}_2)_2 / \mathrm{Ni}_3\mathrm{Sn}_2 / \mathrm{Au} / \mathrm{AuSn}$ sample, and (b) this sample after reflow at $350^{\circ}\mathrm{C}$ for $5\mathrm{min}$ . + +![](images/881acf3b7a507bebdaef7925f608c0b1e5acbb19b62ffff6ab5a7f57087d767a.jpg) +Fig. 30. SEM micrograph of AuSn eutectic structure formed on the Ti/W/W(NiSn)/Ni $_3$ Sn $_2$ /AuSn sample after reflow at $300^{\circ}\mathrm{C}$ for 5 min under forming gas ambient. + +which took place in the system and enhances the adhesion of the AuSn to the W while leaving unreacted W under it to allow good barrier performance. + +# Ti/W/W(Ni3Sn4)/Ni3Sn4/Au/AuSn + +$\mathrm{Ti}(100 \mathrm{~nm}) / \mathrm{W}(50 \mathrm{~nm}) / \mathrm{W}(\mathrm{Ni}_{3} \mathrm{Sn}_{4})(100 \mathrm{~nm}) / \mathrm{Ni}_{3} \mathrm{Sn}_{4}$ (100 nm)/Au(100 nm)/AuSn(3000 nm) system was the final modified system that was evaluated in an attempt to extend the SLFT of the previous systems. In this case, two codeposited layers with different codeposited ratios of W and $\mathrm{Ni}_{3} \mathrm{Sn}_{4}$ were introduced between the + +![](images/f113edeeeb5407b5d4af5cbeb59b41563f02cd1341a5a961c5c648877bb8524c.jpg) + +![](images/b1c81e15a06843e6b0bd12fde6561fa23c2d7b26ba52b77c413d2bbed885610c.jpg) +Fig. 31. EDS spectra of eutectic AuSn solder in the Ti/W/ W(NiSn)/NiSn/AuSn sample after reflow at $320^{\circ}\mathrm{C}$ for 5 min under forming gas ambient; (a) at the white area of the eutectic solder, and (b) at the black area of the eutectic solder. + +solder and the W barrier layer in order to provide a better chemical gradient in the transition from the pure W layer to the AuSn solder. A Sn-rich solder was also introduced onto this sample. + +The motivation for adding the $\mathrm{Ni}_3\mathrm{Sn}_4$ intermetallic compound layer was to improve the adhesion of the AuSn solder to the underlying W layer, while minimizing the solder to barrier interfacial reaction. This is essential in order to eliminate the dilution of the Sn from the AuSn solder, which eventually leads to degradation of the eutectic composition and to premature freezing of the solder. The $\mathrm{Ni}_3\mathrm{Sn}_4$ is a promising candidate to fulfill this goal, since it is saturated with Sn and thus will not consume further Sn from the adjacent AuSn solder. Indeed, SLFT of longer than 200 s while heating at temperatures of $320 - 350^{\circ}\mathrm{C}$ were measured for this sample, and a GFT longer than 60 min was measured. + +In order to evaluate the quality of the solder wetting performance, and the nature of the molten solder flow, four different samples with various metallization pad sizes under the solder pad were prepared and tested. These metallization areas are designed to allow the flow of excess molten solder away from the joint-bonded volume under the device. Figure 32 provides optical micrographs of two of the above submounts after reflow under forming gas at $350^{\circ}\mathrm{C}$ for $5\mathrm{min}$ . One can see that the molten solder has flowed to the edges of the metallization pad and completely covered it, regardless of the metallization area size. Even the largest metallization pad, which is shown on the right-hand micrograph of Fig. 33, has been completely covered with molten solder in a uniform fashion by the end of the reflow cycle. This suggests that excellent solder stability + +![](images/baebb3342d8111f75626714b3f5a24b21103a262aa482fdf4ec8925c08545385.jpg) + +![](images/05edb200410773e3659484d01a5ddac98f9dc5dcddc9d54b8ddf4d4f82efe0f9.jpg) + +![](images/9181de757c801303d2749a36bfafcbbc3f48d69edaf686fcd010da7956f1535a.jpg) + +![](images/12958265d21f3f04bc8d9426194e363cc4144e08b889e1ecafa77e491cc01ff4.jpg) +100μm + +![](images/0d65bed63ac46452f80e90ad150d6a0aa9ba71cd1cabdadf650e0477c20a8a6b.jpg) +Fig. 32. Optical micrographs of four Ti/W/W( $\mathrm{Ni}_3\mathrm{Sn}_4$ )/ $\mathrm{Ni}_3\mathrm{Sn}_4/\mathrm{Au}/\mathrm{AuSn}$ metallized submounts, differing from each other by the dimension of the Au-capped solder flowing area under the solder pad. + +![](images/b0c73619a3ddcf9c65de5a6508520f3814a9ab40a2688055bc08feb871060f71.jpg) +Fig. 33. Optical micrographs of two Ti/W/W( $\mathrm{Ni}_3\mathrm{Sn}_4$ )/ $\mathrm{Ni}_3\mathrm{Sn}_4/\mathrm{Au}/\mathrm{AuSn}$ metallized submounts, after reflow in liquid flux at $350^{\circ}\mathrm{C}$ for 5 min. + +was achieved and, as observed previously, the 80:20 Au:Sn eutectic composition was maintained even during the very long $(>60\mathrm{min})$ reflow process, with an excellent solder-to-barrier wetting performance. Figure 34 provides optical micrograph of the solder after undergoing a reflow cycle at $320^{\circ}\mathrm{C}$ for $5\mathrm{min}$ . One can observe the uniform and integrated appearance of the metallization, which has a typical eutectic phase-separation morphology, indicating that the composition of the resolidified solder remained close to the eutectic composition. As predicted from the phase diagram, the $\mathrm{Au}_{4}\mathrm{Sn}$ and AuSn solid intermetallics were, indeed, + +formed and observed in the frozen AuSn eutectic structure. + +Figure 35 provides the AES spectra of the Ti/W/ $\mathrm{W}(\mathrm{Ni}_3\mathrm{Sn}_4) / \mathrm{Ni}_3\mathrm{Sn}_4 / \mathrm{Au} / \mathrm{AuSn}$ as deposited and after reflow at $320^{\circ}\mathrm{C}$ and $350^{\circ}\mathrm{C}$ for $5\mathrm{min}$ . It is clear that the presence of the $\mathrm{W}(\mathrm{Ni}_3\mathrm{Sn}_4)$ wetting layer significantly improved the integrity of the overall multilayer structure resulting in perfect adhesion of the AuSn solder to the W layer. Thus, the application of a W layer as a reaction barrier under the AuSn solder, in which reaction of the AuSn solder with the barrier metals is essentially eliminated, is seen to provide a superior bonding system. + +![](images/592398369b52db8741e76c8db68ab8d6b905bc13904f56a02349da8744a01461.jpg) +Fig. 34. Optical micrograph of the AuSn eutectic structure formed on the Ti/W/W( $\mathrm{Ni}_{3}\mathrm{Sn}_{4}$ / $\mathrm{Ni}_{3}\mathrm{Sn}_{4}$ /AuSn sample, after reflow at 350 °C for 5 min under forming gas ambient. + +![](images/d7327a0cca84962c43a52da0f98dc9241d4f14726aba01b297e75e3aaec834d6.jpg) +Fig. 35. AES profiles of the Ti/W/W( $\mathrm{Ni}_3\mathrm{Sn}_4$ )/ $\mathrm{Ni}_3\mathrm{Sn}_4$ /Au/AuSn samples, (a) as-deposited, and (b) after reflow under forming gas at $320^{\circ}\mathrm{C}$ for $5\mathrm{min}$ . + +Figure 36(a) shows the DSC results (without scan rate calibration) of the Ti/W/W(Ni₃Sn₄)/Ni₃Sn₄/AuSn sample, which was heat-treated at 320 °C for 10 min. The heat of melting (based on total weight of the system) is shown to be about 0.8165 J g⁻¹. The solder started to melt at 281.3 °C. After annealing at 320 °C for 10 min, on cooling the solder solidified at 277.8 °C with a heat release of 0.8025 J g⁻¹. The difference between the heats of absorption and evolution of this metallization-thin solder system after heat treatment at 320 °C for 10 min is less than 2%, which is within experimental scatter. This is in contrast to the previous + +![](images/626ca135627f94fbed2db8aa7f59ce3822d5d970780d0cc4bb1a759226de13db.jpg) + +![](images/43e2a2a6343abdb61925a34cffb80e51629092569d18c96ac15067e6505e457c.jpg) +Fig. 36. DSC spectra of Ti/W/W( $\mathrm{Ni}_3\mathrm{Sn}_4$ )/ $\mathrm{Ni}_3\mathrm{Sn}_4$ /Au/AuSn sample that was heat-treated at $320^{\circ}\mathrm{C}$ (a) for $10\mathrm{min}$ , and (b) for $10\mathrm{min}$ , cooled to $200^{\circ}\mathrm{C}$ , and heated again to $320^{\circ}\mathrm{C}$ for $60\mathrm{min}$ under $\mathbf{N}_2$ gas ambient. + +DSC study of Ti/Pt/AuSn system, in which the sample, when heated at $320^{\circ}\mathrm{C}$ for $10\mathrm{min}$ , solidified completely and showed no indication of freezing upon cooling through the eutectic melting temperature of AuSn. Such behavior indicates that this metallization-thin solder system is so stable that the total amount of the compounds which comprises the eutectic solder in the system has changed only slightly during the heat treatment. + +In order to evaluate the reproducibility of the melting, the sample which had been annealed at $320^{\circ}\mathrm{C}$ for 10 min was reheated to $320^{\circ}\mathrm{C}$ and held at this temperature for $60\mathrm{min}$ . Figure 36(b) gives the DSC results of this experiment. The heat of melting was measured this time to be $0.8017\mathrm{Jg}^{-1}$ and the melting temperature $280.3^{\circ}\mathrm{C}$ . After heat treatment at $320^{\circ}\mathrm{C}$ for $60\mathrm{min}$ , the sample heat of solidification was measured to be $0.8057\mathrm{Jg}^{-1}$ and the freezing temperature $277.1^{\circ}\mathrm{C}$ . The heat of melting and solidification measured in the previous run and in this run were nearly identical. This provides a strong evidence to the fact that the solder has not been prematurely frozen even after many thermal cycles have been executed. + +Figure 37(a) provides the DSC results of a sample that was heat-treated at $350^{\circ}\mathrm{C}$ for $10\mathrm{min}$ . The heat of melting was found to be $0.8028\mathrm{Jg}^{-1}$ and the melting temperature $281.3^{\circ}\mathrm{C}$ . After heat treatment, the heat of solidification was measured to be $0.7909\mathrm{Jg}^{-1}$ and + +![](images/389227e7f08d8f8e5017b9c967bbc171cb6f434610c0d15b903a7e89c441d6fe.jpg) + +![](images/3ab50f092396c4526df137e16ae833054770d4ec3ca2e8631b7a3442acae7181.jpg) +Fig. 37. DSC spectra of Ti/W/W( $\mathrm{Ni}_3\mathrm{Sn}_4$ )/ $\mathrm{Ni}_3\mathrm{Sn}_4$ /Au/AuSn sample that was heated to $350^{\circ}\mathrm{C}$ (a) for $10\mathrm{min}$ , and (b) for $10\mathrm{min}$ , cooled to $200^{\circ}\mathrm{C}$ and heated again to $320^{\circ}\mathrm{C}$ for $60\mathrm{min}$ under $\mathbf{N}_2$ gas ambient. + +the freezing temperature $276.8^{\circ}\mathrm{C}$ . The heats evolved through melting and solidification are very similar, suggesting again the high stability of this system. + +Multicycle reflows of this sample were conducted at $350^{\circ}\mathrm{C}$ as well. Figure 37(b) shows the DSC results of the second thermal cycle of the testing. The heat of melting was found to be $0.7930\mathrm{Jg}^{-1}$ and the melting temperature $279.9^{\circ}\mathrm{C}$ . After a further heat treatment at $350^{\circ}\mathrm{C}$ for $60\mathrm{min}$ , the heat of solidification was measured to be $0.7843\mathrm{Jg}^{-1}$ and the solidification temperature $275.4^{\circ}\mathrm{C}$ . The heats of melting and solidification measured in the first run and the second run were nearly identical. The fact that the solidification temperature was so well conserved provides strong evidence that the solder composition has not changed through these very severe thermal cycles. + +In summary, it appears that the five layered Ti/W/ W(Ni3Sn4)/Ni3Sn4/Au/AuSn system provides the best bonding metallization scheme of all the systems that were evaluated through the current study. + +# Conclusions + +In this review we attempted to examine the performance of the most commonly used metallurgical schemes for the bonding of high-power laser devices + +onto submounts, and in particular onto CVD-diamond submounts. + +In the standard configuration, a $3\mu \mathrm{m}$ layer of eutectic AuSn (80 wt. $\%$ Au) solder ( $T_{\mathrm{m}} = 278^{\circ}\mathrm{C}$ ) is deposited onto a $0.2\mu \mathrm{m}$ Pt barrier layer, which caps a $0.1\mu \mathrm{m}$ Ti adhesion layer adjacent to the submount. This system is easy to deposit and permits good bonding of the device to the submount. The Ti/Pt/AuSn system, and particularly the Pt/AuSn interface, is very reactive thermodynamically, mainly as a result of the almost spontaneous reactions that take place between Pt and Sn, even at room temperature. As a result the Pt consumes Sn from the AuSn to form interfacial PtSn intermetallics through solid state reactions, diluting the Sn from the solder and thus deteriorating the eutectic stoichiometry composition. The stoichiometry varies even further once the solder melts during the bonding cycle, dissolving more Pt into the solder, and leading to premature solidification of the solder. Thus special precautions must be taken during deposition of the various layers and during the bonding process, to limit heating of the multilayer structure and thus limit metallic interactions. It is clear that the Ti/Pt/AuSn bonding metallization scheme is not a robust one, and provides a very narrow processing window (for both temperature and time) to complete a successful bonding cycle. + +In order to improve the bonding process, a more stable metallurgical system is needed. This may involve either replacement of the AuSn solder by a less reactive one, or the use of an alternative barrier layer metal, or both. For InP-based laser bonding applications, it is, unfortunately, impossible to introduce a more stable solder, which by definition will have higher melting and bonding points, since the laser device undergoes accelerated degradation at temperatures higher than $320^{\circ}\mathrm{C}$ . Therefore the more attractive solution is to evaluate an alternative and advanced barrier metallization schemes under the AuSn solder, in order to stabilize the metallurgical scheme and allow for a superior and more robust bonding assembly. The quality of the advanced bonding metallization schemes was evaluated through the thermodynamic stability of the solder, once melted. This was reflected by both the surface and the global freezing times of the solder, as well as the adhesion of the molten and solidified AuSn solder to the various barrier metals. A minimum interaction between the AuSn solder and the barrier metals is required to eliminate solder deterioration, and thus allow for a long solder nonfreezing period. The applied metallization should, however, undergo some interaction with the solder in order to promote sufficient adhesion of the barrier metals to the solder. The microstructure of the various metallization systems, as deposited and + +after reflow at both $\mathbf{N}_2$ ambient and in liquid flux, were studied in detail, and the correlation between it and the solder refreezing time was evaluated. + +The first set of systems tested contained a conceptually similar structure, of $0.1\mu \mathrm{m}$ Ti as an adhesion layer adjacent to the submount, capped with $0.2\mu \mathrm{m}$ of barrier metal, onto which $3\mu \mathrm{m}$ of AuSn solder was deposited. Ni, Cr and W were evaluated as potentially superior barrier layer metals. All of the above metals were found to be less reactive with the AuSn than was Pt; W was found to be absolutely inert. Surface local freezing and global freezing phenomena were not observed during reflow in forming gas or flux, but the solder had dewetted from the W layer, reflecting insufficient interaction and adhesion of the solder to the W layer. + +Therefore, it was concluded that in order to maintain good stability of the solder on the barrier metal, as observed in the Ti/Pt/AuSn system, but improve the adhesion of the solder, the simple bonding metallization schemes were not sufficient, and more complicated systems had to be applied. Not only are adhesion and barrier metal layers required at the submount and under the solder, but an additional adhesion layer must be added between the barrier and the solder layers, to improve the integrity of the entire system. + +The first multilayer that was studied was the Ti/W/ WAu/AuSn system. The adhesion of the solder was improved somewhat, by introducing a WAu codeposited layer between the solder and the W layer. By reflowing the sample under forming gas at $320^{\circ}\mathrm{C}$ , good integrity of the solder to W layer was achieved; however, dewetting still took place at the higher reflow temperature of $350^{\circ}\mathrm{C}$ . Reflowing the sample at liquid flux resulted in complete delamination of the solder from the W layer, even at the lower temperature of $320^{\circ}\mathrm{C}$ . + +Further improvement in system integrity was obtained by using the Ti/W/W(AuSn)/AuSn system. Introducing an improved ternary wetting layer between the AuSn solder and the W layer resulted in entirely eliminating delamination of the solder from the W layer during reflowing of the sample at $320^{\circ}\mathrm{C}$ . Occasional dewetting of the solder from the W layer was, however, still observed during heating the sample at higher temperatures. The improved stability led to extension of the solder local freezing time to about 25 s, and more than 500 s were required for complete solidification of the solder. + +Final modification and optimization of the bonding scheme were achieved by adding a second wetter layer between the adhesion layer and the solder. The entire metallurgical scheme became a 5-layer structure, with the following composition: Ti/W/W(Ni₃Sn₄)/Ni₃Sn₄/AuSn. This system showed excellent thermal stability at all treatment temperatures. Surface local freezing time was measured to be longer than 200 s and the + +global freezing time was found to be longer than 60 min. No dewetting whatsoever was observed in the reflow samples, regardless of the ambient. Thus, it was concluded that introduction of two adhesion layers between the AuSn solder and the W barrier layer provided the optimum solution for achieving perfect adhesion and integrity of the entire multilayer structure, without sacrificing the excellent inert nature of the AuSn solder to the barrier metal provided by the presence of the W layer. + +Thus, it is seen that in order to achieve a robust bonding metallization system, a sophisticated and advanced multilayer technology must be applied. This provides a great challenge to materials and process engineers and designers. The same skills and talent that traditionally were invested in developing stable and robust metal contacts to the semiconductor devices must be used to develop modern bonding metallization schemes as well. Multilayer structures are essential in order to achieve the required device-to-submount bonded assembly integrity. Furthermore, by developing solutions to both the device contacts and submount bonding metallization schemes, which each have special tasks and goals, the metallurgical engineer has to be concerned about the compatibility of both complicated metallizations. + +# Acknowledgements + +The collaboration and hard work of H.S. Chen, D.D. Bacon, F. Baiocchi, E. Lane and C. Doherty, in realizing and carrying out this study, are greatly appreciated. + +# References + +1 K. Das, V. Venkatesan, K. Miyata, D.L. Dreifus and J.T. Glass, in Y. Tzeng, M. Yoshikawa, M. Murakawa and A. Feldman (eds.), Applications of Diamond Films and Related Materials, Elsevier, Amsterdam, 1991, pp. 301-308. +2 A.T. Collins, Semicond. Sci. Technol., 4 (1989) 605. +3 J.E. Graebner, S. Jin, J.A. Herb and C.F. Gardinier, Appl. Phys. Lett., 60 (1992) 1576. +4 J.E. Graebner, S. Jin, G.W. Kammlott, J.A. Herb and C.F. Gardinier, Nature, 359 (1992) 401. +5 Wall Street Journal, October 13 (1992) +6 Diamond and Structural Carbon News, October (1992). +7 C.C. Lee, C.Y. Wang and G.S. Matijasevic, IEEE Trans. Comp. Hybrids, Manuf. Technol., 14 (1991) 407. +8 J.H. Lau and D.W. Rice, Solid-State Technol., 28 (1985) 91. +9 S. Knecht and L.R. Fox, IEEE Trans. Comp. Hybrids, Manuf. Technol., Chmt-9 (1986) 423. +10 O. Wada and T. Kumai, Jpn. J. Appl. Phys., 40 (1991) L1056. +11 G.S. Matijasevic and C.C. Lee, Proc. IEEE Int. Reliability Physics Symp., Phoenix, AZ, November, 1989. +12 O. Wada and O. Ueda, in A. Katz, S.P. Muraka and A. Appelbaum (eds.), Advanced Metallization in Microelectronics, Mat. Res. Soc. Symp. Proc. 181, MRS, Pittsburgh, 1990, pp. 273-280. + +13 W.G. Bader, Welding Res. Suppl., 48 (1969) 551. +14 O. Wada and T. Kumai, Appl. Phys. Lett., 58 (1991) 908. +15 C.Y. Wang and C.C. Lee, IEEE Trans. Comp. Hybrids, Manuf. Technol., 14 (1991) 874. +16 H. Okamota and T.B. Massalski, in T.B. Massalski (ed.), Binary Alloy Phase Diagrams, AMS, Ohio, 1991, p. 403. +17 P.J. Kay and C.A. Mackay, Trans. Inst. Met. Finish., 54 (1976) 68. +18 R.J. Klein Wassink in Soldering in Electronics, Electrochemical Publications, Ayr, UK, 1989, p. 179. +19 P. Nash and A. Nash, in T.B. Massalski (ed.), *Binary Alloy Phase Diagrams*, AMS, Ohio, 1991, p. 2864. +20 W.J. Tomlinson and H.G. Rhodes, J. Mater. Sci., 22 (1987) 1769. +21 A.C. Harman, Proc. Technical Programme Interepcon 1978, Brighton, UK, pp. 42-49. +22 P.J. Kay and C.A. Mackay, Trans. Inst. Met. Finish., 57 (1979) 169. + +23 N.G. Koopman, T.C. Reiley and P.A. Totta, in R.R. Tummala and E.J. Rymaszewski (eds.), Microelectronics Packaging Handbook, Van Nostrand Reinhold, New York, 1989, p. 395. +24 K. Mizuishi and T. Mori, IEEE Trans. Component Hybrids Manuf. Technol., Chmt-11 (1988) 481. +25 A. Katz, F. Baiocchi, E. Lane, C.H. Lee, C. Hall, J. Doting, C. Gritsbach and K. Harris, J. Appl. Phys., to be published. +26 US Pat. 005 197 654 (1993) to A. Katz, C.-H. Lee, K.L. Tai and Y.-M. Wang. +27 A. Katz, K.-W. Wang, F.A. Baiocchi, W.C. Dautremont-Smith, E. Lane, H.S. Luftman, R.R. Varma and H. Curnan, J. Appl. Phys., 72 (1992) 3808. +28 R. Darshmukh, E. Laskowski and J. Shmulovich, private communication. +29 S. Nakahara, R.J. McCoy, L. Buene and J. Vanderberg, Thin Solid Films, 84 (1981) 185. +30 US Pat. 5 234 153 (1993), to A. Katz, D.D. Bacon, C.-H. Lee, K.L. Tai and Y.-M. 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b/金刚石散热项目/项目前期工作_MD/02_金属化金刚石载板与器件封装/8_Thermal stability of metallized CVD diamond./8_Thermal stability of metallized CVD diamond..md new file mode 100644 index 0000000..4bb6a99 --- /dev/null +++ b/金刚石散热项目/项目前期工作_MD/02_金属化金刚石载板与器件封装/8_Thermal stability of metallized CVD diamond./8_Thermal stability of metallized CVD diamond..md @@ -0,0 +1,132 @@ +# Thermal stability of metallized CVD diamond + +C.D. Iacovangelo + +Chemical Process Technology Laboratory, GE Corporate Research and Development Center, Schenectady, NY 12301, USA + +Received 10 July 1995; accepted 8 November 1995 + +# Abstract + +Four metallization systems, WTi/Au, Ti/Pt/Au, Mo/Au and Nb/Au, have been examined for thermal reliability, solderability, wire bonding and adhesion on chemically vapor deposited diamond (CVDD). Thermal stability up to $600^{\circ}\mathrm{C}$ was determined by conductivity measurements and X-ray photoelectron spectroscopy depth profiles. WTi/Au was found to be the most thermally stable, exhibiting constant resistivity and little interdiffusion after 4 h at $500^{\circ}\mathrm{C}$ . Nb/Au was stable to approximately $450^{\circ}\mathrm{C}$ , where marked diffusion of Nb into the gold caused a rapid rise in resistivity and loss of adhesion. Ti/Pt/Au exhibited slow diffusion of Pt into the Au and a modest rise in resistivity as low as $350^{\circ}\mathrm{C}$ . Only Mo/Au showed significant interdiffusion between the Mo bond coat and CVDD at such low temperatures. This interdiffusion was markedly enhanced by the presence of oxygen. Each of the four metallization systems initially demonstrated excellent adhesion to CVDD, although all eventually lost adhesion due to conversion of the transition metal carbide bond layer to an oxide. + +Keywords: Adhesion; Coatings; Diamond; Diffusion + +# 1. Introduction + +Diamond, being the hardest material, has long enjoyed the position of the ultimate material for most cutting and grinding applications. Diamond, however, has several properties which make it an ideal material for electronic applications, the most salient being thermal conductivity (12-18 W cm $^{-1}$ K $^{-1}$ ), and dielectric constant (5.7). Currently, chemically vapor deposited diamond (CVDD) has experienced limited market entry primarily as heat sinks for laser diodes [1]. Several investigators are exploring large-area (10 cm× 10 cm) multi-chip modules using CVDD substrates (DSMCMs) to remove the excessive heat that is generated in high-density MCMs [2,3]. In this application, the low dielectric constant, which allows close packing of traces and high speed, is an advantage over AlN or BeO. Transparent CVDD films can make ideal windows for a host of applications ranging from X-rays to microwaves. Diamond can even be doped to provide potential semiconductor properties for passive and active electronic devices that could operate at elevated temperatures. Improvements in process reproducibility and deposition rates over large areas, have brought the cost of CVDD in reach of electronic applications. Each of these applications requires cost effective, reliable methods for metallizing diamond with minimal degradation of thermal and electrical performance. + +It has been shown that transition metals can be used as bond layers on CVDD with gold as a top coat for wire bonding and contacts [4,5]. While several authors have recommended these metallizations as ohmic contacts for elevated temperature diamond electronics [6,7], little data exists in the literature on the thermal stability of these metallization systems on CVDD, or their reliability for wire bonding and solder die attach. Katz et al. have shown that the state of the art Ti/Pt/Au metallization system can be used to die attach InP chips to CVDD with thin film solder as deposited but no information was given as to thermal reliability [8]. Katz et al. have also shown that Ti/Ni/Au can overcome the problem of Pt reaction with the Sn in AuSn solders but again no information was given as to thermal stability c' these systems. In many applications such as MCMs, the diamond substrate must go through several thermal cycles often up to an hour at $350^{\circ}\mathrm{C}$ , with repeated exposure when repairs are needed. This puts a significant thermal stability requirement on the metallization system of choice and can dramatically limit the application for many of the standard metallization systems such as Ni/Au [9]. It is critical, therefore, that metallization systems on CVDD be developed that can withstands and long durations in air at elevated temperatures and remain solderable, well adherent, and wire bondable. We will show data for three metallization systems that demonstrate this reliability. + +0040-6090/96/$15.00 © 1996 Elsevier Science S.A. All rights reserved +SSD/0040-6090(95)08236-0 + +# 2. Experimental + +All experiments were conducted on square CVDD plates approximately $2.54\mathrm{cm}$ on a side and approximately 0.025 cm thick. The CVDD was mechanically polished with diamond grit to approximately $300\mathring{\mathrm{A}}R_{\mathrm{a}}$ . The diamond was cleaned in boiling acids, rinsed in deionized water, and dried in methanol and propanol prior to deposition. The CVDD was sputter etched for 6 min at $100\mathrm{W}$ prior to deposition. All metals were d.c. magnetron sputtered in a Plasma Science ARC unit at $0.5 - 2.2\mathrm{Pa}$ argon. The background pressure was typically $2\times 10^{-7}\mathrm{Pa}$ prior to sputtering. The four transition metal bond coats Ti, WTi, Mo, and Nb were deposited at $5\mathrm{Wcm}^{-2}$ . Platinum and gold were deposited at $2.5\mathrm{Wcm}^2$ . The target to substrate distance was approximately $7.5\mathrm{cm}$ which resulted in deposition rates of 45, 65, 48, 60, 75 and $90\mathrm{nmmin}^{-1}$ for Ti, WTi, Mo, Nb, Pt and Au respectively. The WTi was sputtered from a 90:10 alloy target and the composition of the deposit was analyzed with AA and ICP. As shown in Fig. 1, the deposits were $8\%$ Ti and independent of sputtering power or pressure. The stress of the deposit, however, was highly pressure dependent. Fig. 1 shows the intrinsic stress in the deposit as a function of Ar sputtering pressure. The deposits used in these studies were deposited at $1.2\mathrm{Pa}$ with a slight compressive stress. The conductivity of the bond coats, shown in Fig. 2, was independent of pressure for the WTi coating. The resistivity of Nb and Ti, however, increased with pressure, presumably due to incorporation of gas into the deposit. For this reason, the deposits were sputtered at lower gas pressures. Unless otherwise stated in the text, all heat treatments were carried out in air. Temperatures were controlled to $\pm 2^{\circ}\mathrm{C}$ using a Love® controller with a type-K thermocouple. Adhesion was measured with a Sebastian tensile adhesion tester using epoxy coated pull pins with a stress limit of $69\mathrm{MPa}$ . The metallization was patterned into $0.25\mathrm{cm}$ wide by $2.54\mathrm{cm}$ long strips for conductivity measurements. Measurements were made using a Keithley Model 224 current source, Keithley Model 19A multimeter, and a Dumas Model C45 4 point probe. Values given are the average of 3-6 readings for each metallization system. X-ray photoelectron spectroscopy (XPS) surface analysis was obtained using a Physical Electronics Model PHI-45. Depth profiles were obtained by sputter etch + +![](images/9cdfecc389f64e6ef3fc2588cda13c8ba7e82a403db7ea6f92fe4f49c55cb518.jpg) +Fig. 1. Composition and stress of WTi deposits on CVDD as a function of argon sputtering pressure. + +![](images/0e7cb370e8c4abc601427a335929b395780714dc07af5c68794b24c71bd794ef.jpg) +Fig. 2. Resistivity of sputtered bond coats on CVDD as a function of argon sputtering pressure. + +ing at approximately $5\mathrm{nm}\min^{-1}$ using $2.8\mathrm{kV}$ argon ions. Wire bonding was performed with $3.3\times 10^{-3}$ cm gold wire using a K&S model 1926 wire bonding station and a Micro Tester Model BT22 pull tester. + +# 3. Results and discussion + +Most metallization systems on CVDD depend upon transition metals to bond to diamond through the formation of transition metal carbides at the metal/CVDD interface. Earlier [10], we demonstrated excellent adhesion of Cr, Nb, Ti, W, and WTi as bond coats deposited by a variety of techniques including sputtering, pack CVD [11], low-pressure CVD (LPCVD) [12], and electroless plating. Electronic applications require the metallization to undergo a variety of thermal excursions from a few hundred degrees to as high as $500^{\circ}\mathrm{C}$ (e.g. die and seal ring attachment, lid sealing, and lead attach) and remain conductive and oxide free so that reliable soldering and wire bonding can be performed. In this paper we examined the thermal reliability of four metallization systems on CVDD: Ti/Pt/Au 100/200/500 nm, WTi (90/10)/Au 100/500 nm, Mo/Au 100/500 nm, and Nb/Au 100/500 nm. The Ti/Pt/Au was chosen because of its copious use in the CVDD community as an industry standard. This metallization, however, is expensive, difficult to pattern because of the Pt, and interacts with AuSn solder in a detrimental manner [8]. Niobium is unique among the transition metals in that the carbide is readily dissolved in aqueous solutions. Nb/Au was, therefore, selected for the ease of patterning the NbC bond layer by wet etching. Nb has been used in a few applications as a barrier between Au and Al, however, the authors did not show data on the transport of Nb into the gold at these temperatures [13]. Like Nb, WTi has been used extensively as a barrier to Al transport [14] for ceramic packaging. Recently, Naseem et al. [5] showed thermal reliability of WTi/Au on CVDD up to $450^{\circ}\mathrm{C}$ . Mo/Au has been used extensively for high-temperature CVDD applications as an active device. interesting though, most authors do not report any long term stability data even though the intended use is a high temperature application. Rosen et al. [15] have shown that excellent ohmic contacts can be maintained up to $130\mathrm{h}$ in low $P_{\mathrm{O_2}}$ argon atmospheres. + +The stability of these metallizations to thermal processing was examined by monitoring the conductivity of patterned metal stripes after heat treating for up to $4\mathrm{h}$ from 300 to 600 ${}^{\circ}\mathbf{C}$ in air. XPS depth profiles were performed, generally after $4\mathrm{h}$ at each temperature. Plotted in Fig. 3 is the resistivity of WTi/Au, Ti/Pt/Au and Nb/Au metallizations after heat treating from 300 to $600^{\circ}\mathbf{C}$ . As shown, all three metallizations demonstrated good stability to processing at $300^{\circ}\mathbf{C}$ for $4\mathrm{h}$ with no increase in resistivity. XPS depth profiles showed no grain boundary diffusion of the bond coats into the gold and pull tests showed no deterioration in adhesion; all coating exceeded $69\mathrm{MPa}$ . At $400^{\circ}\mathbf{C}$ , the Ti/Pt/Au standard showed a marked rise in resistivity compared with WTi/Au and Nb/Au, which showed no increase in resistivity. In fact, these metallizations demonstrated a slight improvement in conductivity at $400^{\circ}\mathbf{C}$ . The XPS depth profiles of Fig. 4 show the rise in resistivity to be due to Pt diffusion into the gold. No Nb diffusion into the gold was apparent after this thermal excursion. Likewise, the WTi sample showed no detectable transport of Ti out of the alloy. There was a slightly higher oxygen content but this was also seen in the as-deposited sample. By comparison, a pure Ti layer would have shown significant diffusion through the gold, resulting in a several hundred angstrom thick titanium oxide coating [16]. Tungsten which is the slower diffusing metal through gold, was not observed. + +![](images/bdbfb803208fe8df5c7be7780c6bcc2888262ae0f2daef201ac41abf44887adf.jpg) + +![](images/1bb40ae59e3ffc4301f4f7e7c36b5224976589cac5a6b4b5f765bd88511bef82.jpg) + +![](images/3aa48b076012f9979f629c0365d51a6692c918017f5e38f5633b9d894a073108.jpg) +Fig. 3. Electrical resistivity of sputtered coatings on CVDD after heat treating in air from 300 to $600^{\circ}\mathrm{C}$ . + +![](images/54deb301d2002cbe21695291c00f1a9a520b7ef7c65aca75a3c80d8b090d84b7.jpg) + +![](images/92ea6d64cd09241a2783594bc2d1802ef37e74d9950198f8c8f2e269246b744a.jpg) + +![](images/e63202a30442dfde4995a63f53639d31680db9d3da01680dc8b920ae1c616b8e.jpg) +Fig. 4. XPS depth profile of sputtered coatings on CVDD after heat treating in air for $4\mathrm{h}$ at $400^{\circ}\mathrm{C}$ . + +At $500^{\circ}\mathrm{C}$ , both the Ti/Pt/Au and Nb/Au samples had a catastrophic rise in resistivity due to the transport of Pt and Nb into the gold; as shown in Fig. 5. Both samples also exhibited marked oxidation of the diffused metal. This effect was dramatic for the Nb sample which exhibited oxidation through the entire sample sufficient to decarburize the NbC bond layer and cause delamination of the metallization. As shown, the WTi/Au sample remained stable through 4 h at $500^{\circ}\mathrm{C}$ with no interdiffusion or degradation in adhesion. + +These results do not agree with those of Naseem et al. [5], who found significantly more interdiffusion of the WTi into the Au at these temperatures. The exact cause of this discrepancy is unknown. One possibility is the activity of Ti in the alloy. The marked increase in stability of the alloy versus a pure Ti layer is due to the decreased activity of Ti in the deposited alloy. It should be noted, however, that the Ti activity in the deposit can be markedly different from the target depending upon the sputtering conditions, distance from the target, and geographical position relative to the center of the target. This is due to the large difference in mass between the W and Ti ions. The samples used in this study were placed directly under the center of a round target with no motion of the substrate so that no significant variation in W:Ti occurred within the 2.5 cm square. Since no data was given for the composition of the deposit in Ref. [5], it is possible that the deposit was richer in Ti, which would + +![](images/b6cceb2b916e7d3698a415704ff54283d945a7a955a46ed29331d95c7b6c9223.jpg) + +![](images/049d2cd4edda411c407a01f98a2480f1bbc8603770f91d88b43b00586281828f.jpg) + +![](images/99f32263ad5fcdfdb55cf0aa127dcbb672d3a20476a19914dc95ac1c33b5a32e.jpg) +Fig. 5. XPS depth profile of sputtered coatings on CVDD after heat treating in air for $4\mathrm{h}$ at $500^{\circ}\mathrm{C}$ . + +decrease its thermal stability but enhance its adhesion to CVDD. + +The WTi/Au sample remained stable through 1 h at 600 °C; as shown in Fig. 6. After 4 h, however, the resistivity doubled because of excessive oxidation. Like the Nb sample, this oxidation penetrated through to the carbide bond layer causing delamination of the metallization. + +The results for Mo/Au metallization of CVDD were dramatically different than the other three transition metal bond coats. Shown in Fig. 7 are XPS depth profiles from the Au surface through to the CVDD on as-deposited coatings and after $4\mathrm{h}$ at 450 and $550^{\circ}\mathrm{C}$ . The disappearance of the Mo interface is quite striking. XPS analysis also shows a significant increase in oxidation of the Mo with increasing temperature. The previous samples demonstrated classic grain boundary diffusion of the transition metal through the gold with little measurable diffusion into the CVDD. Mo, in contrast, showed massive migration into both the Au and the CVDD. We have not observed this phenomena with any other transition metal. + +This test was repeated with a modification that the entire CVDD was coated with Mo and only the Au was patterned + +![](images/e4fc2aca65a608233d557afc98bf6d5af5da3dfa0e7945b53c23661e5bfbb128.jpg) +Fig. 6. XPS depth profile of sputtered Au/WTi coatings on CVDD after heat treating in air for $4\mathrm{h}$ at $600^{\circ}\mathrm{C}$ . + +![](images/45939e83ab19017ea70192c742f550eca2cb2fd2fa839e7af46b74fc4662b2a6.jpg) + +![](images/9c84b2ddd0444fec9f41a15b8af5a546ec60fc212e3ec8aca57f90d4c62a4022.jpg) + +![](images/76d4bbb228cb27afd8641dce746ed9a162180a6966f9390cfbdecb8d46bfc134.jpg) +Fig. 7. XPS depth profile of Mo/Au coatings on CVDD: top, as sputtered; middle, after heat treating in air for $4\mathrm{h}$ at $450^{\circ}\mathrm{C}$ ; bottom, after heat treating in air for $4\mathrm{h}$ at $550^{\circ}\mathrm{C}$ . + +for conductivity measurements. This insured that the Mo layer was unperturbed by the Au deposition process. XPS depth profiles of an area where the gold was removed leaving only the Mo showed that the Mo had migrated deep into the CVDD. The test was repeated in a $4\%$ hydrogen/argon environment to examine the effect of oxygen on molybdenum oxidation and interdiffusion. Fig. 8 shows that very little Mo diffusion into either the gold or CVDD occurred without the presence of oxygen. + +![](images/5858e24c014d058223add5f6043bcb9914a931a91f076763f2272d5643f24f06.jpg) +Fig. 8. XPS depth profile of Mo/Au coatings on CVDD after heat treating in $5\% \mathrm{H}_2 / \mathrm{Ar}$ for $4\mathrm{h}$ at $550^{\circ}\mathrm{C}$ . + +![](images/6ea0979a612d598cc225b7b1763e076b73e6b5d974f9c226db9cecfb55f41336.jpg) +Fig. 9. Electrical resistivity of Au/Mo/CVDD metallization after heat treating in air and $4\%$ $\mathrm{H}_{2} / \mathrm{Ar}$ . + +The resistivity change with time at temperature for samples treated in air and hydrogen/argon are shown in Fig. 9. In the reducing environment the resistivity was unchanged through $4\mathrm{h}$ at $550^{\circ}\mathrm{C}$ , as observed by Rosen et al. In air, however, the resistivity rose dramatically at temperatures as low as $350^{\circ}\mathrm{C}$ , corresponding to the interdiffusion shown in Figs. 7 and 8. These results demonstrate the critically of any processing of the diamond in air prior to hermetic sealing and the need for extremely low $P_{\mathrm{O2}}$ hermetic seals when this metallization is used. + +Wire bond tests after $4\mathrm{h}$ of heat treating on $\mathrm{WTi / Au}$ and Ti/Pt/Au samples using $3.3\times 10^{-3}$ cm Au wire are shown in Fig. 10. The WTi/Au samples had a mean of 12.9 and 13.3 g at $400^{\circ}\mathbf{C}$ and $500^{\circ}\mathbf{C}$ respectively, with a standard deviation of $0.6\mathrm{g}$ . Most of the pulls resulted in wire failure rather than separation from the substrate. The Ti/Pt/Au sample showed a decrease in mean pull strength from 14.2 to $12.6\mathrm{g}$ , and a slight increase in standard deviation to $1.5\mathrm{g}$ from $400^{\circ}\mathbf{C}$ to $500^{\circ}\mathbf{C}$ . Several of the breaks occurred at the wire/metallization interface. Tensile pull tests showed no decrease in adhesion of either metallization to the CVDD after $4\mathrm{h}$ at $500^{\circ}\mathbf{C}$ . No observable increase in carbide formation was observed for any of the four metallizations. + +# 4. Conclusions + +Three of the four sputtered coating systems examined provided thermally reliable thin film metallization on CVDD: + +![](images/25da677154bbed5dca49820ea41c758a7f2cff3a70cb726083353fc788e03342.jpg) +Fig. 10. Pull strength of $3.3 \times 10^{-3}$ cm gold wires after heat treating for $4\mathrm{h}$ at 400 and $500^{\circ}\mathrm{C}$ . + +Ti/Pt/Au, WTi(90/10)/Au, and Nb/Au. All three demonstrated excellent adhesion to CVDD and stability to elevated temperature processing in air up to $300^{\circ}\mathrm{C}$ for 4 h. The Nb/Au metallization was more stable than the Ti/Pt/Au industry standard at $400^{\circ}\mathrm{C}$ but degraded rapidly at 500 ${}^{\circ}\mathrm{C}$ due to interdiffusion and oxidation of the Nb. The WTi/Au metallization had the best performance, remaining stable at $550^{\circ}\mathrm{C}$ for up to 4 h. After 1 h at $600^{\circ}\mathrm{C}$ , this metallization also failed because of Ti transport out of the WTi. Both the Ti in the gold and the TiC bonding the metallization to the CVDD oxidized, causing poor performance and delamination from the CVDD. All three metallizations provided adequate thermal stability for most electronic applications. The optimum metallization will depend on other factors such as solder compatibility, cost, ease of patterning, oxidation resistance, etc. + +The Mo/Au metallization system behavior in air was unique among the systems studied. While the performance in reducing environments was extremely stable up to $600^{\circ}\mathbf{C}$ , the performance degrade rapidly in air at temperatures as low as $350^{\circ}\mathbf{C}$ . Oxidation of the Mo resulted in an excessive rise in resistivity and massive diffusion of the Mo into both the Au and CVDD layers. + +# Acknowledgements + +The author wishes to acknowledge the assistance of J. Chera for XPS analysis, K. Borst, L. Gruenke, and D. Skelly for sputtering, and E. Jerabek and P. DiConza for their constant help and fruitful discussions. The author also wishes to + +acknowledge J. Pagois for manuscript preparation, and GE Superabrasives for supplying CVDD test samples. The author gratefully acknowledges NRL for administering and ARPA for the financial support of this work under contract #N00014-91-C-2167. + +# References + +[1] AT&T, The world's first commercial application of a novel synthetic diamond material combined with an advanced metallization bonding technique, Wall Street J., B9 (1992). +[2] General Electric, Development of free standing diamond manufacturing technology for thermal management substrates, ARPA #N00014-91-C-2167. +[3] R.C. Eden, Applicability of diamond substrates to multi-chip modules, Int. Soc. Microelectronics Proc. 1, INSIM, Reston, VA, 1991, pp. 363-367. +[4] Jiang Xiang-Liu et al., Diamond Optics V, SPIE, 1759 (1992) 145-153. + +[5] N.A. Naseem et al., Int. J. Microcircuits Electron. Packaging, 16(4) (1993) 257-268. +[6] J. Nakanishi et al., J. Appl. Phys., 76(4) (1994) 2293-2298. +[7] K.L. Moazed, J.R. Zeildner and M.J. Taylor, Mater. Res. Soc. Symp. Proc., 162 (1990) 347-352. +[8] A. Katz, C.H. Lee and K.L. Tai, Mater. Chem. Phys., 37 (1994) 303-328. +[9] C.D. Iacovangelo, New autocatalytic gold bath and diffusion barrier coatings, AESF SUR/FIN'92, Inter. Tech. Conf. Proc., Atlanta, GA, June 22-25, AESF, Orlando, FL, Vol. 1, 1992, p. 51. +[10] C.D. Iacovangelo and E.C. Jerabek, Metallizing CVD diamond for electronic substrates, Int. Soc. Microelectronics Proc., Dallas, TX, ISNIM, Reston, VA, 1993, pp. 132-137. +[11] D.S. Park, US Patent No. 4664991, June 15, 1967. +[12] C.D. Iacovangeo, E.C. Jerabek and R.H. Wilson, US Patent No. 5328715, June 12, 1994. +[13] B.L. Doyle et al., J. Appl. Phys., 53(9) (1982) 6186-6190. +[14] J.M. Oparowski, R.D. Sisson and R.R. Biederman, Thin Solid Films, 153 (1987) 313-328. +[15] M. Rosen, C.A. Hewett, K.L. Moazed and J.R. Zeidler, J. Electrochem. Soc., 139(7) (1992) 2001-2004. +[16] R.S. 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submounts produced by microwave plasma chemical vapour deposition/9_Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition.md new file mode 100644 index 0000000..32284c8 --- /dev/null +++ b/金刚石散热项目/项目前期工作_MD/02_金属化金刚石载板与器件封装/9_Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition/9_Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition.md @@ -0,0 +1,76 @@ +Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition + +This content has been downloaded from IOPscience. Please scroll down to see the full text. + +2012 Quantum Electron. 42 959 + +(http://iopscience.iop.org/1063-7818/42/11/L02) + +View the table of contents for this issue, or go to the journal homepage for more + +Download details: + +IP Address: 137.99.26.43 + +This content was downloaded on 07/01/2014 at 14:38 + +Please note that terms and conditions apply. + +PACS numbers: 42.55.Px; 42.60.Lh; 81.05.ug + +DOI: 10.1070/QE2012v042n11ABEH015042 + +# Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition + +E.E. Ashkinazi, V.V. Bezotosnyi, V.Yu. Bondarev, V.I. Kovalenko, V.I. Konov, O.N. Krokhin, V.A. Oleshchenko, V.F. Pevtsov, Yu.M. Popov, A.F. Popovich, V.G. Ral’chenko, E.A. Cheshev + +Abstract.  We have designed and fabricated submounts from synthetic diamond grown by microwave plasma chemical vapour deposition and developed an economical process for metallising such submounts. Laser diode chips having an 808-nm emission wavelength, 3-mm-long cavity and 130-mm-wide stripe contact were mounted on copper heat sinks with the use of diamond submounts differing in quality. The devices were tested for more than 150 h in continuous mode at an output power of 8 W on diamond with a thermal conductivity of 700 $\bf { W } \delta m ^ { - 1 } \delta K ^ { - 1 } ,$ , and no changes in their output power were detected. On diamond with a thermal conductivity of 1600 W $\mathbf { m } ^ { - 1 } \mathbf { K } ^ { - 1 }$ , stable cw operation for 24 h at an output power of 12 W was demonstrated. + +Keywords: laser diode, diamond submount. + +# 1. Introduction + +The purpose of this work was to increase the output power, brightness, efficiency and operating life of high-power laser diodes (LDs). Solving this set of problems, one will potentially be able to bring the brightness of systems combining the outputs of diode lasers closer to the level of state-of-the-art $\mathrm { C O } _ { 2 }$ lasers and diode-pumped solid-state and fibre lasers and move to large-scale materials processing using the LD output. + +An integral and key part of this issue is the ability to raise the effectiveness of heat removal from the active region of LDs, which has been the focus of many studies, including those reported in Refs [1 – 3]. The removal of extremely high heat flux densities (above 3 – 5 kW cm–2) at minimum possible temperature differences between the active region and heat sink (due to the strong temperature dependence of the emission characteristics of the semiconductor medium) is the most important fundamental limitation to the most complete realisation of the energy potential of the active element of a highpower LD. + +# 2. Synthetic diamond and submounts for laser diode mounting + +The record high thermal conductivity of diamond among all known materials, 2000– 2400 W $\mathrm { m ^ { - 1 } ~ K ^ { - 1 } }$ at room temperature, is the most attractive of its parameters for us. It should be emphasised that diamond as a thermal conductor has appreciable inherent reserves: the thermal conductivity of isotopically pure diamond may reach 3300 W $\mathrm { m } ^ { - 1 } \mathrm { K } ^ { - 1 }$ (natural crystals contain $1 . 1 \% ^ { 1 3 } \mathrm { C } )$ . Note that the thermal conductivity of copper, which is currently widely used as a heat sink material, is about $3 7 0 \mathrm { \ : W \ : m ^ { - 1 } \ : K ^ { - 1 } }$ . + +The use of diamond as a heat dissipation material for mounting high-power laser chips required solving a number of serious problems: growth of diamond wafers more than 300 mm in thickness; wafer cutting, grinding and polishing with necessary parameters; application of adhesive and solder layers for double-sided metallisation; diamond mounting on a heat sink base and laser chip mounting on the diamond; wiring of the lower (positive) electrode because diamond is an insulator; and thermal expansion matching (at 300 K, the thermal expansion coefficient of diamond is $1 \times 1 0 ^ { - 6 } \mathrm { K ^ { - 1 } }$ 1, that of copper is $1 6 . 7 \times 1 0 ^ { - 6 } ~ \mathrm { K ^ { - 1 } }$ , and that of semiconductor heterostructures is about $6 \times 1 0 ^ { - 6 } \mathrm { K ^ { - 1 } } )$ ). + +Diamond films were grown by microwave plasma chemical vapour deposition (CVD), a relatively economical process, on silicon substrates preseeded with ultradispersed diamond under sonication. + +Diamond wafers for injection lasers were prepared through microwave plasma CVD of diamond on silicon substrates, separation of the diamond wafers from the substrates by chemical etching in an acid, laser cutting of the wafers into submounts and polishing. The process involved the decomposition of a hydrocarbon mixture (methane + hydrogen) and carbon deposition on a preheated substrate. + +# 3. Results + +LD chips for an 808-nm wavelength had a 3-mm-long cavity and 130-mm-wide stripe contact. The diamond submounts were $2 \times 4 \times 0 . 3$ mm in dimensions. + +Figure 1 shows power – current characteristics of a cw LD mounted directly on a copper C-mount (LDC) and on an identical C-mount with the use of a diamond submount (LDD). Our life test statistics indicate that the guaranteed maximum output power of the LDC is 7 W. According to preliminary test results, the LDD provided a stable output of 12 W for 24 h. After producing a statistically relevant number of devices and evaluating all their parameters, we will con- + +![](images/5f8d700930cee4e38e479e8d9e750cea6312fab0b92786b16551cbc0c0fbd6c3.jpg) +Figure 1. Power – current characteristics of the LDD and LDC. + +tinue testing, determine the maximum power of the LD design in question and find out its limitations. + +Figure 2 shows the peak emission wavelength as a function of pump current for the spectral envelopes of the LDC and LDD. It is seen that, at a current of 8 A, the peak in the spectrum of the LDD is shifted to shorter wavelengths by almost 2 nm relative to the spectrum of the LDC, which points to a more effective heat removal from the chip through the diamond submount because it acts as a ‘negative thermal lens’. The heat from the 130-mm-wide laser stripe contact is more effectively spread over the entire width (2 mm) of the diamond submount. + +![](images/118745ac717978771497318e174a3b7058fcaf7dee88c0871c189053ca5b6518.jpg) +Figure 2. Peak emission wavelength as a function of pump current for the envelopes of the emission spectra of the LDD and LDC. + +Comparison of the full width at half maximum (FWHM) as a function of pump current for the spectral envelopes of the LDC and LDD (Fig. 3) demonstrates the advantageous effect of the lower active region temperature in the LDD at high currents. + +Thus, mounting an LD chip on a copper C-mount with the use of a synthetic diamond submount we achieved stable cw operation of a 12-W LD at 808 nm for 24 h. Output power data and emission spectra demonstrate that diamond submounts improve the effectiveness of heat flux removal from high-power LDs. + +Acknowledgements.  This work was supported by the Russian Foundation for Basic Research (Grant Nos 11-02-00922-a and 11-02-90025-Bel_a). + +![](images/d116ec6cc090d502630e127e848acd3937007aa2dabd4bd9ca74b87c271500cf.jpg) +Figure 3. FWHM as a function of pump current for the spectral envelopes of the LDD and LDC. + +# References + +1. Seurin J.F., Xu G., Khalfin V., Miglo A., Wynn J.D., Pradhan P., Ghosh C.L., D’Asaro L.A. Proc. SPIE Int. Soc. Opt. Eng., 7229, 722903 (2009). +2. Parashchuk V.V., Ryabtsev G.I., Belyaeva A.K., Bez’yazychnaya T.V., Baranov V.V., Telesh E.V., Vu Doan Mien, Vu Van Luc, Pham Van Truong. Kvantovaya Elektron., 40, 301 (2010) [Quantum Electron., 40, 301 (2010)]. +3. Ashkinazi E.E., Bezotosnyi V.V., Bondarev V.Yu., Kovalenko V.I., Krokhin O.N., Oleshchenko V.A., Pevtsov V.F., Popov Yu.M., Cheshev E.A. Trudy konf. ‘Poluprovodnikovye lazery i sistemy’ (Book of Papers, 8th Belorussian – Russian Workshop ‘Semiconductor Lasers and Systems’) (Minsk, 2011) p.  29. \ No newline at end of file diff --git a/金刚石散热项目/项目前期工作_MD/02_金属化金刚石载板与器件封装/9_Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition/images/118745ac717978771497318e174a3b7058fcaf7dee88c0871c189053ca5b6518.jpg b/金刚石散热项目/项目前期工作_MD/02_金属化金刚石载板与器件封装/9_Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition/images/118745ac717978771497318e174a3b7058fcaf7dee88c0871c189053ca5b6518.jpg new file mode 100644 index 0000000..449d740 Binary files /dev/null and b/金刚石散热项目/项目前期工作_MD/02_金属化金刚石载板与器件封装/9_Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition/images/118745ac717978771497318e174a3b7058fcaf7dee88c0871c189053ca5b6518.jpg differ diff --git a/金刚石散热项目/项目前期工作_MD/02_金属化金刚石载板与器件封装/9_Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition/images/5f8d700930cee4e38e479e8d9e750cea6312fab0b92786b16551cbc0c0fbd6c3.jpg b/金刚石散热项目/项目前期工作_MD/02_金属化金刚石载板与器件封装/9_Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition/images/5f8d700930cee4e38e479e8d9e750cea6312fab0b92786b16551cbc0c0fbd6c3.jpg new file mode 100644 index 0000000..2648755 Binary files /dev/null and b/金刚石散热项目/项目前期工作_MD/02_金属化金刚石载板与器件封装/9_Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition/images/5f8d700930cee4e38e479e8d9e750cea6312fab0b92786b16551cbc0c0fbd6c3.jpg differ diff --git a/金刚石散热项目/项目前期工作_MD/02_金属化金刚石载板与器件封装/9_Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition/images/d116ec6cc090d502630e127e848acd3937007aa2dabd4bd9ca74b87c271500cf.jpg b/金刚石散热项目/项目前期工作_MD/02_金属化金刚石载板与器件封装/9_Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition/images/d116ec6cc090d502630e127e848acd3937007aa2dabd4bd9ca74b87c271500cf.jpg new file mode 100644 index 0000000..eb0ceae Binary files /dev/null and b/金刚石散热项目/项目前期工作_MD/02_金属化金刚石载板与器件封装/9_Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition/images/d116ec6cc090d502630e127e848acd3937007aa2dabd4bd9ca74b87c271500cf.jpg differ diff --git a/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/10_High thermal conductive copper-diamond composites:state of the art/10_High thermal conductive copper-diamond composites:state of the art.md b/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/10_High thermal conductive copper-diamond composites:state of the art/10_High thermal conductive copper-diamond composites:state of the art.md new file mode 100644 index 0000000..9a477b7 --- /dev/null +++ b/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/10_High thermal conductive copper-diamond composites:state of the art/10_High thermal conductive copper-diamond composites:state of the art.md @@ -0,0 +1,790 @@ +# High thermal conductive copper/diamond composites: state of the art + +S. Q. Jia1 and F. Yang1,* + +1 Waikato Centre for Advanced Materials and Manufacturing, School of Engineering, University of Waikato, Hamilton 3240, New Zealand + +Received: 22 July 2020 + +Accepted: 24 September 2020 + +Published online: + +20 October 2020 + +- The Author(s) 2020 + +# ABSTRACT + +Copper/diamond composites have drawn lots of attention in the last few decades, due to its potential high thermal conductivity and promising applications in high-power electronic devices. However, the bottlenecks for their practical application are high manufacturing/machining cost and uncontrollable thermal performance affected by the interface characteristics, and the interface thermal conductance mechanisms are still unclear. In this paper, we reviewed the recent research works carried out on this topic, and this primarily includes (1) evaluating the commonly acknowledged principles for acquiring high thermal conductivity of copper/diamond composites that are produced by different processing methods; (2) addressing the factors that influence the thermal conductivity of copper/diamond composites; and (3) elaborating the interface thermal conductance problem to increase the understanding of thermal transferring mechanisms in the boundary area and provide necessary guidance for future designing the composite interface structure. The links between the composite’s interface thermal conductance and thermal conductivity, which are built quantitatively via the developed models, were also reviewed in the last part. + +# Introduction + +Miniaturization of electronic devices makes it challenge to dissipate heat generated during operation. The insufficient thermal conductivity and heat dissipation rate of heat sinks would increase the electronic devices’ working temperature, and this dramatically degrades its working performance and efficiency. + +Besides, the electronic packaging becomes multiplelayer structures, and this causes more thermally challenging [1]. Until now, the most efficient cooling method for electronic system is still by fans and circulated cooling water, which consumes large amounts of additional electric energy [2]. According to one statistic from the US data center, one-third of the total power consumption is expected to be used in + +cooling the electronic system by 2025 [3]. The heat management problem is the bottleneck for sustaining and improving the performance of the high-power density electronic devices used everywhere today. The utilization of suitable materials with high thermal conductivity is essential for efficient heat dissipation from hot spots and improving the electronic device’s performance. To understand the mechanisms of interface thermal conductance can not only improve the thermal conductivity of the composites, but help assembly the heat management structure efficiently. + +Copper/diamond composites have the potential to be used as the next-generation heat sink materials in advanced electronic devices, and this is because (1) the artificial diamond as a metal matrix composites reinforcement has the highest thermal conductivity (TC) up to 2200 W/(m K) in nature [1, 4]; (2) the coefficient of thermal expansion (CTE) of copper/diamond composites could be tailored to be close to that of semiconductor chip materials (4–6 ppm/K) [5]; (3) the TC of copper/diamond composites could reach 500–900 W/(m K) [6–11]; and (4) the copper/diamond composite is stable at ambient temperature and has isotropic thermal conductivity, which doesn’t limit its wide applications [12]. However, the chemical affinity between the copper and the diamond particles is poor, which makes copper and diamond hardly bond together. The interface between the copper and the diamond particles is very important in the copper/diamond composites, and it has a bridging effect to help transfer the heat between the reinforcement and the matrix. Hence, the interface characteristics can influence the composites’ properties significantly [13, 14]. + +Carbide-forming metal elements such as Cr, B, Mo, Ti, and Zr can be added in the copper matrix in the copper/diamond composites to form a chemical bonding (metal carbides layer) between the copper matrix and the diamond particles, which helps bridge the diamond and copper matrix and improve the thermal boundary conductance [15]. A similar effect could be achieved by using the carbide-forming elements-coated diamond particles to fabricate copper/diamond composites. Most of the earlier literature’s themes on copper/diamond composites are toward the fabrication process of copper/diamond composites with high thermal conductivity value by designing different processing parameters based on materials thermal dynamics principle and + +collate the TC prediction value by different models to the experimental results [16–21]. Some research works characterized the morphology of the interface area and identified phases formed at the interface layer [9, 22–24]. Some of them modified the interface thermal conductance parts in TC models based on the interface characteristics to make the TC prediction value close to the experimental value [19, 25, 26]. For copper/diamond composites, it still lacks a deep understanding on how the interface characteristics affect heat transfer behavior. The research on the thermal conductance problems that have been carried out in other materials systems such as between specific metal elements and dielectric materials (including some dielectric materials like sapphire, diamond) can give us some insights to deal with the similar problems in the copper/diamond composites interface and provide guidance on the interface design [27–29]. We will delve into the related research and try to connect it with the interface thermal conductance problems in the copper/diamond composite. + +In this review, we begin with generalizing and comparing the commonly used fabrication methods of copper/diamond composites, to evaluate their primary experimental results and findings, and then address the factors that could affect the copper/diamond composite’s thermal conductivity. The interface-related problems, which determine the thermal conductivity of copper/diamond composites, are particularly emphasized and discussed in the perspective of classic lattice dynamic theory. Finally, the models for predicting the thermal conductivity of copper/diamond composites combined with the interface thermal conductance are reviewed. + +# Fabrication and thermal conductivity of copper/diamond composites + +# Vacuum hot pressing method + +Vacuum hot pressing (VHP) is a method that could hold pressure and temperature at the same time for powder consolidation (as shown in Fig. 1) [30], which is generally used for fabricating highly densified powder metallurgy materials. VHP usually holds the desired temperature and pressure for longer time compared to the spark plasma sintering (SPS) method (to be mentioned later). This trait promotes the + +![](images/5311b00c01850c8f1c18a7c1e7c4c6da40d146722de2f2df3e8969f14d37516c.jpg) +Figure 1 Vacuum hot pressing diagram [30]. + +element diffusivity between powder particles, thus enhancing the interface bonding strength between the copper matrix and the diamond particles. + +Some experimental results of the copper/diamond composites fabricated by VHP methods with different processing parameters are listed in Table 1. The thermal conductivity value of the fabricated copper/diamond composites varies with the contents of carbide-forming additives (B, $Z \mathbf { r } ,$ and Cr) that are added via using coated diamond particles or copper alloying [18–20, 31, 32]. + +Chu finds that the TC values increase first and then decrease with the $Z \mathbf { r }$ content increasing from 0 to 2.4 wt%, and the thickness of carbide interface layers is 0 nm, 270 nm, 320 nm, 480 nm, respectively, for the composites with Zr contents of 0, 0.8 wt%, 1.2 wt%, and 2.4 wt%, as shown in Fig. 2 [32]. It suggests that there exists an optimal content of carbide-forming additives to obtain a desired thickness + +of interface layers and high thermal conductivity for the copper/diamond composites. The interface layer with optimal thickness leads to high interface thermal conductance and improves the efficiency of heat transfer from the diamond particles to the copper matrix [34]. The low content of metal additives in the copper/diamond composites is not enough to form interface layers that completely cover all diamond particle surfaces, leading to reducing the interface thermal conductance. However, adding a large number of metal additives in the composite is prone to form an interface layer with large thickness, which is detrimental for improving interface thermal conductance. These align well with the experimental results shown in Fig. 2. + +Interface thermal conductance ‘‘hc’’ (as shown in Table 1) represents the thermal transfer efficiency from the diamond particles to the copper matrix. The higher the value is, the better the interface layer functions as a thermal transfer ‘‘bridge,’’ so that higher thermal conductivity of composites could be obtained. + +The ‘‘hc’’ values listed in Table 1 (including the ‘‘hc’’ in the table below) are obtained by prediction, not by experimental measurement. The principles of the prediction process are briefly illustrated in Fig. 3. When calculating ‘‘hc’’ with the model shown in the picture, ‘‘R’’ is defined as thermal resistance, and the relationship between $" R "$ and $^ { \prime \prime } \mathrm { h c } ^ { \prime \prime }$ is: + +$$ +R = \frac {1}{h c} \tag {1} +$$ + +For the case in Fig. 3, $R = R _ { \mathrm { C u / B 4 C } } + R _ { \mathrm { B 4 C } } .$ - $+ R _ { _ \mathrm { B 4 C / d i a m o n d } } . R _ { \mathrm { B 4 C } }$ can be obtained by the equation: $\begin{array} { r } { R _ { \mathrm { B 4 C } } = \frac { d } { \lambda } , } \end{array}$ where $" d "$ is the thickness of the $\mathtt { B _ { 4 } C }$ layer + +Table 1 Thermal conductivity and calculated interface thermal conductance of copper/diamond composites fabricated by VHP + +
AuthorsProcessing parametersRaw materialsDia. (μm)Vf (%)The thickness of the interface layer (nm)TC (W/m K)hc (MW/m2K)
Zhang et al. [33]900 °C and 80 MPa for 30 minW/Cu-coated diamond40055300721-
Sun et al. [18]950 °C and 60 MPa for 20 minB4C-coated diamond200501000687-
Hu et al. [31]950 °C and 50 MPa for 20 minB4C-coated diamond10060100065035
Chu et al. [32]980 °C and 42 MPa for 20 minCu-1.2%Zr powder1105532061561
Kang et al. [20]1150 °C and 20 MPa for 10 minMo2coated diamond7060-596-
Chu et al. [17]950 °C and 40 MPa for 20 minCu-0.8%B powder905828353840
Sinha et al. [34]950 °C and 45 MPa for 30 minCu-0.8%Cr powder1106017060121
+ +‘‘hc’’ is the interface thermal conductance by calculation + +![](images/a98a1eb2c07e0a5c82fcf3e2c96dec0e9c3e4d3bab36491c637051cb9460379b.jpg) +Figure 2 Measured thermal conductivity of Cu(Zr)/Dia composites versus zirconium contents [32]. + +and k is the thermal conductivity of $\mathrm { B } _ { 4 } \mathrm { C } . R _ { \mathrm { C u / B 4 C } }$ and $R _ { \mathrm { B 4 C / d i a m o n d } }$ are the Thermal Boundary Conductance (TBC), which will be addressed thoroughly in ‘‘Factors affecting the thermal conductivity of copper/diamond composites’’ section. This case indicates that the interface thermal conductance comprises two parts: one is the contacted TBC, and the other is the thermal conductance of the interface layer. + +The fabricated copper/diamond composites with different metal additives have varied thermal conductivity (see Table 1). Different metal elements have different intrinsic TC, phonon velocity and carbide transformation of interface layers and their solubility in Cu matrix, which greatly affect the thermal performance of the composites [35]. This finding is also applicable to the studies under the other three processing methods reviewed below. + +![](images/f6ddf1383a59a556bcbec047f1fd8524fbfb791e22f89e59d883ff4cd9cfd1a3.jpg) +Figure 3 Schematic illustration of the interface thermal resistance of $\mathrm { C u - / B } _ { 4 } \mathrm { C }$ -coated diamond composites according to the concept of an electrical resistance analogy [31]. + +Most of the copper/diamond composites are fabricated at the temperature below the molten point of copper $( 1 0 8 0 ~ ^ { \circ } \mathrm { C } )$ (see Table 1) via the vacuum hot pressing method, and this is because the pressure can promote the formation of complete interface layers on the diamond particle surfaces while holding the temperature for desired time. + +# Spark plasma sintering method + +The spark plasma sintering (SPS) method employs a high pulsed DC to heat powder or powder compact fast in an electrically conductive tool under vacuum or inert atmosphere condition, and small uniaxial pressure can be also exerted for densification (the diagram is shown in Fig. 4) [36]. High temperature is generated at the point contact between powder particles, which helps clean and melt the particle surfaces. Sintering necks are formed in a short time, and highly densified billets can be obtained generally. + +Effects of B, Cr, Si, Ti additives on the thermal conductivity of the SPSed copper/diamond composites have been studied (see Table 2). Except for the investigation on the interface layer structure in these studies, the effects of diamond volume fraction are noticed as well. For the fabricated copper/diamond composites with different metal additives added via coating or copper alloying, the results show different correlations. For example, Che finds that the thermal + +![](images/c1b96303e3fb3d3982d1957e6d874d6d59d3a13f09d5cdd8c10329bd5274d9c7.jpg) +Figure 4 Spark plasma sintering diagram [36]. + +conductivity of the composites with Ti-coated diamond is decreased when the diamond volume fraction is increased, as shown in Fig. 5 [37]. But in Ren’s work, the TC of the composites with Cr-coated diamond is increased with increasing the diamond volume fraction from 55 to 70 vol% [41]. In the case of adding Si-coated diamond, the composites’ TC is changed in the way as shown in Fig. 6, when the diamond volume fraction is increased from 40 to 60 vol% [39]. + +According to the rule of mixture, a higher diamond volume fraction in the composite should lead to the composite has higher thermal conductivity, because the diamond has very high thermal conductivity (1000–2000 W/m K). However, the composites with too large volume fraction of diamond are usually not easy to be densified during the consolidation process. It’s a trade-off between the processing and the proportion of diamond particles. Thus, this is why the thermal conductivity of copper/diamond composites is varied with the diamond volume fraction in the trends shown in Figs. 5 and 6. The optimal thermal conductivity of the composites can be achieved under the combination of high relative density and high diamond volume fraction. Considering the high energy the spark plasma instantly contains, the SPS processing temperature is usually maintained below 1000 C for avoiding the damage of diamond particles [5, 42]. + +![](images/03ef695eafd33b4243daa885ea3adbb92906b7c44984ccf02d4ec31901912ca2.jpg) +Figure 5 Variation in the thermal conductivity of Ti-coated diamond/Cu composites with different volume fractions of diamond with the thickness of the Ti-coated layer on the surface of diamond particles [37]. + +# Metal infiltration method + +In this method, diamond particles are homogeneously distributed in a crucible bottom by vibration first, and then the molten copper with metal additives is poured into the crucible. Low pressure (about a few tens MPa) is applied to the composites by pumping in the protected atmosphere in the airtight chamber for achieving densification. The process is illustrated in Fig. 7. Molten copper infiltrates between + +Table 2 Thermal conductivity and calculated interface thermal conductance of copper/diamond composites fabricated by the SPS method + +
AuthorsProcessing parametersRaw materialsDia. (um)Vf(%)The thickness of the interface layer (nm)TC (W/m K)hc(MW/m2K)
Mizuuchi et al. [36]900 °C 80 Mpa for 10 min(1.8–13.8 vol%) B additive30050-689-
Grzonka et al. [24]900 °C 80 Mpa for 10 minCu-0.8 wt%Cr powder18050-589-
Che et al. [37]1000 °C 50 Mpa for 10 minTi-coated diamond, Cu-0.5 wt%Ti alloy18050500630-
Mankowski et al. [38]1000 °C 50 MpaCu-0.8 wt%Cr powder20050-658-
Zhu et al. [39]900 °C 50 Mpa for 3 minSi-coated diamond30050-535-
Bai et al. [40]1030 °C 30 Mpa for 6 minCu-5 wt%B powder20044-660-
Rosinski et al. [9]900 °C 80 Mpa for 5 minCu-5 wt%Cr powder18050100–200--
Ren et al. [41]950 °C 40 Mpa for 20 minCr-coated diamond130 and 400 blended70600–900657-
+ +![](images/5149303b827b723a7f89df48343e1c70e720131941b881ed8c353a4eda6be144.jpg) +Figure 6 Effect of diamond particle volume fraction on the thermal conductivity of coated and uncoated composites [39]. + +the diamond particles, and the bulk materials can be obtained after copper is solidified. A detailed discussion of the infiltration mechanism in preparing copper/diamond composites can be found in Refs. [43, 44]. + +We can see that, from Table 3, the high thermal conductivity (up to 900 W/m K) can be achieved for the copper/diamond composites prepared by this method. This is because high temperature (100–200 C above the molten point of copper) is beneficial for the diffusion of metal elements. This method is favorable for forming a homogeneous interface layer on the diamond particles and overcoming the poor chemical affinity between the copper and the diamond. + +![](images/c2a1fb496dff88a24e90fafbdbd67ef078f94596c7e7fbfd3682fa6e905815ba.jpg) +Figure 7 Schematic illustration of the unidimensional infiltration process [43, 44]. + +In Li’s work, Zr element is added in the copper/diamond composites in the way of copper alloying, and he finds that the interface layer’s thickness of the composites is varied with adding different amounts of $Z \mathbf { r } ,$ thus leading to the composite’s TC value changed, as shown in Fig. 8 [19]. Based on Differential Effective Model (DEM) calculation, it finds that the composite’s TC value has a positive correlation with interface thermal conductance ‘‘hc.’’ ‘‘hc’’ is calculated by the models using the TC value of known composites. Interface thermal conductance correlates with the thermal conductivity of the composites positively. It can be reflected by Eqs. (19) and (20) (see ‘‘Model prediction of effective thermal conductivity of copper/diamond composites’’ section). + +Bai finds that the thin discontinuous jig saw interface layers formed between the copper and the diamond are more favorable for $^ { \prime \prime } \mathrm { h c } ^ { \prime \prime }$ than a plain interface layer; the morphology is shown in Fig. 9c. According to the thermal resistance formula, one $R _ { \mathrm { C u / d i a m o n d } }$ item is calculated in this way: $\begin{array} { r } { \frac { 1 } { R _ { \mathrm { n e w } } } = \frac { 1 } { \mathrm { R c } } + \frac { 1 } { R _ { \mathrm { C u / d i a m o n d } } } . ~ R _ { \mathrm { n e w } } } \end{array}$ RCu=diamond. can be taken as the interface thermal resistance (ITR) that is illustrated in Fig. 10b, Rc as ITR that is illustrated in Fig. 10c, and $R _ { \mathrm { C u / d i a m o n d } }$ as the ITR of bare Cu/diamond interface. In this way, the interface thermal conductance of the jig saw structure shown in Fig. 10b, which is corresponding to the structure in Fig. 9c, is larger than that of the planar structure shown in Fig. 10c, which is corresponding to the structure in Fig. 9d. When the morphology of the interface layer is distributed discretely, due to the formation of a new jig saw structure, the value of $R _ { \mathrm { n e w } }$ is smaller than Rc, and the thermal conductivity of the composite would be improved further. Besides, graphite is formed in the interface area under high temperature without adding the B element, which deteriorates the thermal performance of the composites [50]. Hence, it is necessary to add the appropriate content of carbideforming additives to hinder the formation of graphite. + +The effects of diamond volume fraction and processing parameters such as processing temperature and holding time have been also investigated on the copper/diamond composites prepared by the infiltration method, and the primary findings are similar to those of the composites fabricated by VHP and/or SPS methods. We can notice that the most important + +Table 3 Thermal properties of copper/diamond composites fabricated by metal infiltration method + +
AuthorsProcessing parametersRaw materialsDia. (μm)Vf (%)The thickness of the interface layer (nm)TC (W/ m K)hc(MW/ m2·K)
Abyzov et al. [45]1130 °C 10 Pa for 5 minW-coated diamond18063150520-
Weber et al. [46]1180 °C 0.6 MPa for 2 minCu-1%Cr alloy20060-600-
Cu-0.1%B alloy700
Dong et al. [43]1250–1450 °C for 30–150 min at vacuumTi-coated diamond110--385
Kang et al. [47]1150 °C 0.6 MPa for 10 minCr7C3-coated diamond70651000562-
Li et al. [26]1150 °C 1 MPa for 10 minTi-coated diamond70651292716-
Ma et al. [48]1100 °C 25 MPa for 10 minMo2C-coated diamond120–15060500657-
Wang et al. [6]1150 °C 1 MPa for 30 minCu-xCr alloy and Cr-coated diamond150–18060500810-
Kang et al. [20]1150 °C 20 MPa for 10 minMo2C-coated diamond7060100059645
Li et al. [19]1150 °C 1 MPa for 10 minCu-xZr alloy212–2506140093080
Li et al. [49]1150 °C 1 MPa for 30 minCu-xTi alloy23061200–300752-
Bai et al. [50]1150 °C 1.5 MPa for 30 minCu-xB alloy230673786890.9
Jia et al. [51]1350 °C for 45 minW-coated diamond150–18065-76824.1
Wang [52]1150 °C 1 MPa for 30 minCu-xZr alloy212–250-370930-
+ +![](images/e2b63c6f129e26cc53218d27bb12c4e0150f6633250bef1e531d7af6cb345407.jpg) +Figure 8 Measured thermal conductivity and interface thermal conductance calculated based on the differential effective medium (DEM) model for Cu–Zr/diamond composites [19]. + +factor, the interface thermal conductance, is not considered (see Table 3) in most of their works. We will discuss this key point in the ‘‘Factors affecting the + +thermal conductivity of copper/diamond composites’’ section. + +# High-temperature high-pressure method + +The principle of high-temperature high-pressure (HTHP) method is exerting high pressure and high temperature on the raw mixture materials to fabricate the composite. Usually, the pressure can be up to a few GPa and the temperature is above $1 1 0 0 ^ { \circ } \mathrm { C } .$ . Copper/diamond composites processed by this method can be highly densified bulk materials and have a tightly bonded structure between the copper matrix and the diamond particles. Moreover, the HTHP method is capable of fabricating copper/diamond composites with a very high volume fraction of diamond particles up to 90 vol%, which is limited to the former three processing methods we reviewed above. There are only several works reporting the research on copper/diamond composites fabricated by this method (see Table 4). + +![](images/bb50223aee77cb3ae6a50b99a8e92a134f913ca4910ed8873310d0a774fd4036.jpg) + +![](images/50d2267f3d3fc5a62f1d765d6e1a536aa0a7a035f8ac2ab3c809ebf62174a01b.jpg) + +![](images/af4b6cddd5b207db82f3740fe9960967ca89828976cf961d4769e2d6a9d26dd8.jpg) + +![](images/fb82667b4e9c756210c8ab5a4976b0f0a179270081ce49851bb9893008b43b3d.jpg) +Figure 9 TEM images of the interface structure of Cu– xB/diamond composites: a x = 0 wt%, b x = 0.1 wt%, $\mathbf { c } \ \mathbf { X } = 0 . 3$ wt%, and $\textbf { d } x = 1 . 0$ wt% [50]. + +Chen et al. find that the thermal conductivity of the HTHP-fabricated copper/diamond composite with carbide-forming element additives like Ti, Co, B, Zr is two times higher than that of composites from the pure copper and diamond powders [8]. This is + +because the latter has lower interface thermal conductance than the former. In other words, the carbide layer helps increase the interface thermal conductance. + +HE et al. find that adding different contents of carbide-forming elements like B, Zr leads to forming a different thickness of interfacial carbide layers in the HTHP-fabricated copper/diamond composites, thus resulting in obtaining different thermal conductivity, as shown in Fig. 11 [16, 22]. + +HE et al. used pure copper and diamond powder to fabricate the composites without adding carbideforming elements, and they made a comparison between the HTHP powder metallurgy method and the HTHP infiltration method. The results indicate the samples prepared by the infiltration method can obtain high thermal conductivity of up to 500–700 W/m K. However, the thermal conductivity of the samples fabricated by the powder metallurgy method is quite low. This is because an amorphous carbon layer is formed on the diamond surfaces under such conditions in the HTHP infiltration method, but is not formed by the PM method. The pure amorphous carbon interface layer helps enhance the bonding between copper and diamond particles [53]. The interface characteristics of the samples produced using this method are not well + +![](images/0161789bcf16874d8a6a17f2d71fe37e8dc0c071f3f2a4a02f37a092631682e2.jpg) + +![](images/34404148089597956e3b2f6b8e555c11acdfada0f0db8248e123890700306f94.jpg) +Diamond: + +![](images/7e4363c14ca033f862b0fae0a136bd26713187e5b9829109f08d1d4a288c6e76.jpg) + +![](images/3fdf7217637326073fecc8850e75d85a1ac237a6189e6432806ab13cf3f156ed.jpg) + +![](images/f13144cdf07dcc8b4f1ec277b2bc8278a78b86ff5ee483f6495d575efaa3cad5.jpg) +Cu: + +![](images/164acf235433a0717872bcb582ac4371a2637ca182151b4524a4ae15ddd120e4.jpg) + +![](images/f7ac0745111b43ca3f3b603da319513d0e385a83713c4be59a0d91e363735b22.jpg) + +![](images/fcfc47952ea72bef2d5599a02795fbaf7317d613859ffffe0096e853b775539e.jpg) +BC: +Figure 10 Schematic diagram of interface thermal resistance of Cu–xB/diamond interface: a = 0.1 wt%, b $\begin{array} { r } { x = 0 . 3 ~ \mathrm { w t \% } , } \end{array}$ , and c x = 1.0 wt% [50]. + +Table 4 Thermal properties of copper/diamond composites fabricated by the HPHT method + +
AuthorsProcessing parametersRaw materialsDia. (μm)Vf (%)The thickness of the interface layer (nm)TC (W/ m·K)hc (MW/ m2·K)
He et al. [22]1500 °C 5 Gpa for 10 min1 wt% Zr powder22090-677-
Chen et al. [8]1100 °C 5 Gpa for 10 minCu-1 wt%Co alloy500–60080-619-
Cu-0.3 wt%B alloy688
Cu-0.4 wt%Cr alloy683
Cu-1 wt%Ti alloy683
Chen et al. [53]1200 °C 5.3 Gpa for 10 minCopper and diamond powder20060-207-
Ekimov et al. [54]1300–2000 °C 8 GPa for the 20 sCu-30 wt%Ti alloy20080-900-
He et al. [16]1500 °C 5 Gpa for 10 minCu-0.3 wt%B alloy22090211075020
+ +![](images/df430336633116eed365f03bd16339f01d2df12d3a1e8d57e1dc57527a774042.jpg) +Figure 11 Thermal conductivity and thickness of interface layers measured for the HTHP-fabricated Cu–B/diamond [16]. + +investigated, comparing with that we reviewed above. The high thermal conductivity of copper/diamond composites prepared by the HTHP method could be attributed to the perfect bonding formed between the copper matrix and the diamond particles due to the ultrahigh pressure, as shown in Fig. 12d, e. + +# Electrodeposition method + +This method utilizes the electrochemical principle to effectively synthesize copper/diamond composites in the specific electrolyte, and the process is illustrated in Fig. 13. The electrodes are usually placed horizontally in the electrochemical solutions, diamond particles are first precipitated on the cathode substrate, and then copper is electrodeposited on the + +substrate to fill the gap between the precipitated diamond particles, resulting in the formation of a Cu/diamond composite. Compared to other hightemperature methods we reviewed, this method is conducted at ambient temperature and without pressure in one step, which makes the process simple and cost-effective. Besides, with proper controlling of the synthesizing parameters, copper and diamond can form chemical bonds without carbide-forming additives like other methods. However, this process is still in the early stage [55], and more investigations need to be done to understand the processing. To successfully fabricate a copper/diamond composite with good interface bonding and high thermal conductivity using this method, it needs to appropriately control the parameters such as electric current density, electrolyte types, and solution concentration. The primary processing parameters and thermal conductivity achieved for the fabricated copper/diamond composites are listed in Table 5. + +Avrai et al. successfully synthesized copper/diamond composites using the electrodeposition method and the highest thermal conductivity achieved is about 662 W/m K [57]. They demonstrate that this processing method can completely fill the spaces between the precipitated diamond particles and obtain dense composites. They also prove that the current density during the electrodeposition process is under galvanostatic conditions [56]. Hagio et al. find that the Cu ions in the electrolyte are not easy to concentrate on the surfaces of raw diamond particles compared to SiC-coated diamond, resulting in the + +![](images/7e946fcd8aac97e5b5e375cb5e34a278cc5dc52dd8abc95472c3ab7ed74576ae.jpg) +Figure 12 Micrographs of typical starting diamond powder used (a) and microstructure of the fabricated samples of Cu– diamond (b, c) and Cu–Ti– diamond (d, e) [54]. + +![](images/dbac4e28f4f67854196f56d443e77e13945abc6744564a4e4d32209b81d9c167.jpg) + +![](images/674dd83d264f5df4f4912e29b11877cb5dc5b1fc3f5014d5674ccf016f496a3e.jpg) + +![](images/dfca0cbe51e4cc589225c1c41bb40b13d56f2bf92ee02a3a71724dd9c3365aac.jpg) + +![](images/613c7ee9f7da027d50ec9a54b943514c2f2f96d8c21018b1914d22a03f7f579c.jpg) + +![](images/aef3c62f33a71af73f0f82db45f8eb94ec1402c30a7814422c1643269ca9f90e.jpg) +Figure 13 Schematic illustration of the process for fabrication of Cu/diamond composites using electrodeposition [56]. + +formation of poor interfacial bonding between the copper matrix and the diamond particles [58], and this may be attributed to the low efficiency of electrolysis [59]. Wu et al. find that two competitive additives ((DVF-B, accelerator, DVF-C, inhibitor) can affect the microstructure, crystallization, interfacial combination of the composite materials [59]. Specifically, DVF-B tends to promote copper fully filling in small and micro intervals formed by diamond particles, while DVF-C prefers to restrain copper deposition in large intervals and leveling copper nodule, and the two competitive additives work cooperatively to form the compactly combined interfaces of Cu–diamond composites. Thus, voids/gaps and nodules can be eliminated in Cu/diamond composites by adding the two additives in a proper ratio, leading to the well-combined interface and high TC. The specimen synthesized without and with adding these additives is referred to Process-C and Process-A, respectively, and the corresponding microstructures are shown in Fig. 14. Obviously, good interfacial bonding and flat surface are obtained in the specimen synthesized by Process-A, which would lead to high thermal boundary conductance and thermal conductivity of the composite. + +The electric current density is another crucial parameter for optimizing the structure-induced properties. Cho et al. [60] find that inappropriate current density can lead to porosity formation around the TiC-coated diamond particles during the + +process. As indicated in Fig. 15, the specimen with TiC-coated diamond should have higher thermal conductivity due to the higher thermal boundary conductivity. However, the current density of $5 0 ~ \mathrm { m A } / \mathrm { c m } ^ { \overset { . } { 2 } }$ leads to pores formed, because TiC is electrically conducting, which accelerates the deposition of Cu from different directions as indicated by the arrows in Fig. 15b and easily encloses a porous space. The pores can be eliminated by reducing the electric current density to $1 2 \mathrm { m A } / \mathrm { c m } ^ { 2 } .$ , and the resultant thermal conductivity improves to 557 W/ m-K [60]. + +The diamond particle size can affect the electric current density around the diamond particles [55], which further influence the microstructure of fabricated copper/diamond composites. When the diamond particle size is reduced from 420 to 66 lm, the electrodeposited copper matrix microstructures are changed from large grain/columnar (Fig. 16b) to fine grain/equiaxed (Fig. 16a) [55]. Large diamond particles are not only beneficial to obtain a long heat diffusion length/time in the inside of diamond particles but also improve thermal boundary conductance attributed to larger matrix grains at the vicinity of diamond particles. + +# Powder forging method + +Compared to other fabrication methods, powder forging can rapidly and cost-effectively prepare + +Table 5 Thermal conductivity and calculated interface thermal conductance of copper/diamond composites fabricated by the electrodeposition method + +
AuthorsElectric current density (mA/cm2)Electrolyte and diamond typeDia. (μm)Vf(%)The thickness of the interface layer (nm)TC (W/ m K)hc(MW/ m2K)
Arai et al. [57]5CuSO4·5H2O and H2SO4solution23061-662-
Raw diamond
Hagio et al. [58]100Idem0.012.72-46-
SiC-coated diamond
Arai et al. [56]5Idem10–23049-600-
Raw diamond
Wu et al. [59]20Idem + DVF-B and DVF-C10042-614.87-
Raw diamond
Cho. et al. [55]50Idem66,42042.1,-222,4541.3,3.1
Raw diamond68.2
Cho et al. [60]12Idem40034.7-5579.7
TiC-coated diamond
+ +![](images/1a9d8d005ea11da79287c0629ad3d3181fc39df6675cef9925fe9f9f37ae1ca4.jpg) + +![](images/fa946fb29e2708057a649e1a83c81eeb8d8fdcb3e57bbd81dfd6da9e5fbb3602.jpg) + +![](images/6f37f718fd82b57181a765b69acc387710880385b575462f66c4c38b711d6d48.jpg) + +![](images/0a4a43719994233a106d9ab0213ab32fa8944168d8d5da32e35fadb53997ae34.jpg) + +![](images/3f03f3826be24df085d26edd4cadd479ae2cbcf17a09f7bc8f46fef5c2a20807.jpg) +(f) + +![](images/89bf098a66e27402b4a6e749c0a441925fad780ba01023e15dcc9369627802d1.jpg) +SEM image of process-A composite (Cu–D–A); f interfacial schematic of Cu–D–A composite [59]. + +Figure 14 SEM images showing the microstructure of the thick composites: a, b sliced SEM image of process-C composite (Cu– D–C); c interfacial schematic of Cu–D–C composite; d, e sliced + +![](images/9a284bd240f341100db8c1975c2fcb1d87a9a34958d40c2c37dd9bd260176685.jpg) +Figure 15 Metal matrix growth during the electrodeposition of composite materials containing a particles that do not conduct electricity (i.e., copper plating with uncoated diamond particles) and b particles that conduct electricity, which can cause porosity formations during the co-deposition process (i.e., copper plating with TiC-coated diamond particles) [60]. + +![](images/d9c18b27881a6359beb59fae731e35fbe86315b9ed4d8cf5bde180d081502338.jpg) + +![](images/78ed5a2f337fa7d28d9df375d8adbc2a7ac0a9da35c5f32a42c55a946a15cd4f.jpg) +(a) + +![](images/13c778dbf0153b711ab23698aef9cb1f85454397644e01181eae7b3c423e8528.jpg) + +copper/diamond composites without using expensive manufacturing equipment, and it offers an alternative and potentially more effective fabrication method for producing copper/diamond composites. Yang et al. first applied the powder forging technique to producing copper/diamond composites from the powder mixture of artificial diamond and elemental copper powders [61–64]. The carbide-forming elements such as titanium and chromium are introduced in the materials system via adding alloying + +element powders in the powder mixture or precoated the alloying element on the diamond particle surface. The processing route is shown in Fig. 17, and it primarily includes powder mixing, powder compaction, and then powder forging (heating the compact to the desired temperature in the atmosphere protective chamber using an induction furnace and then forging the hot compact to form a densified copper/diamond composite pancakes/billets). + +![](images/cff55b00173f257cc36cd0449b160d4d42c154d8913c9f2a6bb4610a4353f7ec.jpg) + +![](images/643350955c6b04001708bbad3be945829aa4ab82e34416ed26ee793d0701274d.jpg) +Figure 16 BSE micrographs of copper matrix around a 66-lm diamond particles and b 420-lm diamond particles. Copper grains near 66-lm diamond particles are generally small compared to + +Yang et al. use the powder forging technique successfully fabricated copper/diamond composites with a diamond volume fraction of up to 65 vol%; however, the copper-55 vol% diamond composites have much higher thermal conductivity compared to the composite with 65 vol% of the diamond, attributed to the diamond particles uniformly dispersed in the copper matrix without agglomerations. Through detailed interface characterization and analyses, Yang et al. find that the semi-coherent relationship is established between the interfacial layer (TiC) and the copper matrix, and dispersed nano-spherical particles (TiC) formed on the diamond particles help strengthen the interface bonding between the diamond and the copper matrix. All these characteristics facilitate the heat transfer across the interface layer, contributing to the improvement in thermal conductivity for the forged copper/diamond composites. However, the formation of a thin amorphous carbon + +those near 420-lm diamond particles. The grains near 420-lm diamond particles tend to be columnar, especially near interfaces [55]. + +layer (for the 1050 C forged composite [61]) in the diamond particle (next to the interfacial layer) and deformed structure in the copper matrix have an adverse effect on the thermal conductivity of the formed composites. The hot-forged copper-55 vol% diamond (coated with Ti) composite shows a high thermal conductivity of 550 W/m-K and a CTE of 7 ppm/K at 313 K, and a flexural strength of 418 MPa [62]. These results are better than those of most reported copper/diamond composites (with the diamond particle size of \ 100 lm), suggesting that the powder forging technique is feasible to produce a copper/diamond composite with an acceptable mechanical and thermo-physical performance for practical applications. + +![](images/5fc1df559b78a4d8d4cbe9687884ac987daec451a0e40d7267c1f28856b50124.jpg) +Figure 17 Processing flow chart of fabricating copper/diamond composites by powder forging technique. + +# Factors affecting the thermal conductivity of copper/diamond composites + +# In copper matrix + +Heat transfers in the materials via heat carriers like atom vibration in the form of phonons or electrons at room temperatures or above. For metal material, the electrons dominate the heat transferring. The defects in the materials such as pores, grain boundaries, and dislocations increase the probability of electron scattering during the transmission process, therefore reducing the overall heat transfer ability [28, 65, 66]. Furthermore, the metal additives used for improving chemical affinity between the copper and the diamond, by forming a carbide interface layer, may form solid solutions with copper. The higher the concentration of metal additives is in the copper matrix, the lower the thermal conductivity of the copper matrix would be. The correlation between Cr concentration in the copper matrix and the thermal conductivity of the copper matrix is shown in Fig. 18 [46]. The solutes of Ti and B in the copper matrix degrade the copper matrix’s TC to a different extent [8, 16]. + +# In diamond + +Most of the diamond particles used for preparing copper/diamond composites (as reviewed in ‘‘Fabrication and thermal conductivity of copper/diamond composites’’ section) are synthetic single crystal diamond, and its size is in the range of 30 lm to 300 lm. Besides the completeness of diamond particles (the cracks and damaged parts in the diamond particle are treated as defects), the content of nitrogen mainly affects the thermal conductivity of diamond. According to the nitrogen measurement method, the relationship can be expressed in the formula: + +$$ +\lambda = 2 2 0 0 - 3. 2 7 [ N ] +$$ + +where k is the measured thermal conductivity of diamond and N represents the nitrogen content [67, 68]. The higher the nitrogen content is, the lower the thermal conductivity of the raw diamond particles would be. + +# Interface thermal conductance + +Interface thermal conductance (ITC, ‘‘hc’’ aforementioned in Tables 1, 2, 3, 4, 5), as an indicator, + +![](images/122b67a63ceeace89680a86ab1152e61cca48d861a2e93794cd0ca3c96466220.jpg) +Figure 18 Evolution of thermal conductivity and coefficient of thermal expansion in a Cu–Cr/diamond composite as a function of the chromium content in the matrix [46]. + +determines the ability of heat transferring between the copper matrix and the diamond particles. If the ITC between the copper matrix and the reinforced diamond particles is high, it suggests that the high thermal conductivity of the diamond particles is exploited and the efficiency of heat carriers passing through the ‘‘bridge’’ (interface layer) is high. Otherwise, the thermal conductivity of copper/diamond composites cannot be improved. ITC consists of two parts, thermal boundary conductance (TBC) and thermal conductance of the interface layer, as shown in Fig. 19. G represents TBC and K is the thermal conductivity of the interface layer material. + +If the interface layer consists of more than one layer (as shown in Fig. 19), it would complicate the problem in the way that more interface layers with different characteristics of thermal boundary conductance and thermal conductance need to be considered. These two parts are reviewed in sequence as below. + +# Thermal boundary conductance + +This definition is first appeared in the research on the thermal transfer problem between liquid helium and solids interface under extremely low temperature + +![](images/6d4f5dfa84865af5aa13d0920048d0a7b3ec024b676ffbb3317707044cc74470.jpg) +Figure 19 Diagram of interface thermal conductance. G1 and G1 represent TBC of the boundary, respectively; $K _ { \mathrm { d } }$ is the thermal conductivity of the material in the interface layer. + +(near absolute zero), known as Kapitza resistance [69]. A previous review of this problem can be found in the seminal article authored by Swartz and Pohl [70]. The thermal boundary resistance comes from the experimental observation of the temperature discontinuity when heat flows across the contacted interface. Hence, the mathematical definition of thermal boundary conductance (TBC) is given by: + +$$ +G = \frac {Q}{A \times \Delta T} \tag {2} +$$ + +where Q is the heat flow, A is the area across the heat flow, and DT is the temperature drop across the boundary. Kapitza resistance, the interface thermal resistance, is the inverse of G in Eq. (2), as expressed in Eq. (1). Recall the physical definition of thermal conductivity, we could find that TBC is the thermal conductivity in unit length by definition. The difference lies in that thermal conductivity is a parameter for bulk materials, whereas TBC is for the boundary without thickness. The mean free path of heat carriers is the determinant for the materials’ thermal conductivity, whereas the transmission probability of heat carriers across the boundary is for TBC. This phenomenon is not only found in the helium and solids system but also existed in solid and solid interface [70]. Even for the interface between two identical materials, given the definition above, thermal boundary conductance or nonzero thermal boundary resistance exists. The differences of lattice + +characteristics and density of states for the materials at the two sides across the boundary lead to the discontinuity of temperature, because of the existence of the interface [27, 71, 72]. In contrast, the identical materials in unit length do not have this problem. + +In a common view, electrons are dominant carriers for most of the heat transport in metals while phonons are heat carriers in dielectric and semiconductors at room temperature (above 300 K); the thermal transfer problems for copper/diamond composites application in this review are all at room temperature or above [73]. For the cases in copper/diamond composites, the interface may comprise metal/metal, metal/dielectric (diamond), metal/carbide, and carbide/dielectric contacts. The electronic structures of metal carbides such as titanium carbide, chromium carbide are quite complex. For example, chemical bonding in chromium carbides has a complex mixture of metallic, covalent, and ionic characters, and this means both electrons and phonons could participate in the heat transfer process [74–76]. Hence, what matters is the variety of particles that contribute to the thermal transport process across the interface and to quantify and formulate the process. For metal/dielectric and carbide/dielectric contacts, one material is a nonmetal, and the much related research proves phonons are the dominant heat carrier [28, 29, 77–79]. For metal/carbide contact, it is considered as a metal /dielectric problem in many research works. As analyzed above, the induction of electrons adds extra complexity in the case of copper/diamond composites, because of the interaction between phonons and electrons. Many works on copper/diamond composites ignore the contribution of electrons but focus on phonons, and the predictions of thermal conductivity do have discrepancy compared to the experimental results [50, 80]. Sadasivam uses the Atomic green function method and integrates the electron–phonon coupling effect in the function to deal with metal–semiconductor thermal boundary conductance problems [81]. The structure of the semiconductor is as complex as carbide, so we can put it here as a reference for comparison. Although the results from Sadasivam’s research are very close to the prediction value by only considering phonons, the contributions of electrons and interaction between electrons and phonons are still worth studying in the future. To identify the heat carrier dominant in the heat transport process across the interface helps us use appropriate theories and + +methods to discuss the problem. From the analyses above, phonons are dominant heat carriers in most cases in the heat transfer process that happened in the copper/diamond composites system. + +There are several methods to quantify TBC based on the analysis of phonon characteristics on both sides of the contact. Classical lattice dynamical theory is quite powerful to make prediction; we mainly adopt this perspective to elucidate this problem. The relatively rigorous approaches such as Molecule dynamics simulation and non-equilibrium Green’s function method require accurate atomic-level interface boundary information plus huge computational expenses, which are not included in this review but can be found in other works [27, 82–85]. + +The classical lattice dynamical theory for describing thermal boundary conductance by phonons can be expressed as follows: + +$$ +G _ {\mathrm {k}} (T) = \frac {1}{V} \int \frac {\partial}{\partial_ {T}} \sum_ {\mathrm {k j}} \hbar \omega_ {\mathrm {k j}} D _ {\mathrm {k j}} (\omega) n \left(\omega_ {\mathrm {k j}}, T\right) v _ {\mathrm {k j}} \alpha_ {\mathrm {k j}} \mathrm {d} \omega \tag {3} +$$ + +where subscripts k and j mean phonon wave vector and phonon mode, respectively, $\omega _ { \mathrm { k j } }$ is the phonon angular frequency, ${ \mathbf { } } J _ { \mathrm { k j } } ( \omega )$ is the phonon density of states, $v _ { \mathrm { k j } }$ is the component of phonon group velocity for the interface, $n ( \omega _ { \mathrm { k j } } , T )$ is the Bose–Einstein distribution function, V is the volume of the object. This formula is constructed based on the Landauer theory [82]. This formula means that phonons in the first Brillouin zone with $v _ { \mathrm { k j } }$ velocities transmit through the boundary with the transmission probability of $\alpha _ { \mathrm { k j } }$ . The subscripts k and j are phonon wave vector and polarization, respectively. Thus, the thermal boundary conductance G can be calculated by the integral of $G _ { \mathrm { k } }$ over all of the allowed phonon wave vectors and polarizations. Within this expression, the distinction among the models introduced next lies in the transmission coefficient parameter $\alpha _ { \mathrm { k j } }$ . + +(A) Acoustic mismatch model Acoustic mismatch model theory (AMM) is based on the following assumptions: + +(1) The interface is a perfectly specular plane, phonons incident on the interface are either refracted or reflected, according to the critical angles which are decided by Snell’s law [70], and any elastic and/or inelastic scattering behaviors are ignored; + +(2) Transmission coefficients are determined by the difference of sound velocities and materials density of materials on the two sides across the boundary; one classical formula is given: + +$$ +\alpha_ {\mathrm {k j}} = \frac {4 Z _ {\mathrm {A}} Z _ {\mathrm {B}}}{\left(Z _ {\mathrm {A}} + Z _ {\mathrm {B}}\right) ^ {2}} \tag {4} +$$ + +where subscripts A and B represent the two sides across the boundary; Z is acoustic impedance given by the product of material mass density and sound velocity lying on that side; and + +(3) Phonon frequency $\omega _ { \mathrm { k j } }$ on the side A can only couple with the phonon on the side B with the same frequency; this confines that there are limited phonons that can be transmitted through the interface; especially, the phonon density of states on the two sides is highly mismatched. + +The thermal boundary conductance under the AMM case could be calculated by integrating Eq. (3) substituting $\alpha _ { \mathrm { k j } }$ by Eq. (4) within certain angles (which are decided by assumption (1)). AMM has been proved that the prediction result has good agreement with the experimental data below a very low temperature, about 30 K. This is because highfrequency modes of phonons that participate in the thermal transfer process are not considered by AMM. There are many works about copper/diamond composites we reviewed above to use AMM methods to calculate the thermal boundary conductance [16, 19, 20, 86]. The predictions do have large discrepancies compared to the experimental results. By putting the predicted TBC value in the effective models to calculate the thermal conductivity of copper/diamond composite, which will be reviewed in the following section, the predicted composite’s thermal conductivity is comparable with the experimental measurements. This is an indirect method to verify the effectiveness of the AMM model. In another work [87], authors directly compared the prediction values by AMM to the experimentally measured TBC value by TDTR technique (which will be introduced later), and it clearly shows the discrepancy between the predicted and experimental values. This means the assumptions in AMM theory need to be further refined. + +(B) Diffuse mismatch model The diffuse mismatch model (DMM) is based on different assumptions about how phonons carry heat energy transmitted through the interface, comparing the AMM. Swartz first introduces the concept by assuming that all the phonons are diffusely scattered across the interface [79], not like the case in AMM that only phonons within certain velocity angles can pass. The main assumption is that the maximum frequency is the cutoff frequency of materials with the lower Debye temperature side and phonon frequency $\omega _ { \mathrm { k j } }$ on side 1 (A) can only couple with the phonon on side 2 (B) with the same frequency. As mentioned before, the only difference between these models lies in transmission probability a (because the phonon energy -hx is independent of its directions and modes). In the DMM model, the transmission probability follows: + +$$ +\alpha_ {i} (\omega) = 1 - \alpha_ {3 - i} (\omega) \tag {5} +$$ + +where i can take 1 and 2. $\alpha _ { i } ( \omega )$ represents transmission probability from the side $i ; 1 - \alpha _ { 3 - i } ( \omega )$ is the reflection probability from the other side. Equation (5) indicates transmission probability from side 1 (A) equals reflection probability of side 2 (B) and vice versa. Heat flux from the side i can be expressed as the number of phonons with energy -hx leaving: + +$$ +\sum_ {j} \int_ {0} ^ {2 \pi} \int_ {0} ^ {\pi / 2} \mathrm {d} \theta \cos \theta \mathrm {d} \phi v _ {i, j} D _ {i j} (\omega) n (\omega , T) \alpha_ {i} (\omega) \tag {6} +$$ + +Since it is independent of angles, Eq. (6) becomes: + +$$ +\frac {1}{4} \sum_ {j} [ D _ {i j} (\omega) n (\omega , T) ] \alpha_ {i} (\omega) \tag {7} +$$ + +Heat flux on the two sides across the interface should be equal; based on this, we get: + +$$ +\sum_ {j} v _ {i, j} D _ {i j} (\omega) n (\omega , T) \alpha_ {i} (\omega) = \sum_ {j} v _ {3 - i, j} D _ {3 - i, j} (\omega) n (\omega , T) [ 1 - \alpha_ {i} (\omega) ] \tag {8} +$$ + +Thus, we can have $\alpha _ { i } ( \omega ) \colon$ : + +$$ +\alpha_ {i} (\omega) = \frac {\sum_ {j} v _ {3 - i j} D _ {3 - i , j} (\omega)}{\sum_ {i , j} v _ {i , j} D _ {i j} (\omega)} \tag {9} +$$ + +The thermal boundary conductance under the DMM model then can be calculated by replacing $\alpha _ { i } ( \omega )$ in Eq. (3) with Eq. (9). + +(C) TBC across the metal–metal contacts by electrons When the electrons are dominant as heat carriers across the metal–metal boundary, some descriptions in Eq. (3) are changed, but the principle is preserved. The model can be formulated as: + +$$ +G _ {e, 1 2} = \frac {1}{4} \int_ {0} ^ {\infty} (\varepsilon - \varepsilon_ {F, 1}) D _ {1} (\varepsilon) \frac {\partial f _ {1}}{\partial T} v _ {F, 1} \zeta_ {1 2} d \varepsilon \tag {10} +$$ + +where $G _ { e , 1 2 }$ represents the electron thermal boundary conductance across the metal 1 (A) and metal 2 (B), e and $\varepsilon _ { \mathrm { F } , 1 }$ are electron energy and Fermi energy of the metal 1 (A), $D _ { 1 } ( \varepsilon )$ is the density of state, f is the Fermi–Dirac distribution function in contrast with that of phonons, $v _ { \mathrm { F } , 1 }$ is Fermi velocity, and $\zeta _ { 1 2 }$ is the transmission coefficient. Still, transmission probability is the parameter that needs to be calculated. Some researchers find that assuming electrons transfer across the boundary diffusively can make the prediction and experimental results consistent [88, 89]. Similar to the process for solving the transmission coefficient in the last section ‘‘Diffuse mismatch model,’’ we have: + +$$ +\begin{array}{l} \int_ {0} ^ {\infty} \left(\varepsilon - \varepsilon_ {\mathrm {F}, 1}\right) D _ {1} (\varepsilon) f _ {1} v _ {\mathrm {F}, 1} \zeta_ {1 2} (\varepsilon) d \varepsilon \tag {11} \\ = \int_ {0} ^ {\infty} \left(\varepsilon - \varepsilon_ {\mathrm {F}, 2}\right) D _ {2} (\varepsilon) \left(1 - f _ {2}\right) v _ {\mathrm {F}, 2} \zeta_ {2 1} (\varepsilon) \mathrm {d} \varepsilon \\ \end{array} +$$ + +Then, for approximation, based on the assumptions of (1) electrons on both the metal sides have the same Fermi energy and (2) energy after transmission will not change; we get the formula of $\zeta _ { 1 2 } ( \varepsilon )$ : + +$$ +\zeta_ {1 2} (\varepsilon) = \frac {D _ {2} (\varepsilon) (1 - f _ {2}) v _ {\mathrm {F} , 2}}{D _ {1} (\varepsilon) f _ {1} v _ {\mathrm {F} , 1} + D _ {2} (\varepsilon) (1 - f _ {2}) v _ {\mathrm {F} , 2}} \tag {12} +$$ + +Because the results at room temperature are of our interests, by substituting the $\zeta _ { 1 2 } ( \varepsilon )$ in Eq. (10) with Eq. (12), the final $G _ { e , 1 2 }$ can be expressed as: + +$$ +G _ {e, 1 2} = \frac {1}{4} \frac {\gamma_ {1} v _ {\mathrm {F} , 1} \gamma_ {2} v _ {\mathrm {F} , 2}}{\gamma_ {1} v _ {\mathrm {F}} 1 + \gamma_ {2} v _ {\mathrm {F}} 2} T \tag {13} +$$ + +where $\gamma$ is the coefficient of electron heat capacity. Details on the derivation of this result can be found in Refs [89, 90]. + +(D) Factors influencing TBC The characteristics of the boundary between the two materials affect the thermal boundary conductance. + +(1) Roughness + +Experiments conducted by Patrick’s group demonstrate that the thermal boundary conductance + +is decreased with the increasing roughness of Si substrate coated with Al thin film (as shown in Fig. 20) [91]. The DMM modeling curve in Fig. 20 is derived from Eq. (3), but the roughness factor is introduced into Eq. (3) in the way formulated in the research work [92], and the calculated results align well with the experimental data that are measured by time-domain thermo-reflectance (TDTR) method. The roughness can be regarded as defects in the lattices and arouses phonon scattering, and the decrease in boundary roughness helps improve the TBC, as suggested in Fig. 20. + +Regardless of the silicon roughening procedure (indicated in the plot), the thermal boundary conductance follows a similar relationship with RMS roughness as shown in Fig. 20. The accuracy of the AFM (atomic force microscope) tip is only about 10 nm, and this may lead to limitation and deviation for the measurement, thereby affecting the final results [28]. + +When the length scale of the roughness is larger, the tendency could be different. Researchers try to increase the interfacial contacting area by etching the diamond surface to form a zigzag structure [94], as shown in Fig. 21a, b, and they find that the etched structure could improve the efficiency of the thermal energy transmission. This is because the etched surface could absorb the reflected energy wave (phonons and electrons) many times (as indicated in Fig. 21e), resulting in doubly or more efficient heat transfer and + +![](images/8bc2ff2f10caf519c3a3b29932cd46c5e19c2f23ef0e5ffe5cf4c9075a3a7d8e.jpg) +Figure 20 Room-temperature predicted (line) and measured thermal boundary conductance et al./Si interfaces plotted as a function of interface roughness [91]. Blue dot data are from [91]. Red dot data are from [92], Green dot data are from [93]. + +higher thermal TBC compared to that of un-etched surface [94]. Moreover, the formation of Si nanopillars on the Si surface, as illustrated in Fig. 22, can also enhance the TBC between Si and Al [95]. Compared to the planar surface, the rough surface caused by the formed nanopillar has a larger effective contact area, contributing to the significant increase in TBC up to 90% [95]. + +# (2) Atomic intermixing + +The assumption of an atom sitting on the two sides of the interface which can be contacted without mixing is a perfect situation. When several layers of atom diffuse across the interface boundary to form an intermixing layer; it can affect the thermal transfer process and TBC. One study investigates the intermixing thickness of Cr/Si’s effect on thermal boundary conductance from both experimental and calculation perspectives, and it finds that TBC decreases with an increase in the thickness of the intermixing layer [96]. Figure 23 indicates the predicted results obtained from the modified DMM model have good agreement with the experimental data. The principle of DMM modification is to add a coefficient about the interfacial mixing layer in Eq. (3), and the details can be found in Ref [96], and this applies to the metal/metal interface [89]. + +# (3) Interfacial bonding strength + +Both the interfacial mechanical bonding and chemical bonding affect the TBC. In general, there is a positive correlation between interfacial bonding strength and TBC. Mechanical bonding in the interface area is a kind of physical contact bonding + +![](images/8426bed6d5e53518ff42089811375afad33f8461158886751458931934cbcafa.jpg) + +![](images/f1e59d4456f092e36e73792a652746ee607dc78b5727c03afa970cd1099299d9.jpg) + +![](images/0f05ec69dbaeb7c233da8683b9a6e7f7a642edd082cd40e02424c3ae8cc635cb.jpg) + +![](images/bcc06f336a7ee8b09f028a24abb7d2fd7cd9ead4407c118da3b482f893142c36.jpg) + +![](images/7f936ef01da24ad7fc3a50a24cfd38fc6933ceb0a8b6919ed89a6019f61c1771.jpg) +Figure 21 Schematic diagram of etching pits and interfacial heat flux: a, b etching and interface coupling area; c flat interface, d rough interface, e single coupling at the pits [94]. + +![](images/a1c5922f63a1e873a6b04e2474139a535867920f130d1e1d891375429ecbac2c.jpg) + +![](images/00fe1fa2bdf338af9cc14afc05da1692addcce6bc6590d295456f9bd78d118dd.jpg) + +![](images/aa64533b8787360d90180131ca0b1697ad7f12fa88fcb3fc790dc1238d3ae6ed.jpg) + +![](images/78cacaee4427072642dfd4635def591828dd3296eb212eb48ad2a000c6a9c373.jpg) + +![](images/848cfb8b06d97aa9acad0107bf0e9c76b7b19204a345797b203c06efb1e957d8.jpg) +Figure 22 SEM images of a periodic and b random nanopillars fabricated on Si substrates using EBL, c schematic of a nanostructured Si surface capped with a 100-nm Al layer as the TBC is being monitored by TDTR, d schematic illustration of thermal transport enhancement by a nanopillar in comparison with + +without forming chemical bonds between the two phases. Strong mechanical bonding usually exhibits higher TBC. As shown in Fig. 12e, the copper/diamond composite fabricated by the HPHT method has tight interface bonding, which leads to the composite has higher thermal conductivity. Yang et al. also find that the nano-spherical TiC particles grown on the diamond surfaces can enhance the interlocking with copper matrix (Fig. 24a), thus increasing the interfacial bonding strength, and this is more beneficial for obtaining high TBC compared to the planar TiC interface layer in Fig. 24b [62]. + +The strength of chemical bonding is usually much higher than mechanical bonding; however, copper has a poor chemical affinity with the diamond + +![](images/defd49d87aa9f793e82dd6dd5a23db4f5e98f4c2005eda082c09a584d7c14458.jpg) +Figure 23 Comparison of the VCDMM to the experimental data on samples Cr/Si with different interfacial mixing thicknesses [96]. + +the case of a planar interface (dotted lines indicate that heat can also pass through the sidewall of the nanopillar to cross the interface) and e normalized TBC as a function of the number density of nanopillars for the periodic (circle) and random (square) nanostructured Si/Al interfaces [95]. + +surface, making it hard to form efficient bonding between. Two ways have been used to facilitate the formation of effective chemical bonding between the copper and the diamond. One is forming the carbide layer between the copper matrix and diamond particles, which has been reviewed in ‘‘Fabrication and thermal conductivity of copper/diamond composites’’ section. In this way, the carbide layer has strong chemical bonding with diamond, and the chemical affinity between the copper and the metal carbides is better than that between copper and diamond [19, 97]. The other one is by modifying the surface chemistry. Mark finds that the transition from van der Waals to covalent bonding increases the ITC by *80% after the Au/Qz interface is chemically modified, as shown in Fig. 25 [98], demonstrating that the increased interfacial chemical bonding can increase the TBC because of the covalent bonding is stronger than van der Waals. Hopkins finds that the graphene surface that is functionalized by Ar:O can boost the TBC between the graphene and Al, and this is because the two electrons of O can form covalent bonding and combine with Al, respectively [99]. + +The interfacial bonding strength is hard to measure directly, due to its small scale in the composite materials. Many related research works combine mechanical properties and fractography around the reinforcement of the composites to assess the interfacial bonding [100–102]. For example, Wu et al. derive an equation to evaluate the interfacial bonding energy theoretically [103], which is expressed as: + +![](images/3356136f60ae1e8da93f7c272e114e988d8bd4832129f876fb6e8fd9d757854a.jpg) +Figure 24 Sketch of toughening and reinforcing mechanisms for a 800–Cu/ 55Dia composite and b 1050C–Cu/55Dia composite; and schematic illustration of heat transport in c 800C–Cu/55Dia composite and d 1050-Cu/55Dia composite [62]. + +![](images/06d8ce99dbbff35a8d411aed08ddf6151c59bf29870b86f60c4e1f62fd84454e.jpg) + +![](images/ed785a5887a81fe208d01245be21f8dbf3ff479aa2c50f29d4f1a867c1bcbff2.jpg) + +![](images/175cae67e12cc0a2495c911e24da9ece84a874a9bd6735a8e325408da8b56fef.jpg) +Figure 25 Tuning interface thermal conductance. Plot of interface thermal conductance (G) as a function of the methyl:thiol endgroup ratio for 0%, 25%, 50%, 75% and 100% thiol end groups. Duplicated Al for each ratio was measured. Error bars represent uncertainty in TDTR data fitting [98]. + +$$ +\gamma = \frac {a \sigma_ {\mathrm {c m}} ^ {2}}{1 2 E _ {\mathrm {m}} \left(\left(1 / \bar {\varepsilon}\right) + v _ {p}\right)} \tag {14} +$$ + +$\gamma$ is interfacial bonding energy, $\sigma _ { \mathrm { { c m } } }$ is ultimate tensile strength, a is volume to surface area ratio of the reinforcement, $E _ { \mathrm { m } }$ is Young’s modulus of the matrix, e is equivalent Young’s modulus of the composite, and $v _ { p }$ is the volume fraction of the reinforcement. From Eq. (14), it is apparent that the interfacial + +bonding energy is highly associated with these mechanical parameters, which can be obtained through mechanical testing. Wu et al. also prove that the fracture surface, which has the characteristics shown in Fig. 26c, has the highest interfacial bonding energy. This is attributed to no gaps existed between the diamond particles and the matrix. Furthermore, dimples on the diamond surfaces suggest the metal matrix bonds tightly with the diamond so that the metal side deforms plastically before it detaches from the diamond, compared to the composites that have the microstructures shown in Fig. 26a, b, d. + +(E) Measurement of TBC Time-domain thermos-reflectance (TDTR) is a optical technique, which can measure the surface temperature changes from the surface optical reflectivity without contacting. It has a laser pulse and a pump probe, with which the thermal boundary conductance between two the materials with very small scales can be experimentally quantified. This technique has become popular and standardized over the past 20 years [104]. With the progress of pulsed laser systems, this thermos-reflectance technique can employ ultrafast laser with very short pulses to characterize the temperature changes in the time domain. In this way, it reduces the errors radiative loss and contact issues that are concerning in other measurement techniques like the 3x method [28, 105, 106]. Thorough descriptions of the measurement methods and data analysis can be found in Refs [96, 107–109]. + +![](images/4e0115dddc0b077a58a4467e68d89d503a57614af025408ec4c8c1e9b846e6ec.jpg) + +![](images/7a1923b3ca976208442d96dc34adafc8a3bf7baa51dcd7eeb266b27e387da2a6.jpg) + +![](images/cef2fd5f888b8a26d69b439688a22c2d9cfc40edf231308711a5751ec34a5e29.jpg) + +![](images/dda9f082b65680f427eb03f150b9516ca5725c9a45ea9941e23f0bdd23e4e960.jpg) +Figure 26 FE-SEM micrographs of tensile fracture surfaces of the Al/Ti diamond composites with different plating times of a uncoated, b 30 min, c 90 min, and d 180 min [103]. + +TDTR technique is just applied for investigating the interface thermal conductance problem of copper/diamond composites recently [80, 87, 110]. The general principle is that the interface scenario in the copper/diamond composites can be physically simulated through creating a Ti layer by magnetron sputtering onto a diamond substrate, as shown in Fig. 27. Then, the heat treatment is conducted for forming TiC to simulate the real scenario at the interface of copper/diamond composites. TDTR measurement for this sandwich-structured object can be conducted, and the obtained results can be compared with the theoretical calculation. Therefore, the thermal boundary conductance and interface thermal conductance can be directly measured and compared to the theoretical value (TDTR system is shown in Fig. 28). + +As we explained before, the thermal boundary conductance is very important for the thermal conductivity in the copper/diamond composites. The recent published research works indicate that this + +problem can be dissociated to investigate by measuring TBC between physically simulated layers. Nevertheless, the measured TBC still has some discrepancies compared to the calculated values by AMM and DMM models. + +# Thermal conductance of the interface layer + +Apart from the thermal boundary conductance, the thermal conductance of the interface layer, $\mathrm { K _ { d } }$ (as indicated in Fig. 19), is another important part to determine the thermal conductivity of the fabricated materials, and it is affected by the following three factors. + +(A) Dislocations From the experimental data listed in Tables 1, 2, 3 and 4in ‘‘Fabrication and thermal conductivity of copper/diamond composites’’ section, it shows that the thickness of the interface layer in copper/diamond composites ranges from a few hundred nm up to 2000 nm. In most cases, the mean free path (MFP) of heat carriers (phonons or electrons) at room temperature is quite small, in the order of \ 1–3 nm, and this value is far smaller than the thickness of the interface layer. This means that the thermal transfer in the interface layer of the copper/diamond composites is exactly similar to that transferring in the bulk materials, which has been analyzed in ‘‘In copper matrix’’ section. + +Researchers investigate the quantified relationship between the crystal dislocation and thermal conductance of the interfacial layer, and the results are shown in Fig. 29 [109]. They find that two orders of magnitude of increase in dislocation densities lead to a factor of two reduction in thermal conductance, demonstrating the detrimental effect of dislocations on the thermal conductance. However, it is hard to directly characterize the dislocation density, making + +![](images/313a7c15f3263be3f388b845fce78ad473b5cc3bbd53d7dd336e3d25bb5f2da0.jpg) +Figure 27 Schematics of the TDTR test of thermal conductance between the TiC layer and diamond substrate [80]. + +![](images/08e745f06fae496a445aeaad7e32636f1f5219d2903d560f494afa64a271a082.jpg) +Figure 28 Schematic diagram of key components of a TDTR system [87]. + +it even harder to establish quantitative forms of functions to describe the dependency of thermal boundary conductance on dislocation density. + +(B) The additional interfacial layer If the interface structure is composed of two layers and the thickness for one of the layers is in nano-size, the wave nature of the phonon needs to be considered when the MFP of heat carriers is comparable or larger than the nanosized interface layer. Researchers study the effect of the additional Cu layer, with a thickness of ranging from a few nanometers to 30 nm, on the thermal boundary conductance between the Au and the substrate [113], and find that the TBC in the structure of ${ \mathrm { A u { \mathrm { - } C u { \mathrm { - } A l _ { 2 } O _ { 3 } } } } }$ is increased with increasing the thickness of inserted Cu layer, as shown in Fig. 30. The improvement in TBC is attributed to the electron–phonon coupling. However, similar research work has been done by Jeong et $\mathsf { a l . , }$ and they claim the additional phonon modes provided by the thin Cu layer are the dominant mechanism, not electron– phonon coupling, for improving TBC [49], and the improved TBC is verified by both experimental measurement and DMM prediction. + +(C) Intrinsic property and structure of the interface layer The intrinsic properties of the interface layer can affect the phonon transmission from the two sides of the interface layer. The Debye temperature of the materials that are sitting on the two sides of the interface layer is usually different. According to the assumptions in DMM and AMM, the phonons of the material that is sitting on one side of the interface layer can only couple with the phonons that have + +![](images/19eedd23dff1581ef88dd96faa31b891511cd7318ee330fd43c8123029cd4cc9.jpg) +Figure 29 Thermal boundary conductance at the GaSb/GaAs interfaces with two different dislocations density from the work [111]. TiN/MgO and Bi/diamond data for comparison come from work [112] and [29]. The figure is from [109]. + +identical frequency from the material sit on the other side of the interface layer during transmission process. Large interface vibrational mismatch of the materials on the two sides leads to that very few + +![](images/31ad8c159ca7da6816e31653a2ae4086c03b9c655de35fca588f226c41d68897.jpg) +Figure 30 Influence of the addition of a nanometer-sized copper interlayer on the thermal boundary conductance at the Au/ ${ \bf A } _ { 2 } \mathrm { O } _ { 3 }$ interface. All data and calculations are from Ref [113]. + +phonon modes can transmit across the interface layer, thus resulting in low interface thermal conductance. If the materials that are sitting on the two sides of the interface layer have a Debye temperature ratio of 5–20, the obtained G will be in the range of 8–30 M/ W m2 K [29]. The density of state (DOS) of the materials can be used to explain this. In Fig. 31, the DOS of the two materials has little overlapped area so that the thermal vibration from one side of the interface layer cannot stimulate the vibration of the material on the other side. To mediate the discrepancy, a layer with required properties could be inserted between the two materials. For example, the inserted film leads to both materials have a more overlapped DOS area, as indicated by the hatched region in Fig. 32, and this facilitates more phonon transmission. This method has been experimentally applied and demonstrated in copper/diamond composites [80, 114]. + +The structure of the interface layer can also tune the phonon transmission across the interface. It is possible to design a certain structure to enhance the interface thermal conductance. Insertion of one layer or two layers with mediated Debye temperature aforementioned is a possible method. Is it possible to make a more complex and multiple-layered structure to boost the phonon transmission further? The investigations of the superlattice may shine some + +light on this problem, even though the topic is toward the reduction in its thermal conductivity, which is opposite to the purpose of this review. Superlattice is a structure with two different materials repeating periodically. Phonon as a kind of quasi-particle has wave–particle duality [117]. When the phonon MFP is smaller than the periodic thickness, it transmits in its particle nature (the descriptions of AMM and DMM are based on its particle nature). When the phonon MFP is comparable or larger than the periodic thickness, its wave nature must be considered. For example, in the incoherent regime in Fig. 33, the particle nature of phonons dominates because the interface density is small (periodic thickness is large or compatible with MFP) [118]. With the interface density increasing, more interface thermal resistance appears, and this leads to the decreased thermal conductivity as shown in Fig. 33. The minimum TC corresponding to a periodic thickness shown in the figure is caused by the phonon band folding [119, 120]. After this critical TC, the phonons’ wave nature dominates in the coherent regime. With higher interface density (much smaller periodic thickness, compared to MFP), phonons can transport ballistically before being reflected at the interface or getting scattered, and this leads to the increased thermal conductivity. The key point that needs to be noticed is that, theoretically, the thermal conductivity of superlattice could surpass the original TC when the interface density is high enough (as shown in Fig. 33). This may enlighten that the additional channels of phonon waves could be utilized by ballistic transport or phonon interference to enhance the interface thermal conductance further if the materials have appropriate periodic thickness and MFP. + +# Interface thermal conductance calculation + +Interface thermal conductance (ITC) is composed of two parts, TBC and thermal conductance of the interface layer, and the affecting factors on those two parts have been reviewed in ‘‘Thermal boundary conductance’’ and ‘‘Thermal conductance of the interface layer’’ sections, respectively; however, an equation needs to be established for easily evaluating the thermal conductance of the entire interfacial system. Recalling the definition of thermal conductance and expression in Eq. (2), it is comparable to the concept of electric resistance, which is the voltage per electric current, and it is temperature per heat current + +![](images/5fe81ab7dd7531a92c2a23fa4bc062e070925c5688ffc54a6c93dab8c0519668.jpg) +Figure 31 Illustration of the phonon transmission probability as predicted by the Diffuse mismatch model. Given two interface solids, 1 and 2, with vibrational spectra, $D 1 ( \omega )$ and $D 2 ( \omega )$ , the phonon transmission probability from 1 to $2 , \ \operatorname { T r } _ { 1 - 2 } ( ( \omega ) )$ is dependent on the ratio of the spectra as shown in the figure. + +![](images/9ae59c1face50cb3e148391c92dc46eda9a4488c80ea70ccb3deae654c2dce3b.jpg) +Figure 32 DOS along the z-axis calculated from monolayers containing 72 atoms of material A and C in the enhanced interface (dashed lines) in comparison with the baseline interface [116]. + +for the thermal resistance. The definition of thermal resistance is the inverse of thermal conductance, vice versa. + +Hence, we assume that the model in Fig. 19 is the structure of the interface; the expression of interface thermal conductance can be formulated according to the principle in electrical resistance in series law as: + +$$ +\frac {1}{h _ {\mathrm {c}}} = \frac {1}{G _ {1}} + \frac {a}{K _ {i}} + \frac {1}{G _ {2}} \tag {15} +$$ + +Representative locations of maxima and minima in the transmission probability are labeled on D1(x), $D 2 ( \omega ) ,$ and $\operatorname { T r } _ { 1 - 2 } ( \omega )$ . The transmission probability at $\omega _ { 3 }$ is zero because solid 2 does not vibrate at that frequency [115]. + +![](images/0756dcedcde8d2b5f34220301549f35404525b314f491382166c23d465a07adf.jpg) +Figure 33 Thermal conductivity of ${ } ^ { 1 2 } \mathrm { C } / { } ^ { 1 3 } \mathrm { C }$ graphene superlattices as a function of the interface density. The red hollow circles represent the cases with total length of 240 $\mathbf { A } ^ { \circ } { \mathrm { : } }$ , and the blue solid squares represent those with total length of 480 $\mathbf { A } ^ { \circ } .$ . The black vertical dash line indicates the position of crossover from incoherent to coherent phonon transport [121]. + +where $h _ { \mathrm { c } }$ is the interface thermal conductance; $G _ { 1 }$ and $G _ { 2 }$ represent thermal boundary conductance of the boundaries at both sides of the interface layer. $K _ { i }$ represents the thermal conductivity value of the materials composed of an interface; $" a "$ represents the thickness of interface layer. Thus, the value of $G _ { 1 } , G _ { 2 }$ and $K _ { i }$ should be considered separately based on the above discussion and review to accurately quantify each value. + +The interface thermal conductance problems should be analyzed from every aspect of the interface system. The complexity in this part lies in handling these factors together in the perspective of phonon as heat carriers mathematically. Especially for thermal boundary conductance, after nearly six decades’ research, the underlying mechanisms of the phenomenon are still elusive [27]. The copper/diamond composites provide us the intricate interfacial systems to delve into this interesting problem, and probably these materials can be used to verify the theory from the experimental results. + +# Model prediction of effective thermal conductivity of copper/diamond composites + +We already reviewed some research results of the synthesized copper/diamond composites and factors that influence the thermal transfer efficiency of the composites in the above sections. Some of the factors are described in the form of functions to quantify their influence on the thermal transfer process. Hence, here we review calculation models for predicting the effective thermal conductivity of heterogeneous materials. Before applying the models to copper/diamond composites, we assume that all factors mentioned above are taken into account; for example, we assume that the interface thermal conductance and thermal conductivity values of a copper matrix and diamond are measured or calculated as discussed in ‘‘Factors affecting the thermal conductivity of copper/diamond composites’’ section. + +# Maxwell model + +Due to the similar mathematical nature of electric current and heat current transmitting in materials, the principles of Maxwell’s findings on electric are applied in heat transfer problems. Maxwell is the first to give expressions for the effective electric conductivity of heterogeneous materials in his famous work on electricity and magnetism [122]. The formula is given by replacing the concept of electric conductivity with thermal conductivity: + +$$ +\frac {k _ {\mathrm {c}}}{k _ {\mathrm {m}}} = 1 + \frac {3 V _ {\mathrm {d}}}{\left(\frac {k _ {\mathrm {d}} + 2 k _ {\mathrm {m}}}{k _ {\mathrm {d}} - k _ {\mathrm {m}}}\right) - V _ {\mathrm {d}}} \tag {16} +$$ + +where $k _ { \mathrm { c } } , k _ { \mathrm { m } }$ and $k _ { \mathrm { d } }$ are the thermal conductivity of composite, matrix, and dispersed reinforcement particles, respectively. $V _ { \mathrm { d } }$ represents the volume percentage of reinforcement particles. Maxwell deduces the expression based on the assumption that the dispersed particles with spherical shape are diluted embedded in a continuous matrix and non-interacting. This indicates that the Maxwell model is accurate only for low volume percentage of reinforced particles and the situation of reinforced particles contacting each other is absent. Also, interfacial resistance and size of the reinforced particles are not considered in this model. Since the shape of diamond particles is cuboctahedron and is very close to the sphere, it is appropriate to apply this model to copper/diamond composites. When applying this model to copper/diamond composites, as the interface thermal resistance between the copper matrix and the diamond particles cannot be ignored, large discrepancy between experimental value and model prediction is obvious [20, 24, 40]. One research work shows the large discrepancy between experimental and calculated values when applying the Maxwell model, as shown in Fig. 34 [20]. The blue dot in the picture, corresponding to the composites with uncoated diamond, shows an opposite tendency in contrast with the modeled results. This is because the model is developed under the assumption that no pore or gaps are existed in the matrix or near the interface area. However, with uncoated diamond, the gaps probably appear around the diamond particle reinforcements due to the poor chemical affinity between the diamond and the copper. + +# Hasselman–Johnson model + +Hasselman and Johnson considered the interface thermal conductance among the reinforced particles and the size effect of reinforced particles, they deduce the Hasselman–Johnson model (H–J model) based on the work of Maxwell, and the deduction details can be found in Ref [123]. The expression is given: + +$$ +k _ {\mathrm {c}} = k _ {\mathrm {m}} \frac {\left[ 2 \left(\frac {K _ {\mathrm {d}}}{K _ {\mathrm {m}}} - \frac {K _ {\mathrm {d}}}{a h _ {\mathrm {c}}} - 1\right) V _ {\mathrm {d}} + \frac {K _ {\mathrm {d}}}{K _ {\mathrm {m}}} + \frac {2 K _ {\mathrm {d}}}{a h _ {\mathrm {c}}} + 2 \right]}{\left[ \left(1 - \frac {K _ {\mathrm {d}}}{K _ {\mathrm {m}}} + \frac {K _ {\mathrm {d}}}{a h _ {\mathrm {c}}}\right) V _ {\mathrm {d}} + \frac {K _ {\mathrm {d}}}{K _ {\mathrm {m}}} + \frac {2 K _ {\mathrm {d}}}{a h _ {\mathrm {c}}} + 2 \right]} \tag {17} +$$ + +where $" a "$ is the radius of the dispersed particles; $h _ { \mathrm { c } }$ is the interface thermal conductance. It can also be written in this form: + +![](images/bbbf62d473bee90a82c049e2099848619ed1effdd067d6bb192f334685981a4f.jpg) +Figure 34 Thermal conductivity of copper–diamond composites with uncoated and Mo2C-coated diamond particles [20]. + +$$ +k _ {\mathrm {c}} = k _ {\mathrm {m}} \frac {\frac {k _ {\mathrm {d}} ^ {\text {e f f}}}{K _ {\mathrm {m}}} \left(1 + 2 V _ {\mathrm {d}}\right) + \left(2 - 2 V _ {\mathrm {d}}\right)}{\frac {\kappa_ {\mathrm {d}} ^ {\text {e f f}}}{\kappa_ {\mathrm {m}}} \left(1 - V _ {\mathrm {d}}\right) + \left(2 + V _ {\mathrm {d}}\right)} \tag {18} +$$ + +the meaning of $k _ { \mathrm { d } } ^ { \mathrm { e f f } }$ is given in $\operatorname { E q . }$ (20). + +Since the H–J model considers the interface thermal resistance and the reinforced particle sizes, it has better prediction results than the Maxwell model. However, the model still has the same restriction as the Maxwell model, and it is only accurate for the composite with a low volume percentage of reinforcement. Due to the simplicity of the H–J model, it has been applied in many research works for predicting copper/diamond composites’ thermal conductivity [21, 32, 48, 124]. The H–J model can basically reflect the tendency of diamond volume percentage versus thermal conductivity of the copper/diamond composites after considering the ITC, as shown in Fig. 35 [41], although the discrepancy between the experimental value and prediction is very obvious. + +# Differential effective medium (DEM) model + +DEM model is an improved version of the former ones, which can be deduced directly from the H–J model; the threads and details can be found in Refs [125, 126]. The expressions of DEM are: + +$$ +\left(1 - V _ {\mathrm {d}}\right) \left(\frac {k _ {\mathrm {c}}}{k _ {\mathrm {m}}}\right) ^ {\frac {1}{3}} = \frac {k _ {\mathrm {d}} ^ {\text {e f f}} - k _ {\mathrm {c}}}{k _ {\mathrm {d}} ^ {\text {e f f}} - k _ {\mathrm {m}}} \tag {19} +$$ + +$$ +k _ {\mathrm {d}} ^ {\text {e f f}} = \frac {k _ {\mathrm {d}}}{1 + \frac {k _ {\mathrm {d}}}{a h _ {\mathrm {c}}}} \tag {20} +$$ + +The meanings of symbols in these equations are the same as those in Maxwell and H–J models. Due to its concise form and simplicity, the DEM model is applied in many works related to copper/diamond composite materials [6, 7, 19, 87, 127]. Yang et al. find that the Cu/diamond composite with Ti additive can reach 99% of the predicted value (555 W/m K) by the DEM model [62]. In another work, the DEM model can give an upper bound for the prediction of thermal conductivity of copper/diamond composites, based on the TDTR-measured TBC value $\mathrm { G } ,$ as shown in Fig. 36 [87]. + +DEM model is more stable for fitting the experimental data when changing the diamond size or volume percentage [125]. Overall, the DEM model is quite accurate for the prediction of the thermal conductivity of the composites. The deviation from the experimental value could be caused by the geometry of the diamond and the interface thermal conductance. + +# Comparison of the prediction models + +The sequence of the three models introduced above is that: the latter model is evolved and modified from the former one. Each of the three models has its own applications, which can be found in most of the research works we reviewed in ‘‘Fabrication and thermal conductivity of copper/diamond + +![](images/186b39ec6bc39a32411fb12e80b0d10d49f39ba06d48bb05c729d832a727a652.jpg) +Figure 35 Theoretical curves and experimental values of the thermal conductivity of the composites fabricated using the diamonds coated by 1-lm $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ or 1-lm TiC layer [41]. + +![](images/435d72d7b6a48e3350ec18f4e4ca00b21f69bfa3c4e30f06fd09693980781684.jpg) +Figure 36 Relation between G of Cu/diamond interface and k of Cu/diamond composite established by the DEM model [87]. + +composites’’ section. The classic Maxwell model is easy to be used for predicting the upper limit of thermal conductivity value of copper/diamond composites with a certain volume fraction of diamonds, because the most important factor, interface thermal conductance, in this model $h _ { \mathrm { c } } = \infty$ . In the H– J model, ITC and diamond sizes are considered; the accuracy for prediction is improved compared to the Maxwell model. For the DEM model, a new term $k _ { \mathrm { d } } ^ { \mathrm { e f f } }$ is introduced, which can be understood that the reinforced particle with the size of $" a "$ is considered as a new reinforced particle (a unity) that has an effective thermal conductivity $k _ { \mathrm { d } } ^ { \mathrm { e f f } }$ with ITC $^ { \prime \prime } h ^ { \prime \prime } = \infty$ . This is also reflected in Eq. (18). Thus, DEM can also be used to estimate the interface thermal conductance by inversely considering this issue. The steps are: (1) calculating the $k _ { \mathrm { d } } ^ { \mathrm { e f f } }$ with known $k _ { \mathrm { c } }$ by experimental measurement based on Eq. (19); (2) rewriting Eq. (20) to: + +$$ +\frac {1}{k _ {\mathrm {d}} ^ {\text {e f f}}} = \frac {1}{k _ {\mathrm {d}}} + \frac {1}{h _ {\mathrm {c}}} \frac {1}{a} \tag {21} +$$ + +and then plotting $( k _ { \mathrm { d } } ^ { \mathrm { e f f } } )$ -1 against a-1 to yield the line with slope $h _ { \mathrm { c } } ^ { - 1 }$ , to obtain the ITC. The estimation of uncertainty by this inversed method can be found in Ref [28]. + +# Conclusions and outlook + +Copper/diamond composites have great potential to lead the next generation of heat sink materials for high-power electronic devices use, with potential + +thermal conductivity [ 500 W/(m K) and a thermal expansion coefficient that is tunable to match with that of chip materials, in the range of 4–6 ppm/K. However, their potentially high thermal conductivity is strongly dependent on the quality of the copper/diamond interface, which is disrupted by the poor chemical affinity between the copper and the diamond. The interface layer acts as a ‘‘bridge’’ to improve the efficiency of thermal transferring between the diamond particles and the copper matrix. The interface layer’s characteristics determine the interface thermal conductance that is crucial for having high thermal conductivity for the copper/diamond composites, and it is also vital for obtaining high conductive copper/diamond composite to require the composite has a matrix with low defect level and high-quality diamond particle reinforcement. Through the understanding on how heat carries across the interface layer boundary and transfers in the interlayer structure from physical perspectives, we have identified some factors, such as the boundary surface roughness, atomic intermixing, interfacial bonding strength, contrast of the DOS, and the structure of interfacial layers that could significantly influence the interfacial layer’s thermal boundary conductance and thermal conductance, thus further affecting the thermal conductivity of copper/diamond composites. To engineer a copper/diamond composite for meeting the thermal conductivity requirement of practical applications, we need to have an in-depth understanding on how the interfacial layer’s characteristics influence the heat transfer from the diamond to copper, establish a quantitative relationship between those factors and the resultant thermophysical properties, and expand the relationship including the processing, eventually establishing the processing-interfacial layer characteristics (microstructures)–thermophysical properties relationship for the copper/diamond composites. This understanding is lacking or not unified; currently, more research work needs to be carried in this regard. Furthermore, diamond is very hard and the composite’s machinability becomes very poor when diamond particles are integrated into the copper matrix, making it very challenging to machine copper/diamond composites into specific products for practical use. Near-net shape forming technology for producing copper/diamond composites is necessary, and this would be a game changer that enables us to achieve cost-effective production of copper/diamond + +products with specific geometry requirements and bring the copper/diamond products to a wide market for high-power thermal management requirements. The potential near-net shaping technology for producing copper/diamond composites could be but not limited to additive manufacturing such as selective laser melting, direct energy deposition, or binder jetting. + +# Acknowledgement + +This work was supported by the US Air Force Office of Scientific Research under Award Number FA2386- 17-1-4025. + +Open Access This article is licensed under a Creative Commons Attribution 4.0 International License, which permits use, sharing, adaptation, distribution and reproduction in any medium or format, as long as you give appropriate credit to the original author(s) and the source, provide a link to the Creative Commons licence, and indicate if changes were made. The images or other third party material in this article are included in the article’s Creative Commons licence, unless indicated otherwise in a credit line to the material. 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of Cu-diamond composites with Ti-coated diamond particles produced by gas pressure infiltration/11_Microstructure and thermal conductivity of Cu-diamond composites with Ti-coated diamond particles produced by gas pressure infiltration.md @@ -0,0 +1,269 @@ +# Accepted Manuscript + +Microstructure and thermal conductivity of Cu/diamond composites with Ti-coated diamond particles produced by gas pressure infiltration + +Jianwei Li, Hailong Zhang, Yang Zhang, Zifan Che, Xitao Wang + +![](images/732b133d09324abc692504623b4c1f0025d9ebbaf494584cab42e91008552d37.jpg) + +PII: S0925-8388(15)30178-X + +DOI: 10.1016/j.jallcom.2015.06.062 + +Reference: JALCOM 34404 + +To appear in: Journal of Alloys and Compounds + +Received Date: 10 January 2015 + +Revised Date: 4 June 2015 + +Accepted Date: 7 June 2015 + +Please cite this article as: J. Li, H. Zhang, Y. Zhang, Z. Che, X. Wang, Microstructure and thermal conductivity of Cu/diamond composites with Ti-coated diamond particles produced by gas pressure infiltration, Journal of Alloys and Compounds (2015), doi: 10.1016/j.jallcom.2015.06.062. + +This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. + +# Graphical Abstract + +![](images/bf6a3770143f62ad8d7f72ca21f96557636a7beb1f6552b7095c641336f30870.jpg) + +![](images/9c091c68b26c4998ede85bba13a4d76d0f34bc4c7e2df55c99d30576642caa6d.jpg) + +![](images/7824905a6cc0bd415219e886bcc1b096a02631034eb72d44afdf1ea3b295de43.jpg) + +# Microstructure and thermal conductivity of Cu/diamond composites with Ti-coated diamond particles produced by gas pressure infiltration + +Jianwei Li, Hailong Zhang, Yang Zhang, Zifan Che, Xitao Wang * + +State Key Laboratory for Advanced Metals and Materials, University of Science and + +Technology Beijing, Beijing 100083, China + +* Corresponding author. Tel.: +86-10-82375280; Fax: +86-10-62333447. + +Email: xtwang.ustb@gmail.com + +Abstract: As an attractive thermal management material, diamond particles reinforced Cu matrix (Cu/diamond) composites generally exhibit thermal conductivities lower than expected. To exploit the potential of heat conduction, a combination of Ti coating on diamond particles and gas pressure infiltration was used to prepare Cu/diamond(Ti) composites. A high thermal conductivity of 716 W/mK and a low coefficient of thermal expansion of 5.8 ppm/K at 323 K were obtained in the composites. Auger electron spectroscopy (AES) characterization shows that a TiC layer was formed between Cu matrix and diamond reinforcement, which is responsible for the enhancement of thermal conductivity. The results suggest that Ti coating can significantly promote interface bonding between Cu and diamond and gas pressure infiltration is an effective method to produce Cu/diamond composites. + +Keywords: A. Metal matrix composites; B. Liquid-solid reactions; C. Heat conduction; Auger electron spectroscopy + +# 1. Introduction + +Microelectronic devices with high power density are demanding electronic packaging materials with high thermal conductivity (TC). Due to extraordinarily high TC of diamond (1200–2200 W/mK), diamond particles reinforced metal matrix composites (MMCs) become a popular candidate and are inspiring great interests. Compared with Al/diamond [1,2], Cu/diamond composites have higher TC and lower coefficient of thermal expansion (CTE) in theory and are especially suitable for applications at higher operating temperatures due to higher melting point of Cu [3]. Microelectronic devices can generate significant heat fluxes such as ${ \sim } 5 0 ~ \mathrm { W / c m ^ { 2 } }$ in current electronic chips and up to $2 0 0 0 \ \mathrm { W / c m } ^ { 2 }$ in semiconductor lasers [4]. Owing to high thermal conductivity, Cu/diamond composites as thermal conductive materials can dissipate heat from the microelectronic devices quickly and maintain reasonable temperature in the devices. + +Despite the potential for excellent performance in Cu/diamond composites, poor wettability of copper with diamond acts as a barrier to processing the composites. Even at a high temperature of 1673 K, copper and diamond exhibits a contact angle as large as 128.7° [5]. In order to take full advantage of high thermal conductivity of diamond, surface metallization of diamond particles [6-11] and alloying of Cu matrix [12-15] are commonly used to improve interfacial bonding between Cu matrix and diamond reinforcements. Zhang et al. obtained an enhanced TC value of 493 W/mK by coating Ti on diamond particles [16]. He et al. achieved a high TC value of 677 W/mK in Cu/diamond composites by alloying Zr into Cu matrix [17]. + +In addition to interface modification, microstructure optimization of Cu/diamond composites is also crucial to the enhancement of thermal conductivity. So far several processing techniques like spark plasma sintering (SPS) [18-21], pulse plasma sintering [22], vacuum hot pressing [23], pressureless infiltration [9], high pressure high temperature (HPHT) infiltration [16] and gas pressure infiltration (GPI) [15] have been employed to optimize the microstructure of Cu/diamond composites. Among these processes, GPI is proved to be effective in achieving high TC in Cu/diamond composites [15]. Since Beffort et al. [24] firstly produced Al/diamond composites using GPI in 2006, this route has been followed by several groups [1,25]. However, because of high melting point of Cu and extreme chemical inertness between diamond and Cu, the production of Cu/diamond composites via GPI is rarely reported [15]. Besides, the effectiveness of Ti coating on diamond particles in enhancing thermal conductivity of GPI-produced Cu/diamond composites has not yet been proved. As an alternative route, the combination of surface metallization of diamond particles by Ti coating and promising technique of GPI could be helpful to modify Cu/diamond interface and to achieve high thermal conductivity in the composites. + +In this article, we use gas pressure infiltration to produce Cu/diamond composites with Ti-coated diamond particles. The interface structure is elaborately investigated by using diamond particles dissolved from the produced Cu/diamond composites to determine the effect of Ti coating on thermal conductivity. As a result, a TC of 716 W/mK is achieved, which is 79% higher than 400 W/mK for Cu matrix. + +# 2. Experimental + +Commercially available bulk Cu (purity: 99.95 wt.%, Cuibolin Non-Ferrous Technology Co., China) was used as metal matrix. Synthetic single-crystalline diamond powders (particle size: 62-75 µm, mesh: 200-230, HWD40, Henan Huanghe Whirlwind International Co., China) were used as particulate reinforcements. The diamond has been characterized to contain 150-170 ppm $\mathrm { N } _ { 2 } ,$ giving a TC of 1700-1850 W/mK [26]. The diamond particles were coated with Ti by using vacuum vapor deposition. In doing so, diamond particles and Ti powders (particle size: ~5 µm, TaiYu Materials Science & Technology Co., China) were mixed and then heated at 1098 K for 15 min in a high vacuum induction- heating furnace. + +The Cu/diamond(Ti) composites were produced by a home-made gas pressure infiltration equipment. The Ti-coated diamond particles were tap-packed in a graphite mold and the Cu bulks were placed on top of the diamond powder bed. The installed graphite mold was moved to a graphite heating muff for infiltration. A vacuum better than 0.1 Pa was evacuated before heating, and the graphite mold was then inductively heated to 1423 K and kept for 10 min. Afterwards, high-purity argon gas was pumped into the muff to maintain a pressure of 1.0 MPa. The pressure was kept for 15 min to enable sufficient contact of diamond crystals with copper melts, before switching off the heating source. The infiltrated Cu/diamond(Ti) sample was furnace cooled down to room temperature within 1-2 h. + +The volume fraction of diamonds in the composite can be estimated as follows. Without consideration of pores in samples, the composite density is written as $\rho _ { \mathrm { c } } = \rho _ { \mathrm { d } } V _ { \mathrm { d } }$ + +$+ \rho _ { \mathrm { m } } ( 1 { - } V _ { \mathrm { d } } )$ , where $\rho _ { \mathrm { c } }$ , ρd and $\rho _ { \mathrm { m } }$ are densities of composite, diamond and copper, + +respectively, and $V _ { \mathrm { d } }$ is the volume fraction of diamonds in the composite. For the + +estimation, $\rho _ { \mathrm { c } }$ was examined by the Archimedes, and $\rho _ { \mathrm { m } } = 8 . 9 0 ~ \mathrm { g / c m } ^ { 3 }$ and $\rho _ { \mathrm { d } } = 3 . 5 2$ + +$\mathrm { g } / \mathrm { c m } ^ { 3 }$ were used. The phase structure of the coated diamond particles was detected by + +X-ray diffraction using Cu Kα radiation (XRD, Rigaku DMAX-RB, Japan). Auger + +electron spectroscopy (AES, PHI 700, Japan) was used to investigate interface diffusion + +and reaction between Ti coating layer and diamond. The produced Cu/diamond + +composite samples were electrochemically etched in a 10 vol.% $\mathrm { H N O } _ { 3 }$ and 90 vol.% + +$_ { \mathrm { H } _ { 2 } \mathrm { O } }$ solution for 15 min. The extracted diamond particles were then cleaned in acetone + +for 1 min and the cleaning was repeated for 3 times. The microstructure of the + +Cu/diamond composites was characterized by field emission scanning electron + +microscopy (FE-SEM, SUPRA 55, ZEISS, Germany). The samples as fabricated by GPI + +were grinded by using diamond wheels and then SiC papers, and finally were polished + +by using diamond pastes. The thermal conductivities of the composites were evaluated + +by using $K = \alpha \rho _ { \mathrm { c } } C _ { \mathrm { p } } ,$ where K is thermal conductivity, α is thermal diffusivity, $\rho _ { \mathrm { c } }$ is bulk + +density of sample, and $C _ { \mathrm { p } }$ is specific heat capacity. The thermal diffusivity was + +measured at room temperature by a laser flash apparatus (LFA427/3/G, Netzsch, + +Germany). For the measurement of thermal diffusivity, the as-cast composite samples + +were cut into disk-shaped specimens 10 mm in diameter and 3 mm in thickness. The + +specific heat capacity was measured by a differential thermal analysis (DSC 204, + +Netzsch, Germany). The CTE measurements were conducted on a NETZSCH DIL + +402C thermo-mechanical analyzer at a heating rate of 5 K/min. The samples for CTE + +measurement were made with a cylinder shape of Ф3 mm×25 mm. + +# 3. Results and discussion + +The volume fraction of diamonds was estimated to be 65 vol.% by the density measurements. Fig. 1 shows the XRD patterns of the Ti-coated diamond particles. Besides diamond and TiC, a small amount of Ti phase was also detected with the diamond particles. This means that a layer configuration of diamond/TiC/Ti could be formed in the vacuum vapor deposition. + +The polished surfaces of the Cu/diamond composites with uncoated and Ti-coated diamond particles are shown in Fig. 2. The diamond particles were homogeneously dispersed in the Cu matrix. Due to poor wettability between pure copper and diamond, some uncoated diamond particles were found to dropout from the composite during the polishing process (Fig. 2a). By contrast, the presence of TiC layer strengthens the Cu/diamond(Ti) interface, and no diamond particles were found to fall off, as seen in Fig. 2b. It is thus concluded that the formed TiC layer can improve adhesion between diamond particles and Cu matrix. + +Fig. 3 shows the SEM observation of the fracture surfaces of the composites. The intact diamond particles shown in Fig. 3a indicate poor interfacial bonding in the unmodified Cu/diamond composite because the fracture often occurs at interfaces with weak bonding. For comparison, good interfacial bonding is revealed in the Cu/diamond composites with Ti-coated diamond particles, where plastic fractures occur within Cu matrix and Cu-adhered diamond particles are observed, as indicated by the circles in Fig. + +3b. This indicates that the interfacial bonding strength of Cu/diamond(Ti) is even higher than the intrinsic fracture strength of Cu matrix. + +The energy dispersive X-ray spectroscopy (EDS) scan across a Cu/diamond(Ti) interface is shown in Fig. 4. The EDS analysis was carried out on a grinded and polished surface. The result indicates that the intermediate zone, between Cu matrix and diamond reinforcements, was enriched in Ti. In addition, Ti signal detected in the matrix was higher than that on the diamond. This suggests that the coated Ti could diffuse to Cu matrix at the infiltration temperature of 1423 K. According to the Cu-Ti phase diagram, the maximum solubility of Ti in Cu is nearly 8 at.%. After cooling down, the dissolved Ti could stay in the crystals of Cu matrix. + +Fig. 5 shows the AES results for the Ti-coated diamond particles before infiltration and dissolved from the produced Cu/diamond(Ti) composites. The formation of TiC can be confirmed by the Auger line shape analysis. As shown in Fig. 5c and 5d, C spectra consisted of three peaks at 256, 264 and 274 eV; as shown in Fig. 5e and 5f, Ti spectra exhibited one peak at 419 eV. The analysis suggests the existence of TiC [16]. The initial thickness of TiC layer is 532 nm (Fig. 5a), deduced from the sputtering time with a sputtering rate of 76 nm/min that is calibrated by $\mathrm { S i O } _ { 2 }$ film. For comparison, the thickness of TiC layer on the diamond particles dissolved from the Cu/diamond(Ti) composites increases to 1292 nm (Fig. 5b). The finding means that C and Ti inter-diffuse during infiltration to increase the thickness of TiC. Fig. 5b shows that oxygen was presented through the layer. This indicates that environmental O could be incorporated into the layer during infiltration owing to the limited vacuum used. + +It is interesting to note that the chemical states of Ti and C on the surface of diamond particles have changed after infiltration. On the surface of the diamond particles before infiltration, the chemical states of Ti and C corresponds to TiC (Fig. 5c and 5e). In contrast, on the surface of the diamond particles dissolved from the composites, only one C peak was observed in the spectra (Fig. 5d) and the Ti peak almost disappeared (Fig. 5f). The difference could be explained as follows. In the electrochemical etching process to extract diamond particles, the Cu/diamond composites were designated as the anode. Besides the Cu matrix, the Ti in TiC is also electrochemically dissolved, leaving activated carbon on the surface of diamond particles [27]. As seen in Fig. 5b, the elemental profile of Ti exhibits an upward trend at the beginning of Ar+ sputtering, which provides an evidence to support the explanation. + +Fig. 6 shows a survey of thermal conductivity of Cu/diamond composites. The GPI-produced Cu/diamond composite reinforced with uncoated diamond particles showed a TC of only 134 W/mK, lower than half value of pure copper. Ti coating on diamond particles can greatly enhance the thermal conductivity of Cu/diamond composites. The TC of the Cu/diamond composites prepared by SPS and pressure infiltration (PI) are 493 W/mK [16] and 514 W/mK [28], respectively. For comparison, the TC of the Cu/diamond(Ti) composites prepared by GPI reached 716 W/mK. This suggests that TiC layer can improve the interface bonding effectively and then enhance the thermal conductivity. + +We have calculated the theoretical value of thermal conductivity for Cu/diamond composites using the Maxwell-Eucken model [29] without considering interface + +thermal resistance: + +$$ +K _ {c} = K _ {m} \left[ \frac {2 + K _ {d} / K _ {m} + 2 V _ {d} \left(K _ {d} / K _ {m} - 1\right)}{2 + K _ {d} / K _ {m} + V _ {d} \left(1 - K _ {d} / K _ {m}\right)} \right] \tag {1} +$$ + +where $K _ { \mathrm { c } } , K _ { \mathrm { m } }$ and $K _ { \mathrm { d } }$ are thermal conductivities of composite, matrix and particulate reinforcement, respectively, and $V _ { \mathrm { d } }$ is volume fraction of reinforcements. The calculations take thermal conductivities of 1800 and 400 W/mK for diamond reinforcement and Cu matrix, respectively, and $V _ { \mathrm { d } }$ of 65 vol.%. The calculated value is 1046 W/mK for the GPI-produced Cu/diamond composites, as shown in Fig. 6. It can be seen that the experimentally measured 716 W/mK is about 68% of the theoretical value. Since the thermal conductivity of TiC is 36.4 W/mK [30], much lower than those of both copper and diamond, the thermal conductivity of the composites is lower than the theoretical value. + +While metal/diamond composites are highly densified, acoustic impedance between dissimilar materials will be the main source of interface thermal resistance. In this study, Ti was chosen as a coating element for the metallization of diamond particles because TiC has an acoustic impendance between diamond and Cu. In order to understand directly the effect of Ti coating, we calculate the interface thermal conductance of the Cu/diamond composites. Without considering the microstructure of interface, the interface thermal conductance can be deduced from the acoustic mismatch model (AMM) [31]. In a previous study [16], we obtain the interface thermal conductance of $h _ { \mathrm { c o p p e r / d i a m o n d } } { = } 4 . 8 0 { \times } 1 0 ^ { 7 } \ \mathrm { W / m ^ { 2 } K }$ . Hasselman and Johnson [29] have proposed a theoretical model to predict the thermal conductivity of metal/diamond composites with consideration of interface thermal conductance: + +$$ +K _ {c} = K _ {m} \left[ \frac {2 \left(\frac {K _ {d}}{K _ {m}} - \frac {K _ {d}}{a h _ {c}} - 1\right) V _ {d} + \frac {K _ {d}}{K _ {m}} + \frac {2 K _ {d}}{a h _ {c}} + 2}{\left(1 - \frac {K _ {d}}{K _ {m}} + \frac {K _ {d}}{a h _ {c}}\right) V _ {d} + \frac {K _ {d}}{K _ {m}} + \frac {2 K _ {d}}{a h _ {c}} + 2} \right] \tag {2} +$$ + +where a is the radius of spherical dispersion and $h _ { \mathrm { c } }$ is the interface thermal conductance. When the values of other parameters are known, $h _ { \mathrm { c } }$ can be derived from Eq. (2). Then, the derived $h _ { \mathrm { c } }$ value is regarded as a measure of interface thermal conductance. The values of $K _ { \mathrm { c } } , K _ { \mathrm { m } } , K _ { \mathrm { d } }$ and $V _ { \mathrm { d } }$ are previously given, and the diamond particles have an average size of $a = 7 0 \mu \mathrm { m }$ . The $h _ { \mathrm { c } }$ values are thus calculated to be $3 . 3 6 { \times } 1 0 ^ { 5 } \mathrm { W / m } ^ { 2 } \mathrm { K }$ and $2 . 9 9 { \times } 1 0 ^ { 7 } \mathrm { W } / \mathrm { m } ^ { 2 } \mathrm { K }$ for the Cu/diamond composites with uncoated and coated diamond particles, respectively. It can be seen that Ti coating has greatly increased the interface thermal conductance between Cu matrix and diamond reinforcements, and in turn enhanced the thermal conductivity of the composites. Even so, the $2 . 9 9 { \times } 1 0 ^ { 7 } \mathrm { W } / \mathrm { m } ^ { 2 } \mathrm { K }$ is still lower than the value of $4 . 8 0 { \times } 1 0 ^ { 7 } \mathrm { W / m } ^ { 2 } \mathrm { K }$ for an ideal Cu/diamond interface. + +Actually, the thickness of TiC layer should also be taken into account in the calculation of interface thermal conductance because TiC has a low thermal conductivity. To do so, a model of series electrical resistance analogy [32] is used to calculate the interface thermal conductance of Cu/diamond composites: + +$$ +\frac {1}{h _ {\text {t o t a l}}} = \frac {1}{h _ {\text {c o p p e r / c a r b i d e}}} + \frac {1}{h _ {\text {c a i b i d e / d i a m o n d}}} + \frac {l}{K _ {\text {c a r b i d e}}} \tag {3} +$$ + +where h is interface thermal conductance, K is thermal conductivity, and l is thickness of carbide layer. The thickness of TiC layer is characterized to be 1292 nm based on the AES analysis and the thermal conductivity of TiC is 36.4 W/mK [30]. The $h _ { \mathrm { c o p p e r / c a r b i d e } }$ and $h _ { \mathrm { c a r b i d e / d i a m o n d } }$ are deduced from the AMM model to be $2 . 1 { \times } 1 0 ^ { 8 } \mathrm { W } / \mathrm { m } ^ { 2 } \mathrm { K }$ and $5 . 9 \times 1 0 ^ { 8 }$ + +$\mathrm { w } / \mathrm { m } ^ { 2 } \mathrm { K }$ , respectively. When the carbide layer is taken into consideration, the interface thermal conductance for the Cu/diamond composites is thus evaluated to be $2 . 3 8 \times 1 0 ^ { 7 }$ $\mathrm { w } / \mathrm { m } ^ { 2 } \mathrm { K }$ , a bit lower than without consideration of TiC layer. This finding shows the negative impact of the coating. As such, the coating layer should be as thin as possible in order to improve the interface thermal conductance of Cu/diamond composites. + +In a thermal management system, the packaging materials should have good thermal conductivities. The other concern is the match of CTE between microelectronic devices and electronic packaging materials. Fig. 7 shows the CTE values of the Cu/diamond(Ti) composites. The CTE values were obtained from the instantaneous derivative of the sample relative dilatation vs. temperature. The CTE of the composites increases gradually with temperature. When the temperature is lower than 600 K, the CTE is below 9 ppm/K. At 323 K, the CTE is as low as 5.8 ppm/K, meeting the requirements of electronic packaging materials. For the cooling of personal computer (PC), the temperature of the central processing unit (CPU) should be restricted to 353 K-363 K [33]. As such, we have also obtained the average CTE value in the temperature range from 323 K to 373 K. The CTE value was obtained to be 6.0 ppm/K, by averaging all data points measured in the temperature range. This CTE value matches that of semiconductors. + +# 4. Conclusions + +The Cu/diamond composites with Ti-coated diamond particles were fabricated by gas pressure infiltration. The results suggest that Ti coating on diamond particles + +enhances the interfacial bonding between diamond and copper in the produced Cu/diamond(Ti) composites. As a result, the thermal conductivity of the composites is dramatically increased from 134 to 716 W/mK. The average CTE values is 6.0 ppm/K in the temperature range from 323 K to 373 K. However, the measured thermal conductivity is still lower than the theoretical value. In future research, the thickness of TiC layer could be optimized by adjusting vacuum vapor deposition or gas pressure infiltration parameter to further enhance the thermal conductivity. + +# Acknowledgements + +This work was financially supported by the International Science and Technology Cooperation Program of China (2014DFA51610), National Natural Science Foundation of China (51271017, 51301018) and Fundamental Research Funds for the Central Universities (No. FRF-TP-13-033A2). + +# References + +[1] I.E. Monje, E. Louis, J.M. Molina, Optimizing thermal conductivity in gas-pressure infiltrated aluminum/diamond composites by precise processing control, Composites Part A, 48 (2013) 9-14. +[2] Z. Tan, Z. Li, G. Fan, X. Kai, G. Ji, L. Zhang, D. Zhang, Diamond/aluminum composites processed by vacuum hot pressing: Microstructure characteristics and thermal properties, Diamond Relat. Mater., 31 (2013) 1-5. +[3] K. Yoshida, H. 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Vershinin, M.A. Korukov, Miniature loop heat pipes for electronics cooling, Appl. Therm. Eng., 23 (2003) 1125-1135. + +# Figure captions + +Fig. 1. XRD patterns of the Ti-coated diamond particles. +Fig. 2. Microstructure of the Cu/diamond composites fabricated by gas pressure infiltration with (a) uncoated and (b) Ti-coated diamond particles. +Fig. 3. Fracture surfaces of the Cu/diamond composites fabricated by gas pressure infiltration with (a) uncoated and (b) Ti-coated diamond particles. +Fig. 4. (a) Polished surface of the Cu/diamond(Ti) composites and elemental distribution of EDS scan across Cu/diamond interface for (b) Ti, (c) Cu, and (d) C. +Fig. 5. Auger electron spectroscopy detected elemental depth profile for Ti-coated diamond particles (a) before composite fabrication and (b) dissolved from the Cu/diamond(Ti) composites as well as the corresponding kinetic spectra of C and Ti at various depths: (c) C and (e) Ti for (a) and (d) C and (f) Ti for (b). +Fig. 6. Thermal conductivity of the Cu/diamond composites fabricated by gas pressure infiltration, compared with literature and theoretical value. +Fig. 7. Coefficient of thermal expansion of the Cu/diamond(Ti) composites fabricated by gas pressure infiltration. + +![](images/59053873ee2bb8f9a2df095064e8766c1da38f5a648d8acc3818e314064b295e.jpg) +Fig. 1. XRD patterns of the Ti-coated diamond particles. + +![](images/7b0be86aea1ef701bbebd238581a6ceac91bda81d238c94982d382a4c22245a5.jpg) +Fig. 2. Microstructure of the Cu/diamond composites fabricated by gas pressure infiltration with (a) uncoated and (b) Ti-coated diamond particles. + +![](images/242da818c0c77d29b2badc51f7836e255d3a85cd4cb2f498fea0b0569adc133f.jpg) +Fig. 3. Fracture surfaces of the Cu/diamond composites fabricated by gas pressure infiltration with (a) uncoated and (b) Ti-coated diamond particles. + +![](images/20245d1f51ec89976216b4b6f0dc88f37655fc5b407b71a7d0ff957740c1a3a9.jpg) + +![](images/77f611caf0cd926d45673ce2ed6b4a2d9c5489d62c469a623be70d2d72bb114d.jpg) + +![](images/e9d4cc23e4ae65fbe8c3bff4026cc76f68158e63deb018c45d4767f86257c4bb.jpg) + +![](images/671845b5217a41e73998cc064a7adf0061409d23569cafa241f3161b847b1207.jpg) +C Kα2 +Fig. 4. (a) Polished surface of the Cu/diamond(Ti) composites and elemental distribution of EDS scan across Cu/diamond interface for (b) Ti, (c) Cu, and (d) C. + +![](images/6fd8144f149866435f89ba76ac0d06954288b567d81e4d10ec131de33f7b513e.jpg) + +![](images/60469da11620853a8f420bdc0eff743ed7714734bd3a782c571aba7e0e5d99db.jpg) + +![](images/31271de49eac9b7a3ee27e5f0f01906e694129afd133f235ea433c3655e55933.jpg) + +![](images/3e733341ab8ef4b54bd25125e4a08eda9d40038a0db2d604a52d4091fadde121.jpg) + +![](images/88f8fd7f3d12f57d20ef0a186bee58f4c89c9c03105a5284d1a5302c96341f5c.jpg) + +![](images/56386c1347e45f4201f2537d77bc0e5e181fb4f1de1453e99053721defb5355e.jpg) +Fig. 5. Auger electron spectroscopy detected elemental depth profile for Ti-coated diamond particles (a) before composite fabrication and (b) dissolved from the Cu/diamond(Ti) composites as well as the corresponding kinetic spectra of C and Ti at various depths: (c) C and (e) Ti for (a) and (d) C and (f) Ti for (b). + +![](images/3e0dee7320ede2b92dd5754bdc5aea0674be6c1495cd90a0aa46895e8a7fc23d.jpg) +Fig. 6. Thermal conductivity of the Cu/diamond composites fabricated by gas pressure infiltration, compared with literature and theoretical value. + +![](images/8d764a6301deaf91f6b4d2ab4c83d64553170935b22967097ddf698319a33e57.jpg) +Fig. 7. Coefficient of thermal expansion of the Cu/diamond(Ti) composites fabricated by gas pressure infiltration. + +# Highlights + +The Cu/diamond(Ti) composites are produced by gas pressure infiltration. +A TiC layer is formed between Cu matrix and diamond reinforcement. +A thermal conductivity of 716 W/mK is obtained for the composites. + A coefficient of thermal expansion of 5.8 ppm/K at 323 K was obtained. \ No newline at end of file diff --git a/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/11_Microstructure and thermal conductivity of Cu-diamond composites with Ti-coated diamond particles produced by gas pressure infiltration/images/1df2128c23710b2f1d2b828fe5a76af5526e2c435ecc4bcaa714a2255fdaf367.jpg b/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/11_Microstructure and thermal conductivity of Cu-diamond composites with Ti-coated diamond particles produced by gas pressure infiltration/images/1df2128c23710b2f1d2b828fe5a76af5526e2c435ecc4bcaa714a2255fdaf367.jpg new file mode 100644 index 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particles produced by gas pressure infiltration/images/e9d4cc23e4ae65fbe8c3bff4026cc76f68158e63deb018c45d4767f86257c4bb.jpg b/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/11_Microstructure and thermal conductivity of Cu-diamond composites with Ti-coated diamond particles produced by gas pressure infiltration/images/e9d4cc23e4ae65fbe8c3bff4026cc76f68158e63deb018c45d4767f86257c4bb.jpg new file mode 100644 index 0000000..8967cac Binary files /dev/null and b/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/11_Microstructure and thermal conductivity of Cu-diamond composites with Ti-coated diamond particles produced by gas pressure infiltration/images/e9d4cc23e4ae65fbe8c3bff4026cc76f68158e63deb018c45d4767f86257c4bb.jpg differ diff --git a/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/12_Enhanced thermal conductivity in diamond0copper composites with tungsten coatings on diamond particles prepared by magnetron sputtering method/12_Enhanced thermal conductivity in diamond0copper composites with tungsten coatings on diamond particles prepared by magnetron sputtering method.md b/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/12_Enhanced thermal conductivity in diamond0copper composites with tungsten coatings on diamond particles prepared by magnetron sputtering method/12_Enhanced thermal conductivity in diamond0copper composites with tungsten coatings on diamond particles prepared by magnetron sputtering method.md new file mode 100644 index 0000000..56ccd88 --- /dev/null +++ b/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/12_Enhanced thermal conductivity in diamond0copper composites with tungsten coatings on diamond particles prepared by magnetron sputtering method/12_Enhanced thermal conductivity in diamond0copper composites with tungsten coatings on diamond particles prepared by magnetron sputtering method.md @@ -0,0 +1,239 @@ +# Enhanced thermal conductivity in diamond/copper composites with tungsten coatings on diamond particles prepared by magnetron sputtering method + +![](images/25e05942078202b26246c98c2425db7d8f406058575a7242a47e543805255a31.jpg) + +Jinhao Jia , Shuxin Bai * , Degan Xiong , Jing Xiao , Tingnan Yan + +College of Aerospace Science and Engineering, National University of Defense Technology, Changsha, 410073, China + +# H I G H L I G H T S + +� The magnetron sputtering is a feasible approach to prepare controllable tungsten coatings for the diamond/Cu composites. +� Effect of the tungsten coatings on properties of the diamond/Cu composites was studied. +� The diamond/Cu composites exhibited an excellent TC of $8 7 9 6 \pm 1 1 \mathrm { ~ W ~ m ~ } ^ { - 1 } \mathrm { ~ K ~ } ^ { - 1 }$ and a fine gas tightness of $2 . 1 \times 1 0 ^ { - 1 0 }$ Pa $\mathfrak { m } ^ { 3 } \ : s ^ { - 1 }$ + +# A R T I C L E I N F O + +Keywords: + +Diamond/Cu composites + +Tungsten coating + +Microstructure + +Thermal conductivity + +# A B S T R A C T + +Tungsten coatings of 45–300 nm thickness were sputtered deposited on diamond particle surfaces, and diamond/ Cu composites were obtained through pressureless infiltration of liquid copper into a bed of coated diamond particles. The composition of the as-deposited coatings was investigated. The effects of coating characteristics on the microstructure, gas tightness and thermal conductivities (TC) of the diamond/Cu composite were also discussed. Results showed that the originally smooth diamond surface are uniformly covered by nanosized tungsten cone arrays. Tungsten coatings with a thickness above 93 nm was found to be the most effective in improving the densification of overall composites and minimizing the interface thermal resistance between reinforcement and matrix. The thermal conductivity of composites reached $7 3 1 \pm 1 0 \mathrm { ~ W ~ m ~ } ^ { - 1 } \mathrm { { K } ^ { - } }$ 1 with a 300 nm tungsten coating and achieved a maximum of $7 \dot { 9 } 6 \pm 1 1 \mathrm { ~ \dot { W } ~ m ^ { - 1 } ~ K ^ { - 1 } }$ with 218 nm coating. In addition, the interface thermal resistance between matrix and coated diamond particles was discussed in accordance with the calculations using a theoretical model. + +# 1. Introduction + +Diamond/Cu composites have been extensively studied as a good candidate for three-dimensional integration communication device packing materials due to its excellent thermophysical properties [1–3]. To further improve the mechanical and physical properties of diamond/Cu, it is essential to optimize the interfacial bonding. Some approaches involved alloying Cu with minor amounts of carbide of carbide formers [4–7], but non-metallic elements are unfavourable for developing the full potential of metal matrix, even resulting in noticeably reduced physical and mechanical properties of the composite. Electrodeposition has been used to fabricate diamond/Cu composites, but microstructure defects are common in a combined interface [8,9]. In addition, although high temperature and high pressure technique can + +minimize the interfacial thermal resistance, it cannot realise the near-net shape of diamond/Cu composite with complex structures [10]. The surface metallisation of diamond particles with carbide-forming metals, such as W, Cr, Mo, and Ti, has been effective in overcoming these obstacles [11–16]. They could preserve the intrinsic merits of the metal matrix coupled with property optimisation in diamond/Cu composites. Compared with TiC, Mo C and $\mathtt { B _ { 4 } C }$ coatings, a single-layered WC coating could substantially improve the performance of diamond/Cu, and TC values over 700 W $\bar { \mathbf { m } } ^ { - 1 } \mathbf { K } ^ { - 1 }$ was achieved through the WC-coated diamond/Cu composites filled with a particle volume fraction of approximately 62% [17–19]. However, most cases applying coatings onto diamond are fabricated using a diffusion method or salt bath treatment. In both cases, the diamond carbon reacts with substances or is easily oxided at high temperature, and obtaining a thin and + +![](images/eb417b4a37996892cf2ae7140e51ca0e3eb09cd46f355442cb3fa4b57a0f53c3.jpg) +Fig. 1. Schematic representations of the assembly unit. + +controllable interfacial carbide layer is unlikely due to the nonuniform nucleation on the diamond surface [20,21]. Abyzov et al. $[ 1 6 , 1 8 ]$ found that the diamond/Cu composite presented thermal conductivities ranging from $5 5 2 \pm 6 \mathrm { W } \mathrm { m } ^ { - 1 } \mathrm { \bar { K } } ^ { - 1 } \mathrm { t o } \bar { 6 } 8 6 \pm 8 \mathrm { W } \mathrm { m } ^ { - 1 } \mathrm { K } ^ { - 1 }$ when tungsten based coating with a thickness of 260 nm is introduced onto diamond particles via diffusion method. However, the thickness of the tungsten + +based coating was thicker than the expectation. They concluded that the optimal values of tungsten based coating thickness for preparing diamond-Ag (Cu) composites range from 100 nm to 250 nm. And the modified layer includes tungsten and tungsten carbide caused by the reaction between coating and diamond during the preparation process. The critical issue is controlling the coatings precisely to maintain a minimum interfacial thermal resistance. Magnetron sputtering has been proved to be useful for preparing uniform and controllable thickness coatings [19,22,23]. Chen et al. [24] prepared tungsten coating by the magnetron sputtering method and no significant debonding was observed in diamond/Al composites with 45 nm and thicker tungsten coatings. Yang et al. [14] investigated the effect of tungsten coatings on the TC value of diamond/Al composites and found that magnetron sputtering was feasible for preparing reliable tungsten coatings for diamond/Al composites. Although magnetron sputtering is an novel strategy to prepare thickness-controllable coatings, the effect of tungsten-coating prepared by the magnetron sputtering method for the diamond/Cu composites has not been systematically investigated. + +In the present work, in order to further optimize the performance of diamond/Cu composites, DC magnetron sputtering was applied to deposit tungsten coatings with a thickness range of 45–300 nm on diamond + +Table 1 Experimentally measured values of Gas tightness and thermal conductivity for the three diamond/Cu Composites. + +
CompositesCoatings thickness (nm)Tungsten to carbon mass ratioThermal conductivity (W·m-1·K-1)Gas tightness (Pa·m3·s-1)RC(m2·KW-1)
H-Jideal value
A450.021500 ± 135.1 × 10-82.00 × 10-70.77 × 10-9
B930.043690 ± 88.5 × 10-97.20 × 10-81.03 × 10-9
C2180.061796 ± 117.1 × 10-103.75 × 10-81.66 × 10-9
D3000.076731 ± 102.1 × 10-106.05 × 10-82.24 × 10-9
+ +![](images/c15a1f421a9d266d53e90c449985f7205328cd24a94c3c8993937696a51e365b.jpg) + +![](images/96cb40df1d1307058fe17b8d083deea418134da3bf215c8f78eada82b8c4c0b2.jpg) + +![](images/cf3e970db6744e49140fd66e80054250be1145f0ba60f252d32861827c73cde2.jpg) + +![](images/b9d854817e405426468f4f624ce040f4dcf68c612ecad292ac807153074313ac.jpg) +Fig. 2. Typical cross-section view image: diamond particles magnetron sputtered for (a) 180min, (b) 360min and (c) 540min, (d) SEM images of diamond particles magnetron sputtered for 90min. + +![](images/4dd3beca457f1b72731392bd4496d29361367113e7c7b727ea6c831c146ed25d.jpg) + +![](images/b1806db45016304e4d3029d07535e6e376d57f56fa8a8e1d1f989ec9d2cdac10.jpg) +Fig. 3. AFM observation the surface of as-deposited tungsten on diamond particles surface; (a) Two-dimensional phase image, (b) three-dimensional topographical image. + +![](images/efeadc9f8256b01320a93abdb6833ee3efab805dbb3270c240707870120b258f.jpg) +Fig. 4. XRD analysis of the diamond particles with various tungsten coatings: (a) 45 nm, (b) 93 nm, (c) 218 nm, (d) 300 nm. + +particles. The composition and morphological characteristics of asdeposited coatings were investigated. The main objective was to fabricate Cu/diamond composites with different interfaces and analyse the effect of coatings with different characteristics on the microstructure and thermal conductivity properties of the Cu/diamond composites. Meanwhile, the interface thermal resistance was theoretically analysed to account for the thermal conductivity behavior. + +# 2. Experimental + +Synthetic hexagonal or octahedral diamond particles (HWD40) with a diameter of 150–180 μm and commercial purity oxygen-free Cu (>99.99 wt%) were used as raw materials. A tungsten metal target (99.95%) with dimensions of 100 mm � 5 mm and a grain size of 300 μm was used for the deposition of tungsten atoms in the experiments. + +Tungsten atoms were deposited on diamond particles using a JGCF-600 DC magnetron sputtering system (Beijing Taikono Technology $\mathrm { C o . , }$ Ltd., China). This system was equipped with a rotary sample holder to prepare controllable coatings on the surface of diamond particles. The temperature in the holder was $3 0 0 ^ { \circ } \mathrm { C }$ and monitored by a thermocouple. + +![](images/277fb9872ffb152df488830962e786b964c0cd95fb0888495674d8078a16c43e.jpg) +Fig. 5. XPS patterns of diamond particles coated with 45 nm tungsten. + +For each magnetron sputtering process, 300 g of cleaned diamond was placed into a sample holder which was rotated at 45 rpm. This process was maintained for 90, 180, 360 and 540 min, a thus tungsten coating with different thickness values was obtained. Other details of the coating preparation are described in Refs. [25]. + +Diamond/Cu composites with about 65 vol% of the coated diamond powers were fabricated via pressureless infiltration in a vacuum of $1 0 ^ { - 2 }$ Pa. Before the infiltration, the uniform distribution of the reinforcement materials was ensured by measuring tap density. The tungsten-coated diamond particles were packed in a graphite mould to form a diamond particle bed. A wire mesh cover was placed into a graphite mould, and Cu was placed on the wire mesh cover. The effect of preparation parameters was eliminated by preparing the packed diamond with various tungsten coatings in the same graphite mould. A schematic of the pressureless infiltration unit is shown in Fig. 1. The graphite mould with diamond particles and Cu was then heated in the carbon tube sintering furnace (WZDS-20, Beijing Research Center) to $1 3 5 0 ^ { \circ } \mathrm { C }$ with a heating rate of 26.6 �C/min for 45 min under vacuum status $( < 1 0 ^ { - 2 } \mathrm { P a } )$ . The graphite mould was naturally cooled to room temperature in the + +![](images/17ef425cc73f136057985aa6697a36e4b5df05260c788127028d576288fb4778.jpg) + +![](images/db605e391859e4428b0483264fa75e8adfd81baadb1868fc7e009b61e7620472.jpg) + +![](images/c1c06abc283d2145604554bd057e26d560da552e663d3957aeba9141fc472a21.jpg) + +![](images/18b8de37f0af05469c7e89fb51874b30485a87b53b702f5a4ef32d432e42554f.jpg) +Fig. 6. SEM images of the fracture surfaces of the diamond/Cu composites with different coating thickness: (a) 45 nm, (b) 93 nm, (c) 218 nm, (d) 300 nm. + +furnace. Disc-shaped composite samples with a diameter of 12.7 mm and height of 3 mm were obtained. Materials and processing conditions are summarised in Table 1 in which the composites with approximately 45, 93, 218 and 300 nm tungsten coatings were denoted as ‘A’, ‘B’, ‘C’ and ‘D’, respectively. + +The microstructure of the coated diamond particles and the diamond/Cu composites were observed by an S4800 scanning electron microscope (SEM) equipped with energy dispersive spectroscopy. Atomic force microscopy (AFM) was performed on Dimension Icon to measure the roughness of the tungsten coatings. The interior structural information of diamond/Cu composites was analysed using an NDT/ analyser high-energy X-ray nondestructive detection system. The thermal conductivity of composites was measured using a Netzsch LFA427 laser flash machine. A helium mass spectrometry leak detector ZQJ 542 was used to measure gas tightness. The XPS spectrum was obtained from Krators AXIS Ultra DLD via the X-ray radiation of 1253.6 eV Mg Ka to analyse the chemical state of the coatings. The phase constitutions of the coated diamond particles were examined via X-ray diffraction (D/maX 2550 VB þ X-ray diffraction). + +# 3. Test results and discussions + +# 3.1. Formation of the tungsten coatings on diamond particles + +The thickness AFM and SEM analyses were performed to investigate the surface morphology of the coatings. The thickness of coatings determined by gravimetric data was obtained by accurately comparing + +the mass of diamond particles before and after sputtering. The thickness of coatings was calculated under the assumption that each diamond particle is a sphere with the same size. The calculated results were higher than the actual results because the coatings were thin, and some particles were lost during sputtering. + +The thickness of coatings was detected by breaking the coated diamond particles via thermal vibration combined with mechanical crushing. Subsequently, the exposed cross section was characterised via SEM to acquire coating thickness. The representative observation of the thickness of coatings is shown in Fig. 2 a–c. The profiles of the thickness indicate that the average thickness values of the coatings for 180, 360 and 540 min were approximately 93, 218 and 300 nm, respectively. The thickness of coatings approximately linearly increased with sputtering time. The coating thickness was approximately 45 nm when the sputtering time was 90 min by fitting to least squares. Typical SEM pictures of a diamond particle magnetron sputtered for 90 180, 360 and 540 min are presented in Fig. 2 a–d. A pronounced resemblance in the morphology of the coated diamond particles was exhibited with various magnetron sputtering times. The diamond particle was fully covered, and the coatings were smooth and dense. Moreover, high magnification image (Fig. 2 d) revealed that the diamond particles were evenly covered by submicron particles, and no observable voids were found. + +Fig. 3 a shows the same image as Fig. 2 d in three dimensions and reveals a continuously lumpy character. The FAM observation (Fig. 3 b) topographical image clearly shows nanosized particles, confirming the surface morphology obtained from SEM and highlighting the presence of nanocrystalline particles. + +![](images/0531d462efbaed6cbd5ba36c9cbdba658afb57c183781b9c172e93ac1086669d.jpg) + +![](images/fc5d4a139c9810c587bb8c199779c4ec31bf64d4d7474afec77d91e11606be07.jpg) + +![](images/c8dcc1648ade31da4131e7f7b4fb27f7bf9afb3663bd36fca33a02c142a79f2c.jpg) + +![](images/0cacff90174f260cfe414371e5d6cf353f3d19e8e8e482b8d79085a0cec1b0b2.jpg) + +![](images/f1fe9b9938c13f3cb34509e241b9995439609616c6872b2ffb59a0937fef4b42.jpg) + +![](images/c7cd45e74aee88b187d75694082162441b28b68424c790468fa4ad707089ef83.jpg) +Fig. 7. EDS analysis: (a) the magnified image of fracture surfaces of composite A, (b) The EDS line-scanning analysis of the coated diamond surface from composite A, (c) EDS element distribution mappings of composite A, (d) (e) and (f) EDS mapping distributions of C of composite B, C and D. + +The peak to peak (Ry), mean roughness (Ra) and ten-point height (Rz) of the coating measured by AFM were 6.8, 0.45 and 3.4 nm, respectively. These values indicate a small disordered region between peaks and valleys on the coating surface. Abyzov et al. [3,18] reported that a tungsten-based coating of 100 nm was prepared on the surface of diamond particles via diffusion treatment. However, the Ra of the sur face was 37 nm, which is unfavourable for the TC properties because the electron transport at the metal interfaces is controlled by elastic scattering by disorder. Thus, magnetron sputtering can be a suitable method + +for preparing reliable tungsten coatings. + +# 3.2. XRD and XPS analysis + +In this work, XRD characterised the filler particles as a whole, and XPS spectra captured the local near-surface region. XRD analysis of diamond particles with different coating thicknesses are presented in Fig. 4. The three main peaks at 2θ ¼ 40�, 44� and 74� correspond to the (110), (200) and (211) crystal planes of a body-centred cubic tungsten, + +![](images/474cf221e4909d82153fca2cf2b640547892635605e29006a6960c2fd77cb8ba.jpg) +Fig. 8. XRD pattern of diamond particles with 300 nm coating underwent infiltration process. + +respectively. It indicated that the reaction temperature between tungsten and diamond did not occur at the preparation temperature $( 3 0 0 ^ { \circ } \mathrm { C } )$ , and no carbide was formed during magnetron sputtering. + +The composition and bonding states of the coatings were further characterised by XPS. Fig. 5 shows the XPS fitted results of tungsten 4f from the coating. The symmetry peak at 35.7 and 37.8 eV in the W 4f spectrum were consistent with the characteristics of hexavalent tungsten with intermediate oxidation states based on the NIST XPS database [26–29]. The two lower binding energy peaks of tungsten 4f located at 31.1 and 33.2 eV exhibited that zero-valent tungsten could be verified, which is in conformity with the XRD results. The bond of W–O and W–W was the main bonding mode, and no substantial carbonisation formed during magnetron sputtering. The combined XPS and XRD analyses proved that a coating with a dominant tungsten phase and the bond of W–W and W–O was the main bonding mode on the coating surface. Details of the analysis are provided in Ref. [25]. + +# 3.3. Microstructure of diamond/Cu composites + +Fig. 6 displays the microstructure of bending fracture surfaces of the composite samples. In the fracture mode, the coated diamond was often pulled out from the matrix and the ductile fracture of Cu [29]. However, the composites exhibited considerably different fracture morphologies with respect to various coating thicknesses. Fig. 6 c and d show that the diamond particles were distributed in the Cu matrix uniformly and integrated compactly with the Cu matrix through the carbides as a binder. This phenomenon explains why coatings guarantee the good adhesion between the diamond and the Cu matrix to form a strong mechanical interlocking. Strong interfacial bonding was obtained when WC Cr7C3, $\mathtt { M o } _ { 2 } \mathrm { C }$ and $\mathtt { B _ { 4 } C }$ were utilised as coating [22,30–33]. The fracture surface of composites C and D (Fig. 6 c and d) was distinctive. Almost all diamond particles exhibited transcrystalline fractures, indicating a suitable coating thickness can ensure a strong interfacial bonding between the filler and the matrix. By contrast, Fig. 6 a–b show that the interface gap is evident between the reinforcements and the Cu matrix, suggesting a weak mechanical bonding between the diamond and the Cu matrix. Moreover, bare diamond surfaces were observed on the fracture surface of the coated diamond/Cu composites, as shown in Fig. 6 a and $\mathbf { b } ,$ leading to the diamond particles being intact. + +The typical diamond particle from composite A, as shown in Fig. 7 a, and the EDS line-scanning analysis in Fig. 7 b further confirm that the coating on the diamond particles was discrete irregular carbides. Fig. 7 + +c–f show the EDS mapping analysis of the cross-sectional microstructures of diamond/Cu composites as a representative. A good interfacial adhesion can be observed in Fig. 7 d–f. Moreover, the interface was straight and overlapped, indicating that the coating morphology was stable during infiltration, and the interface became flatter as the thickness of the coating increased. + +XRD analysis shows that only WC and diamond existed in the diamond particles extricated from composite D (Fig. 8). Combined with the XRD results of ref [24], it can be confirmed that different thickness coatings were completely converted to WC. The typical diamond particle from composite $\mathbf { A } ,$ as shown in Fig. 7a, and the EDS line-scanning analysis in Fig. 7 b further confirm that the coating on the diamond particles was discrete irregular carbides. The interface was prone to contraction and separation along the bare diamond surface due to the existence of the sparse coating surface, leaving behind the micropores. Therefore, producing a bond of low thermal resistance and a full metallurgical bonding between the reinforcement and the Cu matrix is difficult. + +Fig. 9 shows the X-ray CT image of the diamond/Cu composite samples. The diamond particles were distributed in the Cu matrix uniformly in Fig. 9 c and d, which is consistent with those in the SEM image (Fig. 6 c and d). With the decrease in coating thickness, the coatings gradually broke up. Liquid Cu phase allows for the aggregation of floating particles and a descending Cu-rich zone, as shown in Fig. 9 b. The inner structure of composite A is similar to that of composite B, as shown in Fig. 9 a and b. The only difference is that a few pores exist inside composite $\mathbf { A } ,$ indicating that the homogeneity of diamond/Cu composites is worsened with the decrease of coating thickness. It is understood that the fracture surfaces of composite A is strongly influenced by porosity. When the thickness increased to approximately 93 nm, microstructure gets worse due to the following reasons: firstly, diamond particles with thin coatings are prone to interfacial failure due to shrinkage and rupture of coatings. As a result, carbon affinity becomes poor, inducing the rearrangement of diamond particles. Secondly, the difference in the density of Cu and coated diamond is greater due to the relative long contact time at high temperature; thus, the diamond particles are easy to shift. Thirdly, the currently used infiltration temperature of 1300 �C could decreased to liquid Cu viscosity, which facilitate diamond particles to float during pressureless infiltration. Therefore, according to the infiltration kinetics, it is essential to increase the coating thickness and increase the contact time between diamond particles and copper melt to yield diamond/copper composites with homogenous microstructure. + +# 3.4. Properties of diamond/Cu composites + +Experimentally obtained properties of diamond/Cu composites with various coatings are shown in Table 1. The resulting composites had thermal conductivity values ranging from 487 $\boldsymbol { \mathrm { \nabla } } \mathrm { m } ^ { - 1 } \mathrm { K } ^ { - 1 } \mathrm { t o } 8 0 7 \boldsymbol { \mathrm { W } } \mathrm { m } ^ { - 1 }$ $\mathrm { K } ^ { - 1 }$ . The thermal conductivity of composites firstly increased and then decreased with the increase in coating thickness. The highest thermal conductivity of $7 9 6 \pm 1 1 ~ \mathrm { W ~ m } ^ { - 1 } ~ \mathrm { K } ^ { - 1 }$ was achieved in the case of composites with 218 nm coatings. Further increment in the thickness of tungsten coating led to the TC value of the composite decreasing by 8.1% because the thick coating layer introduced a large interfacial thermal resistance. In the absence of shrinkage or rupture, the thinner the coating, the lower the interfacial thermal resistance. This behavior favoured improving thermal properties, but it was not a dominant factor in composite A. The lowest thermal conductivity value was measured in the case of composite A, which used the diamond with the 45 nm tungsten coating. The gas tightness data (Table 1) corresponding to the diamond/Cu system were consistent with the TC mentioned above. The gas tightness of composite A was approximately $5 . 1 \times 1 0 ^ { - 8 }$ , which is much higher than that of composites C and D. In the case of Al/diamond composites [14] fabricated by vacuum infiltration, tungsten coating with 45 nm was used to improve wettability and enhance the thermal + +![](images/3d13390d2e47d2fcecb801e1a1dcea6b0ea7aa5e606c64ca3fdf37180d8e4271.jpg) + +![](images/7f035b3f9cb6ac0a94b862a0a01c30faf76c67a48b4e2e5153de5f4e25812834.jpg) + +![](images/ec2f7f0e3fa0f69308f7133ee92f414f6a0a5ab6e2ccde386dfcbf0b5c765f33.jpg) + +![](images/43933af6d45c965d16cdef132b93ad26f8b3d1372f264778e66b388e29ad5209.jpg) +Fig. 9. X-ray imaging of diamond/Cu composites (a) with 45 nm coatings, (b) with 45 nm coatings, (c)with 218 nm coatings, (d) with 300 nm coatings. + +Table 2 Parameters of materials for interfacial thermal resistance calculation. + +
MaterialDensity (g·cm-3)Heat capacity (J·Kg-1·K-1)Sound velocity (m·s-1)Thermal conductivity (W·m-1·K-1)
Diamond3520510127041500
Copper89603852970400
WC157003554706173
+ +conductivity to a level of 622 $\mathbf { W \ m } ^ { - 1 } \mathbf { \ K } ^ { - 1 }$ . Compared with microstructures, a bare diamond surface was not found in Al/diamond composites with 45 nm tungsten coatings because the liquid Al infiltration maintained a relatively low temperature $( \mathrm { < } 9 0 0 ^ { \circ } \mathrm { C } )$ , and the reaction between W and diamond occurred above $9 8 3 . 6 \ ^ { \circ } \mathrm { C } ,$ which was lower than the reaction temperature between tungsten and diamond $( \geq 9 8 3 . 6 ^ { \circ } \mathbf { C } )$ [22]. These results suggest that the decreased thermal conductivity of composite A can be attributed to the added internal microdefects, which effectively hinder the heat-conducting flow. Thus, the optimal values of tungsten coating thickness prepared by magnetron sputtering for forming diamond/Cu composites should exceed 93 nm. Composite D has + +further potential because it achieved a relatively high thermal conductivity, especially considering its relatively low leak rate. + +We considered the incompleteness of the tungsten carbide interface. Based on the Hasselman and Johnson (H-J) model [34]. The interfacial thermal resistance $( R _ { C } )$ can be deduced using measured values, which can be written as: + +$$ +R _ {c} = \frac {\mathrm {a}}{\lambda_ {D i a}} \times \left[ \frac {V _ {D i a} \times \left(\frac {\lambda_ {D i a}}{\lambda_ {C u}} - 1\right) \left(2 + \frac {\lambda_ {C}}{\lambda_ {C u}}\right) + \left(2 + \frac {\lambda_ {D i a}}{\lambda_ {C u}}\right) \left(1 - \frac {\lambda_ {C}}{\lambda_ {C u}}\right)}{V _ {D i a} \times \left(2 + \frac {\lambda_ {C}}{\lambda_ {C u}}\right) + \frac {2 \lambda_ {C}}{\lambda_ {C u}} - 2} \right] \tag {1} +$$ + +where $\lambda _ { C } , \lambda _ { C u }$ and $\lambda _ { D i a }$ are the effective thermal conductivities of the composite, Cu matrix $( \lambda _ { C u } = 4 0 0 )$ and diamond particles $( \lambda _ { D i a } = 1 5 0 0 )$ respectively; $\mathrm { V } _ { D i a }$ is the volume fraction of diamond particles (65%); a is the average radius (180 μm) of the reinforcement. + +The calculated results are shown in Table 1. Both an unusual decrease in the interface thermal resistance of the composite with an increase of the coating thickness, and a predictable increase in interface thermal resistance with increasing thickness were observed. The interface thermal resistance values deduced from H-J moldel of the three diamond/Cu composites (A, B and C) slightly decreased with the + +increase in coating thickness. To the best of the authors’ knowledge, the weaker the adhesion at the interface between the coated diamond and Cu, the lower the strength of the composite and the higher the thermal resistance of the interface. Combined with previous research findings [25], the high thermal conductivity obtained was associated with appropriate interface thickness and good composite density. + +Meanwhile, the analysis thermal resistance of the perfect boundary can be analogous to a series circuit in which the total thermal resistance is composed by the sum of three terms $( R _ { \mathrm { c } } = R _ { W C / C u } + R _ { D i a m o n d / W C } +$ $R _ { W C } ) { : }$ diamond/WC interface RDiamond/WC, WC/Cu interface $R w C / C u$ and the resistance of the WC layer itself $R _ { W C } .$ . Interface thermal resistance $R _ { W C / C w } \ R _ { D i a m o n d / C u }$ and RDiamond/WC can be calculated using the acoustic mismatch mode (Eq. (2)) [35,36]. + +$$ +R = \frac {2 \left(\rho_ {m} \nu_ {m} + \rho_ {\mathrm {p}} \nu_ {p}\right) ^ {2}}{C _ {m} \cdot \rho_ {m} ^ {2} \cdot \nu_ {m} ^ {2} \cdot \rho_ {p} \cdot \nu_ {p}} \left(\frac {\nu_ {p}}{\nu_ {m}}\right) ^ {2} \tag {2} +$$ + +where $\rho ,$ C and v are the density, specific heat capacity and Debye velocity of the matrix (longitudinal or transverse). Subscripts m and p refer to materials on the matrix and diamond side. $R _ { W C }$ was calculated by Eq $R _ { W C } = L / \lambda$ , where L is the average thickness of tungsten, which was recalculated into the average thickness of corresponding carbide. Table 2 lists the related data used in calculations based on Refs [25]. Amongst the composites, composite D had the highest Rc $( 2 . 2 4 \times 1 0 ^ { - 9 }$ $\mathrm { m } ^ { 2 } { \cdot } \mathrm { K W } ^ { - 1 } )$ , and composite A had the lowest Rc $( 0 . 7 7 \times 1 0 ^ { - 9 } \mathrm { m ^ { 2 } { \cdot } K W ^ { - 1 } } )$ . The obtained result implies that the thermal transport across interface is the most efficient in composite A, but composite A has the lowest thermal conductivity. The results only reflect a perfect bonding condition because they ignored the gap between discrete carbides. Thus, a thin coating could deteriorate the thermal conductivity of composites because a volume contraction and discontinuous carbide layer could not effectively improve the interfacial bonding of the composites. + +# 4. Conclusion + +Tungsten coatings with different thicknesses were deposited onto diamond particles via magnetron sputtering, and diamond/Cu composites were successfully fabricated via pressureless infiltration. A thickness of 218 nm is the optimum combination to improve the densification of overall composites and minimize the thermal boundary resistance between diamond particles and copper. Meanwhile, the densification of overall composites was improved with thick coatings, and the gas tightness reached 2.1 Pa $\mathbf { m } ^ { 3 } \thinspace \bar { \mathbf { s } ^ { - 1 } }$ with a 300 nm tungsten coating. Decreasing the coating thickness to 93 nm will worsen the interfacial bonding, resulting in the heterogeneous microstructure and performance degradation. The lowest thermal resistance of the diamond-WC-copper interface is $3 . 7 5 \times 1 0 ^ { - 8 } \mathrm { ~ m } ^ { 2 } \mathrm { ~ K W } ^ { - 1 }$ in accordance with the calculations using the H-J model. + +# Declaration of competing interest + +The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. + +# CRediT authorship contribution statement + +Jinhao Jia: Conceptualization, Methodology, Investigation, Writing - original draft. 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0000000..eaebb59 --- /dev/null +++ b/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/13_Optimized thermal properties in diamond particles reinforced copper-titanium matrix composites produced by gas pressure infiltration/13_Optimized thermal properties in diamond particles reinforced copper-titanium matrix composites produced by gas pressure infiltration.md @@ -0,0 +1,192 @@ +# Optimized thermal properties in diamond particles reinforced coppertitanium matrix composites produced by gas pressure infiltration + +![](images/3fb71ad3ab983e19669a079a18aa3ac5a9485d838a4615a429b7ae498aba3618.jpg) + +Jianwei Li a , Hailong Zhang a , Luhua Wang a , Zifan Che a , Yang Zhang a,1 , Jinguo Wang b , Moon J. Kim b , Xitao Wang a,⇑ + +a State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, China + +b Department of Materials Science and Engineering, The University of Texas at Dallas, Richardson, TX 75080, USA + +# a r t i c l e i n f o + +Article history: + +Received 9 May 2016 + +Received in revised form 12 August 2016 + +Accepted 1 October 2016 + +Available online 5 October 2016 + +Keywords: + +A. Metal-matrix composites (MMCs) + +B. Interface/interphase + +B. Thermal properties + +E. Liquid metal infiltration + +# a b s t r a c t + +Interface modification is crucial to exploit high thermal conductive potential of diamond in the metal matrix composites reinforced with diamond particles (Cu/diamond composites). With an attempt to modify the Cu/diamond interface, we add a carbide-forming element of Ti to the Cu matrix and use a liquid-phase processing technique to attain sound interfacial bonding. The Cu-xTi/diamond composites were characterized by using scanning electron microscopy, transmission electron microscopy, and Xray diffraction. The interface layer is confirmed as TiC, the amount of which increases with increasing Ti concentration in the Cu-xTi alloy matrix. As the Ti concentration increases, the thermal conductivity of the Cu-xTi/diamond composites first increases and then decreases, giving an optimized thermal conductivity of 752 W/m K and a coefficient of thermal expansion of 6.50 106 /K at x = 0.5 wt.%. The results show that an appropriate amount of Ti addition in Cu matrix can enhance the thermal conductivity of Cu/diamond composites. + + 2016 Elsevier Ltd. All rights reserved. + +# 1. Introduction + +The miniaturization of electronics is urgently demanding packaging materials with superior thermal conductivity. Effective thermal management is fundamental to consistent long-term performance and reliability of electronic devices [1,2]. As a promising candidate, diamond particles dispersed copper matrix (Cu/diamond) composites are attracting more and more interests [3–17]. In preparing Cu/diamond composites, the major restriction originates from the mismatch in thermal expansion coefficient between diamond and copper, which easily induces large thermal stress. In addition, the non-carbide forming nature of copper and the large contact angle between diamond and Cu mean an absence of strong chemical bonds at the interface. Consequently, the diamond particles are inclined to delaminate from the pure copper matrix after cooling down. Poor interfacial bonding becomes a main obstacle to obtaining high thermal conductivity in the Cu/diamond composites. The thermal conductivity of the Cu/diamond composites is reduced even with increasing the amount of diamonds [4]. + +At present, it is popular to improve the interfacial bonding by introducing a third component between diamond and copper, usually through surface metallization of diamond particles [3,7,8,11,18] or through alloying of copper matrix [6,19,20]. Among broad researches, TiC is suggested to connect the diamond reinforcements with the metal matrix tightly. TiC has a Gibbs free energy of 170.5 kJ/mol [21], which is favorable to its formation at the interface. Through Ti coating on diamond particles, the derived Cu-Ti/diamond composites display a thermal conductivity as high as 630 W/m K, owing to TiC formation at the interface [22]. Nevertheless, the coating on diamond is expensive and demanding. Moreover, the coating is oxidized during high-temperature processing [7]. The presence of oxygen on TiC surface strongly inhibits the interaction between the carbide and the molten Cu matrix, causing non-wetting conditions [23]. By means of alloying Ti to Cu matrix, the diamond particles and the Cu-Ti alloys can react directly, and the formation of oxide layer on the diamond particles is thus avoided. From the above, the route of Cu matrix alloying is proved to be more feasible in the fabrication of Cu/diamond composites. + +Here we add Ti to Cu matrix to demonstrate the enhancement of thermal conductivity in the Cu-Ti/diamond composites. Intuitively, the amount of Ti addition should be limited in order to maintain high thermal conductivity in the Cu matrix. Literature + +has not reported detailed information about the effect of alloying element content on the microstructure and thermal properties of the composites. Several groups [20,24] prepare the Cu-Ti/ diamond composites using a solid-state sintering route, but obtain relatively low thermal conductivities. A liquid-state processing route of gas pressure infiltration has been successfully utilized to achieve high thermal conductivity in Cu/diamond composites [7,16]. The gas pressure infiltration is capable of realizing dense microstructure and promoting interfacial reaction in the diamond composites. So far the promising route has not been used to prepare the Cu-Ti/diamond composites. + +Another difficulty in Cu/diamond composites community comes from delicate characterization of the Cu/diamond interface. Due to huge difference in hardness between diamond and metal, it is extremely tough to mill an eligible specimen for transmission electron microscopy (TEM) observations. For this reason, actual interface structure of the Cu/diamond composites is rarely reported [10,15], which is certainly critical to understand the thermal conducting mechanism involved. Taking advantage of focused ion beam (FIB) technique, we have developed thin foils containing the Cu/diamond interface and obtained the state-of-the-art images depicting carbides at the interface. + +In this article, we produce the Cu-Ti/diamond composites by the gas pressure infiltration and investigate the effect of Ti addition on the microstructure and thermal properties with respect to diamond surface state and interface structure of the composites. The effect of interfacial layer thickness on thermal conductivity of the Cu/diamond composites is elucidated. The thermal conductivity and coefficient of thermal expansion (CTE) of the Cu-Ti/diamond composites are correlated with various modeling schemes. + +# 2. Experimental + +The Cu-xTi/diamond composites were produced by a gas pressure infiltration method. Fig. 1 shows the schematic drawing of the gas pressure infiltration device. The starting materials were HHD90 synthetic single-crystalline diamond powders with an average particle size of 230 lm (Henan Huanghe Whirlwind Co., China) and Cu-xTi alloys (x = 0.3, 0.5, 2.0 wt.%). The Cu-xTi alloys were melted by a vacuum induction route using 99.999 wt.% Cu and 99.99 wt.% Ti bulks (TaiYu Materials Science & Technology Co., China). The infiltration was conducted at 1423 K for 30 min, under an Ar gas pressure of 1.0 MPa. The details of the infiltration are referred to elsewhere [7]. + +X-ray diffraction (XRD, Rigaku DMAX-RB, Japan) was used to characterize the phase structure of the composites. The polished surface and fracture surface were observed by field emission scanning electron microscope (SEM, ZEISS SUPRA 55, Germany). A dual beam FIB system (FEI Nova 200 FIB, USA) was used to mill the Cu-Ti/diamond samples to thin foils. The FIB-milled samples were characterized using a scanning transmission electron microscope (STEM, JEOL, ARM200, Japan). Besides, the metal matrix was electrochemically etched and the diamond particles were collected to directly characterize the interface structure [25]. The composites were electrochemically etched in a 10 vol.% HNO3 and 90 vol.% H O solution for 5 min. The sample acts as the anode and a steel plate as the cathode. The current density was controlled by a direct current power supply and the current was fixed at 1 A. After electrochemical etching, the samples were cleaned ultrasonically in an acetone bath for 40 s. + +The thermal diffusivity (a) was measured by a laser flash method (LFA427, NETZSCH, Germany) with an international standard using a disc 10 mm in diameter and 3 mm in thickness. Three times were repeated for one sample to get an average value of the thermal diffusivity. The thermal conductivity (K) can be calculated + +![](images/40390edff2090b83ea6422ae80745187e88492c165ba738bc37587836fe6b9d1.jpg) +Fig. 1. Schematic drawing of the gas pressure infiltration device. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.) + +according to the equation K = aqc, where q is sample density and c is specific heat. The density was measured by the Archimedes method using alcohol as the immersion medium, and the measurement was repeated three times for each sample. The heat capacities were calculated from the rule of mixture based on mass fraction of each component. The CTEs of the composites were measured by a dilatometer (DIL 402C, NETZSCH, Germany) with a heating rate of 5 K/min in the temperature range between 323 K and 573 K, using a cylinder 5 mm in diameter and 25 mm in length. + +# 3. Results and discussion + +Fig. 2 shows the representative SEM images of the polished surface and fracture surface of the Cu-0.5 wt.%Ti/diamond composite. The diamond particles are found to be uniformly distributed in the Cu matrix (Fig. 2a) and no noticeable defects such as cracks or flaws are observed at the interface. The pull-out of diamond particles is rarely seen in the polished surface, which indicates strong interface bonding between the diamond particles and Cu matrix. The diamond particles maintain their original shape without any degradation. The observations demonstrate excellent quality of the composites produced by the gas pressure infiltration. + +As seen from the fractured surface, the Cu matrix closely adheres to the surfaces of the diamond particles (red circles in Fig. 2b). Some diamond particles are found to fracture transgranularly. This fracture only occurs when the interfacial bonding strength is higher than the fracture strength of the diamond particles. The synthetic diamond has some flaws inside. It can be concluded that Ti addition has significantly improved the interfacial bonding between the Cu matrix and diamond particles. + +In order to study the effect of Ti addition on the phase composition, the XRD patterns of the Cu-Ti/diamond composites were characterized, as shown in Fig. 3a. The results indicate the coexistence of Cu, diamond and TiC phases in the composites. Fig. 3b shows the XRD patterns of the diamond particles extracted from the Cu-Ti/diamond composites. Only diamond and TiC phases are detected. The morphology of the diamond particles is clearly seen in the inset to Fig. 3b. The diamond particle surface is covered uniformly by the carbides. Owing to the interfacial reaction between the Cu-Ti matrix and the diamond, TiC is formed during the infiltrating process by the reaction between Ti and C atoms (Ti + C ? TiC). Thermodynamic calculation shows that titanium carbide will be produced by the reaction of carbon with copper-tin-titanium alloys that have a titanium activity larger than 0.1 at 1400 K and close to 1.0 at 1423 K [26], which is the infiltration temperature + +![](images/695e3dc5f77f060ba856c18f41b80b81c0ca11f9928f7bd31600114d0d4b7d18.jpg) + +![](images/ca0ccad06bd849707d7b732cdec139fddc402fc8e5f10361510df2eb07449e3d.jpg) +Fig. 2. SEM observations of (a) the polished surface and (b) the fractured surface of the Cu-0.5 wt.%Ti/diamond composite. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.) + +in this study. Accordingly, the Ti atoms dissolved in the Cu matrix will react with the C atoms on diamond surface. The peak intensity grows stronger with increasing Ti content, indicating more amount of TiC formed. The interfacial bonding of the Cu-Ti/diamond composites is therefore enhanced by the interfacial phase TiC. + +To further study the effect of Ti addition on interface microstructure, the FIB-milled thin foils were used for SEM characterizations. As shown in Fig. 4a and b, it is clear that the composites consist of three distinct zones: the matrix, the interphase and the diamond. The diamond is found to be bound tightly to the matrix. The carbides act as a binder between the diamond and the matrix. The average size of the carbides is found to increase from 100 nm to 200–300 nm when Ti content increases from 0.3 wt.% to 0.5 wt.%. The carbides are not continuous and they are separated by the Cu alloy matrix. The result is explained by the heterogeneous nucleation of carbides at interface and the following preferential growth along certain crystal orientations from the diamond into the Cu alloy matrix. This is similar to carbide characteristics in the Al/diamond composites [27]. When Ti content is further increased to 2.0 wt.%, a continuous layer of carbides is formed, with a thickness of roughly 300–500 nm. The SEM image in Fig. 4c was taken from diamond particles extracted from the composite sample. As the carbide grains grow into the Cu alloy matrix, the transport of carbon source from the diamond is restrained and then new carbide grains nucleate between the primary carbides. Accordingly, continuous carbides are derived on the diamond surface. + +The thermal conductivity K of the Cu-Ti/diamond composites exhibits an apparent variation with Ti addition (Fig. 5). With + +![](images/2ded01fa22ae66c562758e0360e97b52830f77513ff13f1ad18b816270b41c10.jpg) + +![](images/e65db19da40573b35f3785cd31930cbea6f50bd71ceaca1fbe902bd2e5d3f442.jpg) +Fig. 3. XRD patterns of (a) the Cu-Ti/diamond composites produced by gas pressure infiltration and of (b) the diamond particles extracted from the Cu-Ti/diamond composites. The inset to (b) shows SEM images of a diamond particle extracted from the Cu-0.5 wt.%/diamond composite. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.) + +increasing Ti content the K value firstly increases and then decreases, attaining a maximum value of 752 W/m K at 0.5 wt.% Ti. The unmodified Cu/diamond composite shows a low K value of 141 W/m K, owing to poor interfacial bonding between diamond and Cu. As the formed TiC connects the two phases tightly, phonon transfer across the interface is intensified and a high K value is obtained. Nevertheless, a redundant amount of Ti addition results in a K value as low as 123 W/m K at 2.0 wt.% Ti, even lower than that for unmodified Cu/diamond composite. Since TiC possesses a low K of just 21 W/m K [28], lower than both Cu matrix and diamond, the TiC layer is harmful to heat transport from diamond to Cu. Combined with the SEM images in Fig. 4, the variation of thermal conductivity is explained by the models in Fig. 6. The total interfacial thermal resistance in the transition region $R _ { \mathrm { i } }$ can be assumed as a parallel of the thermal resistances RTiC and RCuTi, where the subscripts TiC and CuTi refer to TiC layer and Cu alloy matrix, respectively. When Ti content increases from 0.3 to 0.5 wt.%, the average size of the carbides at the interface is found to increase from 100 nm to 200–300 nm. The decrease in number of grain boundary is beneficial to heat transfer across the interface. When Ti content is below 0.5 wt.%, the carbides are not continuous and they are separated by the Cu alloy matrix. The remaining Cu alloy is still a continuous network and helps to maintain high ther- + +![](images/adce335795992b8d386489a7a3f846c355ecb158fc32b31698efce01a63347ea.jpg) + +![](images/489fc224403777ce2652de9c31a52a107619f0f5eec54446587a66cdf415f2ed.jpg) + +![](images/628a136be6e245bd3bb511ab80fab075f367e1576caf93328da65b8fab4d60d2.jpg) +Fig. 4. Electron microscopy images showing the interfaces between diamond and Cu-xTi matrix: (a) x = 0.3 wt.%, TEM, (b) x = 0.5 wt.%, TEM, and (c) x = 2.0 wt.%, SEM. + +![](images/0a003c4ac9ce07c8cd98b30d65212d692858253b45db897ab5a7d3cce5004b7f.jpg) +Fig. 5. Variation of thermal conductivity of the Cu-Ti/diamond composites with Ti content. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.) + +mal conductivity. With increasing Ti content from 0.3 to 0.5 wt.%, the interfacial bonding could be improved with mechanical interlocking by larger TiC particles that are penetrated deeper into the Cu matrix. The above-mentioned two factors promote the thermal conductivity of the Cu-xTi/diamond composites. However, the gap between TiC particles decreases with increasing Ti content. This could impair the thermal conductivity of the composites because of the narrowed channel for heat transfer. By a compromise, the thermal conductivity increases from 710 to 752 W/m K with increasing Ti content from 0.3 to 0.5 wt.%. When Ti content is further increased to 2.0 wt.%, a continuous TiC layer is formed, which separates the diamond and the Cu alloy matrix completely. Since TiC has a low thermal conductivity, the thermal conductivity of the Cu-xTi/diamond composites drops dramatically. + +The resultant 752 W/m K is protruding among K values reported for Cu-Ti/diamond composites: 670 W/m K via spark plasma sintering (SPS) [20,22] and 608 W/m K via pressureless sintering method [24]. However, the value is still lower than the theoretical value of 1047 W/m K, which is calculated by using a differential effective medium (DEM) model [29]: + +$$ +(1 - V _ {d}) \left(\frac {k _ {c}}{k _ {m}}\right) ^ {\frac {1}{3}} = \frac {\left(k _ {d} ^ {\text {e f f}} - k _ {c}\right)}{\left(k _ {d} ^ {\text {e f f}} - k _ {m}\right)} \quad \text {w i t h} \quad k _ {d} ^ {\text {e f f}} = \frac {k _ {d}}{1 + \frac {k _ {d}}{a h c}} \tag {1} +$$ + +where $k _ { \mathrm { c } } , k _ { \mathrm { d } }$ and $k _ { \mathrm { m } }$ are the thermal conductivities of composite, dispersed reinforcement and matrix, respectively. The a and $V _ { \mathrm { d } }$ are the radius and volume fraction of spherical reinforcement, respectively, and the $h _ { \mathrm { c } }$ is interfacial thermal conductance. Accordingly, the concentration of Ti addition requires further fine optimization to achieve higher thermal conductivity in the Cu-Ti/ diamond composites. This job is left for future study. + +Fig. 7(a) shows the CTE values for the Cu-0.5 wt.%Ti/diamond and unmodified Cu/diamond composites. The CTE of the Cu-0.5 wt.%Ti/diamond composite was $6 . 5 0 \times 1 0 ^ { - 6 } / \mathrm { K }$ at 323 K, which is compatible with the CTE of semiconductor materials [30] and is lower than reported values for Cu/diamond composites [8,31,32]. The unmodified Cu/diamond composite has a higher CTE of $7 . 2 6 \times 1 0 ^ { - 6 } / \mathrm { K }$ at 323 K. When the bonding between the matrix and the diamond is weak, the Cu matrix would dominate the CTE value of the composite. On the contrary, in the Cu-0.5 wt.%Ti/diamond composite, both the Cu matrix and the diamond would contribute to the CTE of the composite because of the sound interface bonding. The CTE value increases with increasing temperature, which is consistent with the Grueneisen theory [33]. The CTE value increases a little bit slower after 473 K. In the fabricated Cu/diamond composites, the thermal stress exhibits such as tensile stress in the Cu matrix and compressive stress in the diamond. When the specimens are heated from room temperature, the residual stresses generated during fabrication are + +![](images/ac0852538725794acc2958b99b0edfc65c6996dae638a16202d6a5dfce62d7ca.jpg) + +![](images/2dae600f304f246b95991289c856acea52e562f82a84e317900bf89c254dfbd3.jpg) + +![](images/fb456db99362764478cb479ab06af5f424c8ee075b45d0c65605ee9cfdbea5f2.jpg) + +![](images/5c338cfc936ea3a38e431fad28358edcf3e3aef0b36a076f0edaa3ce4302fdac.jpg) +Fig. 6. Interfacial thermal resistance in the Cu-Ti/diamond composites: (a) schematic illustration of Cu-TiC-diamond transition region and (b) a parallel analogue of thermal resistance in the transition region. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.) + +![](images/7fca5c3a50d8b160b33b80b7513923c23cc408b78c79ee44bb22d5e6c8ac0b5c.jpg) +Fig. 7. (a) Coefficient of thermal expansion of the Cu-Ti/diamond composites and thermal linear expansion curve for a heating and cooling cycle of the Cu-0.5 wt.% Ti/diamond composite; (b) comparison of the CTEs between predictions and experimental results. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.) + +released gradually and this is helpful to increase CTE; at higher temperature, the promotion to CTE of residual stress does not occur any more. Therefore, the increasing rate of CTE is lowered between 473 K and 573 K. + +Fig. 7(b) shows both the measured and predicted CTE values for the Cu/diamond composites. Two existing models are applied to calculate the theoretical values of CTE [34,35]: + +(1) The Turner model + +$$ +\alpha_ {c} = \frac {V _ {r} K _ {r} \alpha_ {r} + V _ {m} K _ {m} \alpha_ {m}}{V _ {r} K _ {r} + V _ {m} K _ {m}} \tag {2} +$$ + +(2) The Kerner model + +$$ +\alpha_ {c} = V _ {r} \alpha_ {r} + V _ {m} \alpha_ {m} + V _ {r} V _ {m} \frac {\left(\alpha_ {r} - \alpha_ {m}\right) \left(K _ {r} - K _ {m}\right)}{K _ {r} V _ {r} + K _ {m} V _ {m} + \frac {3 K _ {r} K _ {m}}{4 G _ {m}}} \tag {3} +$$ + +where a, V, K and G are the CTE, volume fraction, bulk modulus and shear modulus, and c, r, m are indexes of composite, reinforcement and matrix, respectively. In this study, $V _ { \mathrm { d i a m o n d } } = 0 . 6 1$ and $V _ { \mathrm { C u } } = 0 . 3 9 .$ . The other parameters [17] used for calculation are $\mathcal { D } _ { \mathrm { d i a m o n d } } = 2 . 3 0 \times 1 0 ^ { - 6 } / \mathrm { K } , \quad \mathcal { D } _ { \mathrm { C u } } = 1 6 . 4 2 \times 1 0 ^ { - 6 } / \mathrm { K } , \quad K _ { \mathrm { d i a m o n d } } = 5 8 0 \mathrm { G P a }$ , $K _ { \mathrm { C u } } = 1 4 0 \mathsf { G P a }$ , $G _ { \mathrm { d i a m o n d } } = 3 6 0 \mathrm { G P a }$ and $G _ { \mathrm { C u } } = 4 9 \mathrm { G P a } .$ . Unlike the Turner model, the Kerner model takes shear strain into consideration. In Fig. 7(b) the two Kerner curves are obtained by reversing the role of copper matrix and diamond reinforcement, since the diamond and the copper have comparable volume fractions in the composites. Without regard to the effect of the third component (alloying element or interphase) on the CTE, the calculations for the Cu/diamond composites are obtained with the three models. As shown in Fig. 7(b), the experimental results are all in good agreement with the values predicted by the Kerner copper matrix model, confirming the presence of shear effects in the composites. The measured CTEs for Cu/diamond composites are apart from the Kerner diamond matrix model. The deviation is understood by the fact that the diamond particles do not form a continuous phase, as shown in Fig. 2(a), although they possess a higher volume fraction than copper phase. The measured CTE for the Cu-0.5 wt.%Ti/diamond composite is below the Kerner copper matrix line. This could be ascribed to the effect of Ti alloying on the shear modulus of the copper matrix. + +The thermal expansion curve is shown in the inset to Fig. 7(a). A residual strain $\left( \varepsilon _ { \mathrm { p } } \right)$ of $0 . 4 0 1 1 2 \times 1 0 ^ { - 3 }$ is derived after a thermal cycle from 323 K to 640 K. The low $\varepsilon _ { \mathrm { { p } } }$ value is acceptable for practical applications of electronic packaging materials. + +# 4. Conclusion + +The Cu-xTi/diamond composites $( x = 0 . 3 , ~ 0 . 5 , ~ 2 . 0 \mathrm { w t } . \% )$ were produced by the gas pressure infiltration method, with a thin and discontinuous layer of TiC formed at the interface. The presence of TiC is favorable to the significant enhancement of interfacial thermal conductance, giving a high thermal conductivity of 752 W/m K at $\chi = 0 . 5 \mathrm { w t . \% }$ . The Cu-0.5 wt.%Ti/diamond composite displays a coefficient of thermal expansion of $6 . 5 0 \times 1 0 ^ { - 6 } / \mathrm { K }$ that is compatible with semiconductors. The finding suggests that Ti alloying in the metal matrix is an effective method to achieving high thermal conductivity in Cu/diamond composites. + +# Acknowledgements + +This work was financially supported by the National Natural Science Foundation of China (No. 51271017, 51301018) and International Science and Technology Cooperation Program of China (No. 2014DFA51610). + +# References + +[1] Zweben C. Advances in high-performance thermal management materials: a review. Covina, CA, ETATS-UNIS: Society for the Advancement of Material and Process Engineering; 2007. +[2] Zweben C. Thermal materials solve power electronics challenges. Power Electron Technol 2006;32:40–7. +[3] Zhang C, Wang R, Cai Z, Peng C, Feng Y, Zhang L. Effects of dual-layer coatings on microstructure and thermal conductivity of diamond/Cu composites prepared by vacuum hot pressing. Surf Coat Technol 2015;277:299–307. +[4] Muhammad Z, Raza K, Khalid FA, Mabrouki T. 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b/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/14_Effect of coating on the microstructure and thermal cond/14_Effect of coating on the microstructure and thermal cond.md @@ -0,0 +1,231 @@ +# Effect of coating on the microstructure and thermal conductivities of diamond–Cu composites prepared by powder metallurgy + +Shubin Ren ⇑ , Xiaoyu Shen, Caiyu Guo, Nan Liu, Jianbing Zang, Xinbo He, Xuanhui Qu + +Institute for Advanced Materials and Technology, University of Science and Technology Beijing, Beijing 100083, China + +# a r t i c l e i n f o + +Article history: +Received 2 September 2010 +Received in revised form 16 June 2011 +Accepted 21 June 2011 +Available online 5 July 2011 +Keywords: +A. Metal–matrix composites +A. Coating +B. Thermal properties +B. Interface +E. Sintering + +# a b s t r a c t + +Diamond–Cu composites from the direct combination of diamond and Cu show low thermal conductivities due to weak interface and high thermal resistance as a result of chemical incompatibility. In this paper, a new method is proposed to strengthen interfacial binding between diamond and Cu by coating strong carbide-forming elements, e.g., Ti or Cr on the surface of the diamond through vacuum microdeposition. Interfacial thermal resistance of diamond–Cu composites is greatly decreased when diamond particles are coated by a Cr or Ti layer of a certain thickness before combining with Cu. Thermal conductivity is also increased several times. Cr coating can reduce more effectively interface thermal resistance between diamond and Cu than Ti coating. Moreover, it has a smaller negative impact on the thermal conductivity of the Cu matrix, resulting in higher thermal conductivity of Cr-coated diamond–Cu composites. Through the vacuum micro-deposition technology, Cr on the diamond particle surface is present in the form Cr7C3 near diamond and a pure Cr outer layer at 2:1. The optimum thickness is within 0.6–0.9 lm; at this depth, the thermal conductivities of 70 vol% diamond–Cu composites can be increased four times and reach as high as 657 W/m K. In this work, an original theoretical model is proposed to estimate the thermal conductivities of composite materials with an interlayer of a certain thickness. The predicted values from this model are in good agreement with the experimental values. + +- 2011 Elsevier Ltd. All rights reserved. + +# 1. Introduction + +Diamond–Cu composites have drawn much attention as new generation heat sink materials in the field of electronic packing because of their potential thermal conductivity higher than 500 W/m K, which is more than twice the value of current SiCp–Al composites or W–Cu alloys, and has lower CTE matching that of the chip [1–3]. Diamond–Cu composites can be prepared by powder metallurgy or metal infiltration; both have been widely used in the production of other heat sink materials, such as SiCp–Al or SiCp–Cu composites [4,5]. However, regardless of preparation techniques, the thermal conductivity of diamond–Cu composites is much lower than that of Cu itself if diamond particles compound directly with Cu. The reason for this is that diamond and Cu are chemically incompatible and their interface shows weak mechanical bonding; as a result, their interface thermal resistance is so high that the composites show low thermal conductivities. In order to reduce the interface thermal resistance between diamond and Cu, some methods must be used to turn their mechanical and physical bonding state into a chemical and metallurgical bonding. Based on a previous study, alloying Cu with minor amounts of carbide + +formers, e.g., Cr or B can improve bonding strength and thermophysical characteristics by forming a thin nano-sized chromium carbide layer at the interface between diamond and Cu [1]. + +In this work, a new method is proposed to solve poor interface binding between diamond and Cu. In the proposed method, the diamond surface is coated with strong carbide-forming elements, e.g., Ti, Cr, Mo or W. These metal elements can act as middle layers that strengthen the interface; in addition, they can also protect the diamond powder from the atmosphere and reduce the degree of graphitization at high temperatures. + +In this paper, both Ti and Cr metals are selected as the surface coating elements of diamond particles for two reasons: (1) the bonding force between Cu and Mo or W are much smaller than that between Cu and Ti or Cr because the solid solution degree of Mo or W in copper is zero; and (2) the process of coating Ti or Cr on the surface of the diamond particle can be carried out by vacuum micro-deposition technology, a method widely used in the manufacture of diamond tools. Although sputtering can be used to coat Ti or Cr metal on the surface of diamond particles, Wang [6] found that the resulting bonding state between diamond and Cr or Ti coating is still physical, and the interface bonding between diamond and Ti or Cr is so weak that Ti or Cr metal layer can be separated from the surface of diamond in the course of the collision with milling balls in the subsequent process of mixing with Cu powder. In this + +case, a second heat treatment must therefore be adopted to strengthen the interface binding by the formation of carbides between diamond and pure Ti or Cr. Magnetron sputtering is believed to be more suitable for W or Mo coating on the surface of the diamond. In comparing the two methods mentioned above, the method to coat Ti or Cr on the surface of the diamond is easier, cheaper, and more suitable for industrial production than that in which W or Mo is used. In this paper, both Ti and Cr are therefore chosen as the coating elements, with which to study the effect of surface metal-coating of diamond particles on the microstructure and thermal conductivities of diamond–Cu composites. + +# 2. Experimental procedures + +# 2.1. Preparation of composite + +The diamond particle used in this work is an MBD4 grade diamond with a nitrogen content of 250 ppm, purchased from Yanshan University, Hebei, China. And its thermal conductivity is estimated to be about 1450 W/m K according to the level of nitrogen content. To obtain a volume fraction of diamond particle in the range of 55–70 vol% in diamond–Cu composites, two kinds of diamonds with designated size, 125–106 lm (mesh 120/140) and 45–38 lm (mesh 325/400) were mixed at a weight ratio of 3 to 1, at which the mean size was about 100 lm and the maximum fraction of diamond can reach as high as 70 vol% of the powder bed. The surface of diamond particle was coated with chromium or titanium of a certain thickness through vacuum micro-deposition technology. And the thickness of Ti or Cr coating can range from several nanometers to ten microns by adjusting the deposition temperature and time, as well as the weight ratio of diamond to coating raw materials such as $\mathrm { T i C l } _ { 3 }$ and $\mathrm { T i H } _ { 2 }$ powders (for Ti coating) or $\mathrm { C r C l } _ { 3 }$ and $\mathrm { C r H } _ { 3 }$ powders (for Cr coating) . In this work, the coating process is performed at the Yanshan University, China, the developer of vacuum microdeposition technology. Electrolytic copper powders with $d _ { 5 0 }$ of 80–90 lm and a purity >99.9% (purchased from General Research Institute for Nonferrous Metals, Beijing, China) were mixed by ball milling with surface-pretreatment diamond powder at a designed volume fraction. This was then synthesized through the spark plasma sintering (SPS) using a sintering pressure of 37–43 MPa and a sintering temperature of 930–950 C for 15–22 min. Given that the sintering process parameters have been optimized in our previous work, the detailed process about SPS are no longer discussed in this paper. + +# 2.2. Property testing + +Cylindrical disk specimens with a diameter of 10 mm and a thickness of 3 mm were produced for the testing of thermal diffusivity and specific heat. Thermal conductivity is specifically calculated as the product of density, thermal diffusivity, and specific heat. In this work, thermal diffusivity and specific heat were measured by Xenon pyrometry with LFA 447 Nanoflash from Netzsch, Germany using the laser flash method and calorimetric techniques, respectively. The bulk density of the composites was measured using a method based on Archimedes’ law and compared with the theoretical density. Microstructure of the composite was observed on the LEO1450 Scanning Electron Microscope (SEM). XRD analysis was carried out on a Siemens D 5000 X-ray Diffractometer using Cu radiation. + +# 3. Results and discussion + +Fig. 1 shows surface morphology of (a) raw diamond particles, (b) Cr-coated diamond particles, and (c) Ti-coated diamond particles; it + +also shows that the surface of diamond particles are evenly coated by Cr or Ti. The designed thickness of Cr or Ti coating layer through vacuum micro-deposition is -1.5 lm. + +Thermal conductivities and relative densities of diamond–Cu composites made from three kinds of diamond powder by SPS are shown in Fig. 2. The thermal conductivity of the uncoated diamond (55 vol%)–Cu composite is 178 W/m K, which is much lower than that of Cu itself (378 W/m K). Coating the diamond with a 1.5 lm thick Ti layer increased the thermal conductivities of the prepared composites to 349 W/m K. When the Ti layer was replaced by a Cr layer, the thermal conductivities of the corresponding composites reached as high as 528 W/m K. The change of relative density of prepared composites in Fig. 2 shows that surface pretreatment of the diamond particle by Ti or Cr raised the density of diamond–Cu composites from 98.2% to 99.5% or higher. + +The thermal conductivities of composites are directly related to the thermal conductivities of both matrix and reinforcement, as well as the microstructure of the composites, especially the interface structure. In this experiment, the three kinds of composites demonstrated a marked difference in interface structure (Fig. 3). + +Fig. 3 shows the interfacial microstructures of the composites made from the compositions corresponding to Fig. 1. The interface gap (shown by the arrow) is obvious between the diamond particles and the Cu matrix; the surface of the diamond is intact, indicating that binding between diamond and Cu is very weak (Fig. 3a). The separation between diamond and copper (Fig. 3a), is generated during the cooling process due to their chemical incompatibility and the different thermal expansion coefficients. Assuming all pores in the composites are considered as gaps between diamond and Cu, and that the shape of the diamond particle is spherical, the thickness of the gap (x) could be estimated from the porosity of composites according to the following equation: + +$$ +\left(\frac {a}{a + x}\right) ^ {3} = \frac {\left(1 0 0 - V _ {2}\right) \cdot V _ {1}}{\left(1 0 0 - V _ {2}\right) \cdot V _ {1} + V _ {2}}, \tag {1} +$$ + +where a and $V _ { 1 }$ are the size and volume fraction of diamond particle, respectively, and V2 is the porosity of corresponding composites. Thus, a density of 98.2% means that the thickness of the gap is about 1 lm, which is close to the observed value by SEM. The diamond–Cu interface became denser when coated by a Ti or Cr layer; at the interface, C and Cu elements diffused into the Ti or Cr interlayer, the Ti or Cr in the interlayer also diffused into the Cu matrix (Fig. 3b and c). To further determine phase composition of these composites, Fig. 4 gives X-ray diffraction patterns of Cr-coated diamond and the corresponding diamond–Cu composites. Given the identical phase change of Cr-coated and Ti-coated diamond, X-ray diffraction patterns of Ti-coated diamond and the corresponding composites are not given in this paper. $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ and Cr phases both appeared in Cr-coated diamond particles; however, the Cr phase disappeared after the formation of composite materials. The same phase variation also appeared in Ti-coated diamond particles and the corresponding composite materials: both TiC and Ti appeared in the Ti-coated diamond particles, and the Ti phase also disappeared in the formed diamond–Cu composites. + +It should be mentioned that although the surface of the diamond particle was coated by a pure Ti or Cr layer, the carbides coexisted in the coating. This special structure is directly related with the chosen plating process. The mechanism of vacuum micro-deposition process can be described in the following reaction equations: + +$$ +\begin{array}{l}2 \mathrm {T i C l} _ {3} (\mathrm {g}) + 3 \mathrm {T i H} _ {2} (\mathrm {s}) \rightarrow 5 \mathrm {T i} (\mathrm {s}) + 6 \mathrm {H C l} (\mathrm {g})\\\mathrm {T i} (\mathrm {s}) + \mathrm {C} (\text {d i a m o n d}) \rightarrow \mathrm {T i C} (\mathrm {s})\end{array}\tag {2} +$$ + +and + +![](images/d4431ede2a6e099e4439ca42650abe8c3d70de32594272b166c4bb5ca6cd0442.jpg) + +![](images/a93253665906aa7da7a5e36e8fee8f2bd7cd249ebbacc242886c64f7fc9cacca.jpg) + +![](images/9fbbf6144a8b7aa4d26ca9971fa601da27efe1512726a37de579cce12a2efc0a.jpg) +Fig. 1. Surface morphology of diamond particles (a: raw, b: Ti-coated, c: Cr-coated). + +![](images/a208a15b82ae8689158e2c2bfad564e34ea5af95afdf2f3da8d48cc77c54fe5c.jpg) +Fig. 2. Thermal conductivity and relative density of diamond (55 vol%)–Cu composites from diamonds with different coating materials. + +$$ +\begin{array}{l}\mathrm {C r C l} _ {3} (\mathrm {g}) + \mathrm {C r H} _ {3} (\mathrm {s}) \rightarrow \mathrm {C r} (\mathrm {s}) + 3 \mathrm {H C l} (\mathrm {g})\\7 \mathrm {C r} (\mathrm {s}) + 3 \mathrm {C} (\text {d i a m o n d}) \rightarrow \mathrm {C r} _ {7} \mathrm {C} _ {3} (\mathrm {s}).\end{array}\tag {3} +$$ + +To coat Ti on the surface of the diamond powder, diamond particles were mixed with $\mathrm { T i C l } _ { 3 }$ and $\mathrm { T i H } _ { 2 }$ powders at a certain proportion; with the increase of temperature, $\mathrm { T i C l } _ { 3 }$ gasified and reduced to Ti atoms by reacting with ${ \mathrm { T i H } } _ { 2 } .$ These generated reactive gaseous Ti atoms deposited on the surface of the diamond particles and then reacted with diamond to form the TiC layer at a plating temperature higher than the forming temperature of TiC. With the increase of plating time, the C atom at the surface of the diamond particle diffused into the TiC layer and continued to react with the Ti atoms through Eq. (2). In the process, diffusion of C atoms and deposition of Ti atoms occurred simultaneously; because the former was slower than the latter, a pure Ti layer remained on the TiC layer. The Ti-coated layer on the surface of diamond actually consisted of a TiC layer and an outer layer of pure + +Ti. Wang [6] and Li [7] studied the thickness ratio of two layer and lattice relationship by small angle X-ray diffraction and TEM, respectively; they found that the thickness ratio of the TiC layer and the Ti layer was about 2:1 and a coherent interface of (1 0 0 0)Ti//(1 1 1)TiC existed between Ti layer and TiC layer. Cr coating has a similar layer structure: the layer is made up of a $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ layer close to the diamond surface and a pure Cr outer layer. Based on these results, the thickness of carbide layer was 1 lm, and the pure Ti or Cr layer was 0.5 lm thick in this experiment. + +A middle layer with a strong binding force with the diamond and copper caused the interface structure of diamond–Cu composites to change from mechanical bonding into chemical, metallurgical bonding; at the same time, the composites also reached full density. The interface layer of prepared diamond–Cu composites can be described by the following forms (according to the different surface states of diamond): + +diamond|air gap|Cu (for raw diamond); + +diamond|TiC|copper solid solution containing titanium (for Ticoated diamond); and + +diamond|Cr7C3| copper solid solution containing chrome (for Cr-coated diamond). + +It should be noted that pure Ti or Cr was not added to the interface structure. The composition of Cu matrix was analyzed using energy scattering spectroscopy, and the matrix contained 1.8 wt% Ti or 2.2 wt% Cr. These values are close to the theoretical estimates of 1.9 wt% Ti or 2.6 wt% Cr in the Cu matrix if a pure 0.5 lm thick Ti or Cr layer on the surface of the diamond particle completely dissolves in copper. Moreover, the XRD result in Fig. 4 also shows that no pure metal coating existed in the composites. Based on both facts, we did not add pure Ti or Cr onto the interface structure. + +The resulting layer structure is what the authors aim for; it is also the reason for using the vacuum micro-deposition process instead of sputtering and other processes. The Cr or Ti layer formed by vacuum micro-deposition as an interlayer could enable the formation of + +![](images/dafb20d796b165e29bd04a337ed78722a8a6c263c00ae0d1ce314bc85ec107c0.jpg) + +![](images/b5e39e57c683429b41ec55821e6d24088433db8882333fe7ecf5a6e416519b55.jpg) + +![](images/9635895a85eee576082cb00af92bb43da8303583da07d8b733a99fbc68f1f20b.jpg) +Fig. 3. SEM of diamond–Cu composites made from diamond with different surface treatments (a: raw diamond, b: Ti-coated, c: Cr-coated). + +![](images/5a6e7a837734422cd7676a8ccf417c011d1114a591f28def812a636b778dec1e.jpg) +Fig. 4. XRD of Cr-coated diamond (a) and corresponding composites (b). + +![](images/2d9869a8af3c5bae8c088f9df8064ab946330da6e5088137989febe06ed9a9fb.jpg) +Fig. 5. Composite schematic diagram of Ti-coated or Cr-coated diamond and Cu. + +metallurgical bonding between diamond and Cu in two ways: by the solution of pure Ti or Cr in the Cu matrix and by the formation of a metal carbide layer, (Fig. 5). Wang [6] has researched on the interface between diamond and coated-Ti or Cr by Magnetron sputtering, and found that if sputtering is used to coat Ti or Cr on the diamond, the surface is going to be coated by pure Ti or Cr layer with no carbide layer; a second heat treatment must therefore be adopted to form carbides between diamond and pure Ti or $\mathsf { C r } ,$ making the coating process complex and difficult to control. + +After analyzing the interface structure, we present in the following section a theoretical argument on the change of thermal conductivities of diamond–Cu composites with the surface state of diamond. + +The thermal conductivities of the composites depend on the thermal conductivity of each component: volume fraction, distribution and size of the reinforcement, density, and bonding between matrix and reinforcement. Many researchers have constructed theoretical models to describe the impact of these factors on the thermal conductivities of composites. Of these, the H–J model is believed to be most accurate because it takes into account the combined + +Table 1 The values of the parameters and calculated thermal resistance. + +
Materialρ(g/cm3)G(GPa)C(J/kg K)λ(W/m K)R×10-8m2KW-1
Diamond[12]3.52504.45151450Rdiamond|TiC=0.15
Cua8.9642.9381.5381RTiC(1μm)=5.98
TiC[12]4.93208.956416.7RTiC|Cu-1.8wt%Ti=0.92
Cr7C3[12]6.92148.454319.11Rdiamond|Cr7C3=0.31
Cu-1.8 wt%Tiα8.8445.2382189RCr7C3(1μm)=5.23
Cu-2.2 wt%Crα8.9152.7381.7370RCr7C3|Cu-2.2wt%Cr=0.36
+ +a Experimental values. + +effects of particle size, volume fraction, and interfacial thermal resistance; the model [8] is as follows: + +$$ +\lambda_ {c} = \lambda_ {m} \left[ \frac {2 \left(\lambda_ {p} / \lambda_ {m} - \lambda_ {p} R _ {c} / a - 1\right) V _ {p} + \lambda_ {p} / \lambda_ {m} + 2 \lambda_ {p} R _ {c} / a + 2}{\left(1 - \lambda_ {p} / \lambda_ {m} + \lambda_ {p} R _ {c} / a\right) V _ {p} + \lambda_ {p} / \lambda_ {m} + 2 \lambda_ {p} R _ {c} / a + 2} \right], \tag {4} +$$ + +where $\lambda _ { c }$ refers to the thermal conductivities of the composites, W $\mathfrak { m } ^ { - 1 } \mathfrak { K } ^ { - 1 } ; \lambda _ { m }$ is the thermal conductivity of the matrix, $\mathrm { { w } } \mathrm { { m } } ^ { - 1 } \mathrm { { K } } ^ { - 1 } ;$ ; $\lambda _ { p }$ is the thermal conductivity of the reinforcement, W m $^ { - 1 } \mathsf { K } ^ { - 1 } ; V _ { p }$ is the volume fraction of the reinforcement (vol%); a is the size of the reinforcement particle (m); and Rc is interfacial thermal resistance $( \mathrm { m } ^ { 2 } \mathrm { K W } ^ { - 1 } )$ . For the three kinds of diamond–Cu composites in this experiment, the values of their $\lambda _ { p } , V _ { p } ,$ and a are the same, but $\lambda _ { m }$ and hc are different. If the diamond particles are directly composited with copper, $\lambda _ { m }$ value in the prepared composites becomes the thermal conductivity of pure copper, 381 $\mathrm { ~ w ~ m ~ } ^ { \setminus 1 } \mathrm { K } ^ { - 1 }$ . If the diamond particle is coated with a chromium or titanium layer with a thickness of 1.5 lm using vacuum micro-deposition, a pure, 0.5 lm thick layer on the outer surface of diamond becomes completely soluble in the copper matrix, and the $\lambda _ { m }$ value for the Ti-coated diamond–Cu composites becomes the thermal conductivity of Cu–1.8 wt%Ti alloys, which is 189 $\mathsf { W } \thinspace \mathrm { m } ^ { - 1 } \thinspace \mathrm { K } ^ { - 1 }$ by experimental measurement. The $\lambda _ { m }$ value of Cr-coated diamond–Cu composites is the thermal conductivity of Cu–2.2 wt%Cr alloys, which is 370 W $\mathrm { m } ^ { - 1 } \mathrm { K } ^ { - 1 }$ by experimental measurement. Rc cannot be measured from the experiment. In terms of the Acoustic Mismatch Theory, the interfacial thermal resistance Rc can be estimated by a simple Debye model [9,10] expressed by: + +$$ +R c = \frac {4}{\rho_ {m} \cdot C _ {m} \cdot v _ {m} \cdot \eta}, \tag {5} +$$ + +where $\rho _ { m }$ is the density, $C _ { m }$ the specific heat, $\nu _ { m }$ the Debye velocity of the matrix, and g is the average probability for the transmission of phonons across the interface into the particles. With the assumption that transverse phonons carry most of the heat, the value of $\nu _ { m }$ can be estimated by [11]: + +$$ +v _ {m} = \sqrt {G _ {m} / \rho_ {m}}, \tag {6} +$$ + +where $G _ { m }$ is the shear modulus of the matrix. The value of g can be estimated by [9,10] using: + +$$ +\eta = p \cdot q, \tag {7} +$$ + +where p is the transmitted probability of a phonon incident within the critical angle $\theta _ { c } ,$ which can then be calculated by [9,10] using: + +$$ +p = \frac {4 Z _ {m} \cdot Z _ {p}}{\left(Z _ {m} + Z _ {p}\right) ^ {2}}. \tag {8} +$$ + +In the above, $Z _ { m } , Z _ { p }$ is the acoustic impedance of the matrix and reinforcement, which can be calculated by $Z = \rho \cdot \nu .$ + +In Eq. (7), q is the fraction of the phonons incident on the interface within the critical angle h from the matrix side, because only these phonons have the opportunity of being transmitted: + +$$ +q = \frac {1}{2} \sin^ {2} \theta_ {c} = \frac {1}{2} \left(\frac {v _ {m}}{v _ {p}}\right) ^ {2}. \tag {9} +$$ + +Table 2 Measured values of thermal diffusivity, specific heat, density, relative density of all prepared diamond–Cu composites, as well as the preparation conditions. + +
Raw materialsThickness of Cr7C3 layer (μm)Vp(vol%)Sintering pressure (MPa)/ temperature (°C)/time (min)Density (g/cm3)Relative density (%)Specific heat (J/g K)Thermal diffusivity (mm2/s)Thermal conductivity (W/m K)
Diamond (D1:126-106 μm; D2:45-38 μm) D1:D2 = 3:1 + Cu powder05537/930/155.86198.20.44069.02178
0.25537/930/155.86898.40.44087.14225
6037/940/155.60698.50.44596.61241
7037/950/155.06798.40.452116.58267
0.55537/930/175.94799.80.441224.20588
6037/940/175.64099.50.445242.25608
7037/950/175.12199.50.452283.84657
1.05540/930/205.91699.50.442201.92528
6040/940/205.65399.60.446216.16545
7040/950/205.11699.60.451256.14591
1.55543/940/225.91399.60.440180.65470
6043/940/225.64099.50.445192.04482
7043/950/225.10599.50.453219.67508
2.05543/950/225.90499.60.441165.15430
6043/950/225.63799.60.445173.41435
7043/950/225.10599.60.453190.26440
+ +Thus, Eqs. (5)–(9) can be integrated into Eq. (10) as expressed by: + +$$ +R c = \frac {2 \left(\rho_ {m} v _ {m} + \rho_ {p} v _ {p}\right) ^ {2}}{C _ {m} \cdot \rho_ {m} ^ {2} \cdot v _ {m} ^ {2} \cdot \rho_ {p} \cdot v _ {p}} \left(\frac {v _ {p}}{v _ {m}}\right) ^ {2}. \tag {10} +$$ + +In this experiment, two interfaces and a middle layer in the prepared diamond–Cu composites have been found to exist. Therefore, Rc is the sum of three thermal resistances in series, that is, for Ti-coated diamond–Cu composites, Rc is signed by $\scriptstyle R _ { \mathrm { d - T - C } } , R _ { \mathrm { d - T - C } } =$ Rdiamond|TiC + RTiC + RTiC|Cu–1.8wt%Ti, and for Cr-coated diamond–Cu composites, Rc is signed by $R _ { \mathrm { d - C - C } } ,$ Rd-C–C = Rdiamond|Cr7C3 $+ R _ { \mathrm { C r 7 C 3 } } + R _ { \mathrm { C r 7 C 3 } | \mathrm { C u - 2 . 2 w t \% C r . } }$ . When Rdiamond|TiC or Rdiamond|Cr7C3 is calculated, we can assume that the carbide is the matrix and diamond is the reinforcement. Moreover, when $R _ { \mathrm { T i C | C u - } }$ 1.8wt%Ti or $R _ { \mathrm { C r 7 C 3 | C u - 2 . 2 w t \% C r } }$ is calculated, TiC or $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ is the reinforcement, and the Cu alloy is the matrix. Values for $R _ { \mathrm { T i C } }$ or $R _ { \mathsf { C r 7 C 3 } }$ are from the following equation: + +$$ +R = b / \lambda , \tag {11} +$$ + +where b is the thickness of the carbide layer, and k is thermal conductivity of the carbide.Table 1 lists all values of parameters in Eq. (10). And the thickness value (b) of both $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ and TiC layer is 1 lm. By substituting these parameters into Eq. (10) and (11), $R _ { \mathrm { d i a m o n d } | \mathrm { T i C } } ,$ + +![](images/b7de8f3f2442ab28d3a427b14dca49f172b1e96daf3850440379a25f0d7dd2e2.jpg) +Fig. 6. Theoretical curves and experimental values of the thermal conductivities of the composites from diamonds coated by 1 lm $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ or 1 lm TiC layer. + +$R _ { \mathrm { T i C } } , R _ { \mathrm { T i C | C u - 1 . 8 w t \mathcal { U } T i } } , R _ { \mathrm { d i a m o n d | C r 7 C 3 } } , R _ { \mathrm { C r 7 C 3 } } ,$ and $R _ { \mathrm { C r 7 C 3 | C u - 2 . 2 w t \% C r } }$ can all be obtained (Table 1). Thus, for Ti-coated diamond–Cu composites, $R _ { \mathrm { d - T - C } } = 7 . 0 5 \times 1 0 ^ { - 8 } \mathrm { m } ^ { 2 } \mathrm { K W } ^ { - 1 } ;$ for Cr-coated diamond–Cu composites, $R _ { \mathrm { d - } C - C } = 5 . 9 \times 1 0 ^ { - 8 } \mathrm { m } ^ { 2 } \mathrm { K } \mathrm { W } ^ { - 1 }$ . Given that the bonding between raw diamond and Cu is mechanical, Eq. (10) is not suitable for theoretical calculations of Rc of the composites. We can approximate the thermal resistance of air layer (1 lm in thickness) as the interface thermal resistance Rc, signed by $R _ { \mathbf { d - A - C } } .$ Eq. (11) can be used to calculate $R _ { \mathbf { d } - \mathbf { A } - \mathbf { C } } ,$ , where b = 1 lm and k is the thermal conductivity of air (0.023 W/m K); $R _ { \mathrm { d - A - C } } = 4 . 3 \times 1 0 ^ { - 5 } \mathrm { \ m } ^ { 2 } \mathrm { K } \mathrm { W } ^ { - 1 }$ . Thus, by substituting every interface thermal resistance into Eq. (4), theoretical values of the thermal conductivities of the composites can be obtained. Fig. 6 shows the theoretical curve based on Eqs. (4) and (10) of the thermal conductivities of the composites as well as the experimental values. The experimental value is close to the theoretical one, indicating that the calculation models are reasonable. + +As proven by analysis, Cr coating on the diamond surface can reduce more effectively the interface thermal resistance between diamond and Cu than Ti coating; it also has a smaller negative impact on the thermal conductivity of the Cu matrix. The Cr-coated diamond–Cu composites exhibit higher thermal conductivities. In addition, we have also shown the effect of the carbide layer thickness on the surface of diamond on the thermal conductivity of the + +![](images/74676c1fb4a25a2acb2cd3df952c3986036773f73e3bb634ae009af166423a62.jpg) +Fig. 7. The change of theoretical and experimental values of the thermal conductivities of Cr-coated diamond–Cu composites with different volume fractions of diamond with the thickness of $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ layer on the surface of the diamond particle. + +Table 3 Interfacial thermal resistance of diamond/Cu composites prepared by different methods. + +
Raw materialPreparation methodInterfacial thermal resistance (m2K W-1)References
Cu/diamondP/M(SPS, 37 MPa /930 °C/15 min)190 × 10-8In this work
Cu/diamondP/M(conventional hot pressing, 950 °C, 60 min)200 × 10-8[1]
Cu/diamondP/M (ultra-high pressure, 4.5 GPa, 1180 °C, 15 min)3.36 × 10-8[13]
Cu/Cr-coated diamondP/M(SPS, 37 MPa, 950 °C, 17 min)4.957 × 10-8In this work
CuCr/diamondP/M(directly heated hot pressing with heating/cooling rates of 100–150 K/min, 950 °C, 30 min)1 × 10-8[1]
+ +diamond–Cu composite. In this work, the thickness of $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ layer on the surface of diamond can be accurately controlled by comprehensive adjustments of coating temperature, time and the weight ratio of diamond to $\mathrm { C r C l } _ { 3 }$ and CrH3 powders. Table 2 summarizes all experimental values of thermal diffusivity, specific heat, density of all prepared Cr-coated diamond–Cu composites samples. + +Fig. 7 shows the change and comparison of theoretical and experimental values, based on Eqs. (4), (10), and (11), of the thermal conductivities of Cr-coated diamond–Cu composites with the thickness of the $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ layer. Change of the theoretical curve means that the decrease of $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ layer thickness helped increase the thermal conductivities of the composites. However, observing the experimental value, it is close to the theoretical one except that the thickness of $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ layer is 0.2 lm. When the thickness of $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ layer became less than 0.5 lm, the thermal conductivity of the composite dropped greatly due to the fact that $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ became too small to effectively combine diamond with Cu. The interface binding also became weak so the interface thermal resistance remained high, resulting in a composite with a very low thermal conductivity. In addition, with a high thickness of the $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ layer, the thermal conductivity of the composite became low because $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ itself had a low thermal conductivity that became detrimental to interfacial heat transfer. Under the premise of ensuring the ability to form a chemical combination between diamond and $\mathsf { C u } ,$ the thickness of the $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ layer should be as small as possible. From the shadow region in Fig. 6, the optimum thickness for the $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ layer ranged from 0.4 to 0.6 lm, indicating that the thickness of the chromium compound layer (composed of $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ and Cr at 2:1 on the diamond surface by vacuum micro-deposition technology) should be controlled within $0 . 6 \substack { - 0 . 9 \mu \mathrm { m } }$ , where the thermal conductivity of the diamond–Cu composites with 70 vol% diamond can reach as high as 657 W/m K, and the interface thermal resistance can be estimated to be about $4 . 9 5 7 \times 1 0 ^ { - 8 } { \mathrm { m } } ^ { 2 } { \mathrm { K W } } ^ { - 1 }$ . Table 3 summarizes interfacial thermal resistance of diamond–Cu composites prepared by various methods, reported in the recent literatures, as well as our work. Observed from Table 3, it is obvious that thermal resistance between diamond and Cu relies heavily on raw material and/or preparation methods. Although Cr metal layer (actual as chromium carbide) is adopted as the transition layer between diamond and Cu in [1] and our work, the thermal resistance in our work is higher than that in [1]. As a result of lower thermal resistance, the thermal conductivity in [1] is near to that in this work although the volume fraction of diamond is only 42 vol%, much lower than 70 vol% in our work. This may be attributed to the different formation ways of chromium carbide between diamond and Cu. The in situ formation of carbide can result in thinner layers with lower interfacial thermal resistances so that CuCr/diamond composites in [1] exhibited higher thermal conductivity. + +# 4. Conclusion + +(1) Diamond–Cu composites from direct combinations of diamond and Cu showed low thermal conductivities due to a weak interface and high thermal resistance resulting from chemical incompatibility. The diamond–Cu interface was + +strengthened when the surface of the diamond particle was coated by Cr or Ti through vacuum micro-deposition technology before forming the compound with Cu. As a result, interfacial thermal resistance in prepared diamond– Cu composites became greatly reduced and their thermal conductivities increased several times. + +(2) Cr-coating on the surface of the diamond reduced more effectively the interface thermal resistance between diamond and Cu than Ti coating; moreover, it had a smaller negative impact on the thermal conductivity of the Cu matrix. Therefore, Cr-coated diamond–Cu composites exhibited higher thermal conductivities than Ti-coated diamond– Cu composites. +(3) Cr-coating on the diamond particle by vacuum microdeposition technology was present in the form of $\mathrm { C r } _ { 7 } \mathrm { C } _ { 3 }$ near diamond and a pure Cr outer layer, whose ratio in thickness was 2:1. The optimum thickness of the Cr layer fell within 0.6–0.9 lm, at which the thermal conductivities of 70 vol% diamond–Cu composites reached as high as 657 W/m K, making them very interesting materials for electronic substrate materials. + +# Acknowledgment + +This study was financially supported by the National Natural Science Fund of P.R. China, No. 51004010. + +# References + +[1] Schubert Th, Trindade B, Weißgärber T, Kieback B. Interfacial design of Cubased composites prepared by powder metallurgy for heat sink applications. Mater Sci Eng A 2008;475(1–2):39–44. +[2] Weber L, Tavangar R. On the influence of active element content on the thermal conductivity and thermal expansion of Cu–X (X = Cr, B) diamond composites. Scripta Mater 2007;57(11):988–91. +[3] Schubert T, Ciupin´ ski Ł, Zielin´ sk W, Michalski A, Weißgärber T, Kieback B. Interfacial characterization of Cu/diamond composites prepared by powder metallurgy for heat sink applications. Scripta Mater 2008;58(4):263–6. +[4] Pech-Canul MI, Katz RN, Makhlouf MM. Optimum conditions for pressureless infiltration of SiCp performs by aluminum alloys. J Mater Process Technol 2000;108(1):68–77. +[5] Lee HS, Hong SH. Pressure infiltration casting process and thermophysical properties of high volume fraction SiCp/Al metal matrix composites. Mater Sci Technol 2003;19(8):1057–64. +[6] Wang YH. Preparation, structure, properties and application of titanium coating on diamond abrasive. Dissertation for the doctoral degree in engineering, Yanshan University, China; 2002. +[7] Li GB, Wang TB, Deng CY. Ti coating on diamond and interface microstructure. Diamond Abrasives Eng 2007(2):14–7. +[8] Hasselman DHP. Effective thermal conductivity of composites with interfacial thermal barrier resistance. J Compos Mater 1987;21(6):508–15. +[9] Every AG. The effect of particle size on the thermal conductivity of ZnS/ diamond composites. Acta Metall Mater 1992;40(1):123–9. +[10] Swartz ET. Thermal boundary resistance. Rev Mod Phys 1989;61(3):605–68. +[11] Wang JJ, Yi XS. Effects of interfacial thermal barrier resistance and particle shape and size on the thermal conductivity of AlN/PI composites. Compos Sci Technol 2004;64(10–11):1623–8. +[12] Yu QX. On the historical development and future prospect of cutting tool material. Mech Eng 2002(1):9–12. +[13] Yoshida K, Morigami H. Thermal properties of diamond/copper composite material. 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a/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/15_The Influence of the Carbide-Forming Metallic Additives (W, Mo, Cr, Ti) on the Microstructure and Thermal Conductivity of Copper-Diamond Composites/15_The Influence of the Carbide-Forming Metallic Additives (W, Mo, Cr, Ti) on the Microstructure and Thermal Conductivity of Copper-Diamond Composites.md b/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/15_The Influence of the Carbide-Forming Metallic Additives (W, Mo, Cr, Ti) on the Microstructure and Thermal Conductivity of Copper-Diamond Composites/15_The Influence of the Carbide-Forming Metallic Additives (W, Mo, Cr, Ti) on the Microstructure and Thermal Conductivity of Copper-Diamond Composites.md new file mode 100644 index 0000000..dc1d43a --- /dev/null +++ b/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/15_The Influence of the Carbide-Forming Metallic Additives (W, Mo, Cr, Ti) on the Microstructure and Thermal Conductivity of Copper-Diamond Composites/15_The Influence of the Carbide-Forming Metallic Additives (W, Mo, Cr, Ti) on the Microstructure and Thermal Conductivity of Copper-Diamond Composites.md @@ -0,0 +1,270 @@ +Article + +# The Influence of the Carbide-Forming Metallic Additives (W, Mo, Cr, Ti) on the Microstructure and Thermal Conductivity of Copper–Diamond Composites + +Arina V. Ukhina 1,*, Dina V. Dudina 1,2 , Maksim A. Esikov 2 , Dmitrii A. Samoshkin 3 and Sergei V. Stankus 3 + +1 Institute of Solid State Chemistry and Mechanochemistry, Siberian Branch of the Russian Academy of Sciences, Kutateladze Str. 18, Novosibirsk 630128, Russia +2 Lavrentyev Institute of Hydrodynamics, Siberian Branch of the Russian Academy of Sciences, Lavrentyev Ave. 15, Novosibirsk 630090, Russia +3 Kutateladze Institute of Thermophysics, Siberian Branch of the Russian Academy of Sciences, Lavrentyev Ave. 1, Novosibirsk 630090, Russia +Correspondence: auhina181@gmail.com; Tel.: +7-(383)-332-40-02 + +Abstract: In this study, carbide-forming metallic additives $( \mathrm { W } , \mathrm { M o } , \mathrm { C r } , \mathrm { T i } )$ were introduced into the copper matrix to improve the wettability of diamond particles in the copper–diamond composites. The samples were prepared by Spark Plasma Sintering (SPS) and Hot Pressing (HP) at $9 2 0 ^ { \circ } \mathrm { C } .$ . The phase composition, microstructure and thermal conductivity of the samples were investigated. The influence of the carbide-forming additive concentration, the sintering method as well as the nature of the metal introduced into the copper matrix on the thermal conductivity of copper–diamond composites was determined. Titanium ensured a more significant wettability improvement at the copper–diamond interface. This is due to its higher solubility in copper in comparison with other metals (W, Mo, Cr) and the possibility of its diffusion through the copper matrix to the diamond surface resulting in the formation of a closer contact at the copper–diamond interface. + +![](images/5395d6df6e61078e36c985ff4f55a241fa6adc1a896e29336450335216466f83.jpg) + +check for updates + +Citation: Ukhina, A.V.; Dudina, D.V.; + +Esikov, M.A.; Samoshkin, D.A.; + +Stankus, S.V. The Influence of the + +Carbide-Forming Metallic Additives + +(W, Mo, Cr, Ti) on the Microstructure and Thermal Conductivity of + +Copper–Diamond Composites. J. + +Compos. Sci. 2023, 7, 219. + +https://doi.org/10.3390/ + +jcs7060219 + +Academic Editor: Francesco + +Tornabene + +Received: 21 April 2023 + +Revised: 10 May 2023 + +Accepted: 25 May 2023 + +Published: 26 May 2023 + +![](images/72f3fb6e53cfdceb0430daf7add8cd2f436c78a87c0cf51ce031d32cb7b08c33.jpg) + +Copyright: © 2023 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/). + +Keywords: diamond; copper; thermal conductivity; microstructure; scanning electron microscopy + +# 1. Introduction + +The rapid development of microelectronics and the industry requires new efficient heat sink materials to increase the performance of devices without increasing their size. The standard materials commonly used as heat sinks (copper, thermal conductivity 400 $\mathbf { W _ { \lambda } m ^ { - 1 } K ^ { - 1 } } .$ , and aluminum, thermal conductivity 250 $\hat { \mathsf { W } } \thinspace \mathrm { m } ^ { - 1 } \thinspace \mathrm { K } ^ { - 1 } )$ may no longer be efficient. Materials with higher thermal conductivities and lower coefficients of thermal expansion are sought. A promising candidate for high-performance heat sink elements is synthetic diamond, which has a higher thermal conductivity compared with copper [1,2] and is less expensive than natural diamond. Determining factors governing the distribution of solid-phase transformation products is a key problem pertaining to the development of the synthesis methods of composite materials. The solution of this problem allows obtaining composites with controllable morphological characteristics of the components and improved properties. The high hardness and thermal conductivity of diamond, as well as its chemical inertness, allow us to use it as a component of materials for abrasive processing, cutting tools, heat dissipation materials and as inert substrates. However, the crystals of synthetic diamond micropowders cannot be consolidated into bulk objects. This makes it impossible to produce heat sink plates or other products directly from the micropowders. The development of metal–diamond composites with high thermal conductivity requires solving a number of scientific problems, the main one being the elucidation of the mechanisms of phase, structural and morphological transformations at the metal–diamond interface. Those transformations determine the wettability of diamond by the metal during + +the formation of a composite material. In this connection, in order to fabricate metal– diamond composites with high thermal conductivity, it is necessary to solve the problem of low wettability of diamond particles by the metal in order to minimize the thermal resistance at the metal–diamond interface and to prevent the formation of voids, which cause significant deterioration of the properties of the composites. To increase the metal-matrix adhesion to the diamond surface, two main approaches are used: preliminary modification of the diamond surface (formation of thin coatings) and addition of an alloying element directly into the metal matrix. + +Recent studies have been devoted to the preparation of copper–diamond and aluminum–diamond composites with high thermal conductivity [3–7]. Copper has a higher thermal conductivity than aluminum and does not form carbides, which can create an additional barrier to heat transfer. Unfortunately, copper has a poor wettability towards the diamond surface, which causes the pore formation at the copper–diamond interface and reduced heat transfer. For this reason, the thermal conductivity of copper–diamond composites can be lower than that of pure copper [8–10]. + +To overcome this problem, several approaches have been used: the preliminary modification of the diamond surface prior to sintering [11–13], the direct modification of the copper matrix with carbide-forming metals to react in situ [14–18], and the optimization of the synthesis parameters of the composites [7,19]. In [11], the effect of Ti- and W-containing coatings of different thicknesses (50, 100 and 150 nm) on the surface of diamond obtained by magnetron sputtering on mechanical properties (bending strength) and thermal conductivity of copper–diamond composites was described. This study shows that the strengthening mechanisms depend largely on the metal, from which the coating is formed. Thus, when Ti is used as a coating, the bending strength of the material gradually increases with increasing coating thickness but the thermal conductivity first increases and then decreases. In contrast, when W is used as a coating, the bending strength is lower than in the case of a Ti coating of the same thickness. In [12], the use of Ti coatings (thickness 50 nm) on the diamond surface is discussed. Copper–diamond composites were obtained at different temperatures $( 8 0 0 ^ { \circ } \mathrm { C }$ and $1 0 5 0 ^ { \circ } \mathrm { C } )$ by hot forging of cold-pressed powder pre-forms. It was found that, at $8 0 0 ^ { \circ } \mathrm { C } ,$ , the roughened coating surface promotes better adhesion to the copper matrix. In [13], the use of Ti-coated diamonds with coating thicknesses from 65 nm to 850 nm produced by magnetron sputtering was described. The thermal conductivity of the copper–diamond composites in that case nonlinearly depended on the coating thickness. First, the thermal conductivity increased but then decreased with the coating thickness. In [14], the effect of the concentration of the chromium additive on the thermal and mechanical properties of copper–diamond composites was investigated. The composites were obtained by field-assisted sintering. The thermal conductivity of the composites reached a maximum at a chromium concentration of 2 wt.% (339.2 W $\mathrm { m } ^ { - 1 } \mathrm { K } ^ { - 1 } )$ and, with a further increase in the chromium content, the thermal conductivity decreased. The authors relate this effect to the fact that initially a coating consisting of $\mathrm { C r } _ { 3 } \mathrm { C } _ { 2 }$ is formed on the surface of the diamond particles, which promotes their wetting by the copper matrix. The introduction of higher chromium contents causes a reduction in the thermal conductivity of the matrix and the composite as a whole. In [17], a significant increase in the thermal conductivity was observed, when zirconium was added to the copper matrix (930 $\mathbf { W _ { \lambda } m ^ { - 1 } K ^ { - 1 } } )$ . The samples were obtained by a gas pressure infiltration route. The zirconium concentration in the samples was between 0 and 1 wt.% Zr. The authors also explained an increased thermal conductivity by the formation of a ZrC layer on the surface of diamond particles. The authors of [20] reported a thermal conductivity of 688 W $\mathbf { m } ^ { - 1 } \mathbf { K } ^ { - 1 }$ for a composite containing 0.3 wt.% B added to the copper matrix (the composites were obtained by the high-pressure high-temperature method). Close thermal conductivity values were obtained by adding 0.4 wt.% Cr and 1 wt.% Ti. The use of elements that form carbides $( \mathrm { T i } , \mathrm { S i } , \mathrm { C r } , \mathrm { Z r } )$ to modify the copper matrix was described in [20–23]. + +It is important to note that the additive metals differ in their solubility in copper. Aluminum, titanium and silicon have a high solubility in copper [7], which allows them to + +diffuse to the boundary between the matrix with the diamond and form coatings. This leads to the improvement of thermal and mechanical properties of the composites. However, with an excess of additives in the matrix, the opposite effect occurs: the total thermal conductivity of the composite may be reduced because of a reduction in the thermal conductivity of the matrix. Metals such as tantalum, molybdenum and tungsten, although forming carbides, are not soluble in copper, so in this case coatings on diamond can be formed upon a direct contact of these metals with diamond. + +Various methods have been described for the preparation of composites: pulse plasma sintering (PPS) [24], Spark Plasma Sintering (SPS) [25,26] and Hot Pressing (HP) [26,27]. As different concentrations of metals (additives) in the copper matrix were used in different publications, it is difficult to compare the effectiveness of metals at improving the thermal conductivity of the composites. + +In this paper, the effect of additives (soluble (Ti) and insoluble (Cr, Mo and W) in copper) introduced into the copper matrix on the thermal conductivity of copper–diamond composites obtained by the SPS and HP methods was studied. + +# 2. Materials and Methods + +Powders of synthetic diamonds (MBD10, China, particle size 100 µm) and copper (PMS-1, Russia, average particle size 40 µm) were used as raw materials. Ti (PTM-1, Russia), Cr (PH1M, Russia), Mo (MFP, Russia), W (Russia) were used as additives to modify the copper matrix. + +The powder mixtures containing 0.15, 0.35, 0.7 and 2 vol.% of the additive (W, Mo, Cr, Ti) were thoroughly blended in a mortar. The diamond concentration in all samples was 50 vol.%. To produce copper–diamond composites in the bulk state, the SPS and HP methods were used. SPS experiments were carried out on a Labox 1575 facility (SINTER LAND Inc., Nagaoka, Japan) at 920 ◦C with holding times of 3 min and 10 min. A custombuilt set-up (The Institute of Automation and Electrometry SB RAS, Novosibirsk, Russia) was utilized to conduct HP experiments. HP was conducted at 920 ◦C with a holding time of 15 min. The temperature during both SPS and HP experiments was measured using a pyrometer focused on the die wall at its mid-plane. The difference between the two synthesis methods is that, in the SPS, the sample is heated by electric current directly passing through the graphite die and the sample, whereas in the HP, the sample and the die are heated by radiation from the external heating elements. + +A uniaxial pressure of 40 MPa was consistently applied to the samples. The consolidation of composites during HP was performed in argon at a pressure of 0.1 MPa, while the SPS experiments were carried out under the forevacuum conditions (at a residual pressure of 10 Pa). The heating rate in both methods was 50 ◦C min−1. To protect the inner walls of the die and the flat ends of the punches, a layer of graphite foil was used. + +Examinations of the microstructure of the copper–diamond composites by scanning electron microscopy were conducted using a Hitachi TM-1000 Tabletop Microscope (Japan). Each composite sample was fractured and the fracture surface was examined. The X-ray diffraction (XRD) technique was used to investigate the phase composition of the samples. The XRD patterns were acquired using a D8 ADVANCE X-ray diffractometer (Bruker AXS, Germany). The step size was 0.02◦ 2θ for all samples. The XRD patterns were recorded from the flat surface of the sintered samples. Additional analyses were carried out for samples with 2 vol.% titanium. A piece of sample was mechanically ground in a mortar to extract the diamond particles from the compacts. The obtained diamond particles were placed on a single-crystal Si-cuvette to record a XRD pattern of these particles. The Rietveld refinement technique implemented in the TOPAS 4.2 software (Bruker AXS, Germany) was used to determine the parameters of the copper crystalline structure. + +The laser flash method was used to measure the thermal diffusivity (a) of the copper– diamond composites, using an LFA-427 setup (Netzsch, Germany). An average value for + +the temperature diffusivity coefficient was obtained by five measurements. The thermal conductivity (λ) of the composites was then calculated as follows: + +$$ +\lambda = a \rho C p, +$$ + +In this equation, a represents the thermal diffusivity, ρ represents density, and CpIn this equation, a represents the thermal diffusivity, ρ represents density, and Cp represents the specific heat capacity of the composite. The specific heat capacity of therepresents the specific heat capacity of the composite. The specific heat capacity of the composite was estimated using the rule of mixtures, employing the values of pure coppercomposite was estimated using the rule of mixtures, employing the values of pure copper $( 0 . 3 \dot { 8 } \mathrm { J } \mathrm { g } ^ { - 1 } \mathrm { K } ^ { - 1 } )$ and diamondnd diamond ( $\mathsf { \Gamma } ( 0 . 5 \mathrm { J } \mathrm { g } ^ { - 1 } \mathrm { K } ^ { - 1 } )$ ). + +# 3. Results and Discussion3. Results and Discussion + +Figure 1 shows the microstructure of copper–diamond composites with differentFigure 1 shows the microstructure of copper–diamond composites with different concentrations of tungsten obtained by HP. It can be seen that pores are present at theconcentrations of tungsten obtained by HP. It can be seen that pores are present at the copper–diamond interface. The thermal conductivity of these composite is lower than thatcopper–diamond interface. The thermal conductivity of these composite is lower than that of pure copper (Table 1) but higher than that of copper–diamond composites obtainedof pure copper (Table 1) but higher than that of copper–diamond composites obtained using unmodified copper (150 Wusing unmodified copper (150 W $\mathbf { m } ^ { - 1 } \mathbf { K } ^ { - 1 } )$ [28].[28]. + +![](images/8c9784263cc2d5491ea22ff47c2bd0f89eb18e4517c18447887f93a423f5017d.jpg) +(a) + +![](images/3463ac2e7889d57f24670890b5212bf5c2b49771d95851df115e210137b10678.jpg) +(b) + +![](images/59f4217309cbee58a8074c4ee0e8efc397d81ed44493a301d03fc16a7439655f.jpg) +(c) + +![](images/0f43bca6af4072cefb7d0fc93e455cd19214a81d013547d580905d7f3474bdbc.jpg) +(d) +Figure 1. Microstructure of the copper–diamond HP-ed composites with a W-modified copper ma-Figure 1. Microstructure of the copper–diamond HP-ed composites with a W-modified copper matrix: trix: 0.15 vol.% (a), 0.35 vol.% (b), 0.7 vol.% (c) и 2 vol.% (d). Black areas are diamond, gray areas0.15 vol.% (a), 0.35 vol.% (b), 0.7 vol.% (c) и 2 vol.% (d). Black areas are diamond, gray areas are are copper and white areas are tungcopper and white areas are tungsten. + +Table 1. Sintering conditions, relative density and thermal conductivity of the copper–diamond composites. + +
Sample NumberSintering MethodAdditiveConcentrations of the Additive, Cu and Diamond, Vol.%Sintering Time, minRelative Density, %Thermal Conductivity, W m-1K-1
AdditiveCuDiamond
1HPW0.1549.85501590 ± 1188 ± 9
2HPW0.3549.65501591 ± 1192 ± 10
3HPW0.749.3501588 ± 1203 ± 10
4HPW248501590 ± 1172 ± 9
5SPSW0.749.350380 ± 1170 ± 9
6SPSW0.749.3501086 ± 1190 ± 10
7SPSMo0.749.3501091 ± 1194 ± 10
8SPSCr0.749.3501089 ± 1385 ± 19
9SPSTi0.749.3501086 ± 1420 ± 21
10HPTi0.749.3501589 ± 1178 ± 9
11SPSTi0.3549.65501082 ± 1362 ± 18
12SPSTi248501082 ± 167 ± 3
13HPTi0.3549.65501590 ± 1265 ± 13
14HPTi248501577 ± 163 ± 3
+ +The XRD patterns of the sintered composites are shown in Figure 2. No carbide phases were detected in the samples, only peaks corresponding to copper, diamond and tungsten phases are present. + +![](images/b8c96ea704af949c44e2671b5796bdfe8624242737de6e2bbd29685f1d49616d.jpg) +igure 2. XRD patterns of the composites with a W-modified copper matrix.Figure 2. XRD patterns of the composites with a W-modified copper matrix. + +1. Sintering conditions, relative density and thermal conductivity of the copper–diamond In the W-modified matrix series, composites obtained with the addition of 0.7 vol.% omposites.of W demonstrate the highest thermal conductivity (Table 1). The thermal conductivity of the composites first increases, reaching a maximum at 0.7 vol.%, and then decreases. Concentrations of the Additive, Cu and Sintering Relative Thermal Con-This effect can be explained by the fact that, at first, the addition of tungsten to the matrix + +contributes to increasing the wettability of diamonds by the modified copper matrix. As the thermal conductivity of tungsten is much lower than that of copper, a further increase in thethe tungsten concentration leads to a decrease in the thermal conductivity of the compotungsten concentration leads to a decrease in the thermal conductivity of the composites.sites. + +Since the composite with 0.7 vol.% W has the highest thermal conductivity in theSince the composite with 0.7 vol.% W has the highest thermal conductivity in the W-W-modified matrix series, additional experiments were carried out to obtain compositesmodified matrix series, additional experiments were carried out to obtain composites with with the same composition using another sintering method, SPS. When the holding timethe same composition using another sintering method, SPS. When the holding time durduring SPS was only 3 min, the thermal conductivity of the composite was lower than thating SPS was only 3 min, the thermal conductivity of the composite was lower than that of of the composite obtained by HP (Table 1). Therefore, the holding time was increased tothe composite obtained by HP (Table 1). Therefore, the holding time was increased to 10 10 min, which resulted in an increase in the thermal conductivity of the composite, somin, which resulted in an increase in the thermal conductivity of the composite, so that it that it became close to that of the composite obtained by HP with 15 min sintering timebecame close to that of the composite obtained by HP with 15 min sintering time (203 and (203 and 190190 W m−1 K− $\mathbf { W _ { m } { } ^ { - 1 } K ^ { - 1 } }$ , respectively).ly). + +In order to investigate the influence of the nature of the metal additive on the thermal conductivity of the composites, experiments were carried out using 0.7 vol.% of Mo, Crconductivity of the composites, experiments were carried out using 0.7 vol.% of Mo, Cr and Ti additives. The XRD patterns of the sintered samples are shown in Figure 3. For samples with the W and Mo additives, there are peaks corresponding to metals, but for the samples with the Cr and Ti additives, only copper and diamond peaks are present in the XRD patterns. As Cr and Ti possess lower atomic numbers, their concentration in the composites may be too small for the detection by the XRD phase analysis. + +![](images/15c5066853734b7538c4c835435db333df876d63125b58bbe26f32b5e379054c.jpg) +Figure 3. XRD patterns of the composites with 0.7 vol.% Me $( \mathrm { M e } = \mathrm { W } , \mathrm { M o } , \mathrm { C r } , \mathrm { T i } )$ -modified copper matrix. + +As seen in Table 1, the composite with 0.7 vol.% Ti has the highest thermal conductivity $( 4 2 0 \mathrm { W } \mathrm { m } ^ { - 1 } \mathrm { K } ^ { - 1 } )$ . The detailed microstructure of this composite is presented in Figure 4. It can be seen that the copper-based matrix wets the diamond surface well. As Ti has a higher solubility in Cu (~5 at.% at $9 2 0 ^ { \circ } \mathrm { C }$ and ~3 at.% at $1 0 0 0 ^ { \circ } \mathrm { C } )$ in comparison with W, Mo and $\mathrm { C r } ,$ solid solutions form during sintering of the mixtures, which improves the wettability of diamond by the matrix (the titanium concentration in the sample is less than 5 at.%). + +![](images/dee560c2c28c4ae5d9b9fee52c9a026df8f5866cd448ea0c6818a43d017e4666.jpg) +Figure 4. Microstructure of the copper–diamond composite obtained by SPS with titanium Ti addi-Figure 4. Microstructure of the copper–diamond composite obtained by SPS with titanium Ti tive 0.7 vol.%: lower magnification ×1200 (a) and higher magnification ×2500 (b). Black areas areadditive 0.7 vol.%: lower magnification ×1200 (a) and higher magnification ×2500 (b). Black areas diamond, gray areas are copper.are diamond, gray areas are copper. + +We believe that, during sintering, a metal soluble in copper (Ti) can diffuse throughWe believe that, during sintering, a metal soluble in copper (Ti) can diffuse through the copper matrix to the interface with diamond to form carbides, whereas in the case ofthe copper matrix to the interface with diamond to form carbides, whereas in the case of copper-insoluble metalcopper-insoluble metals $( \mathsf { W } , \mathsf { M o } , { \mathsf { C r } } ) ,$ r), carbides can be formed only if a direct contact of, carbides can be formed only if a direct contact of these these metal particles diamond is established. The carbides are formed according to reametal particles diamond is established. The carbides are formed according to reactions: + +$$ +\mathrm {W} + \mathrm {C} _ {\mathrm {D}} \rightarrow \mathrm {W C} \tag {1} +$$ + +$$ +2 \mathrm {W} + \mathrm {C} _ {\mathrm {D}} \rightarrow \mathrm {W} _ {2} \mathrm {C} \tag {2} +$$ + +$$ +\mathrm {M o} + \mathrm {C} _ {\mathrm {D}} \rightarrow \mathrm {M o C} \tag {3} +$$ + +$$ +2 \mathrm {M o} + \mathrm {C} _ {\mathrm {D}} \rightarrow \mathrm {M o} _ {2} \mathrm {C} \tag {4} +$$ + +$$ +3 \mathrm {C r} + 2 \mathrm {C} _ {\mathrm {D}} \rightarrow \mathrm {C r} _ {3} \mathrm {C} _ {2} \tag {5} +$$ + +$$ +\mathrm {T i} + \mathrm {C} _ {\mathrm {D}} \rightarrow \mathrm {T i C} \tag {6} +$$ + +° = –21 kJ, Δr2G° = −11 kJ, Δr3G° = −19 kJ, Δr4G° = The Gibbs free energy of these reactions at 920 $^ \circ C$ kJ, Δr5G° = −20 kJ, Δr6G° = −87 kJ. was calculated using FACT on-It can be seen that the formation of all carbline calculator (http://www.crct.polymtl. $\mathrm { c a } / \mathrm { F A C T } / )$ modynamically possible, and TiC is and are given below (per 1 mol of the mosatoms): $\Delta _ { \mathrm { r } 1 } G ^ { \circ } = - \dot { 2 } 1 \mathrm { k J } , \Delta _ { \mathrm { r } 2 } G ^ { \circ } = - 1 1 \dot { \mathrm { k J } } , \Delta _ { \mathrm { r } 3 } G ^ { \circ } = - 1 9 \mathrm { k J } , \Delta _ { \mathrm { r } 4 } G ^ { \circ } = - 2 2 \mathrm { k J } , \Delta _ { \mathrm { r } 5 } \dot { \mathrm { G } } ^ { \circ } = - 2 0 \mathrm { k J } ,$ $\Delta _ { \mathrm { r 6 } } \mathrm { G } ^ { \circ } = - 8 7$ not only the solubility of the carbide-forming metal in copper, but also the kJ. It can be seen that the formation of all carbides is thermodynamically thermodynamic characteristics of the carbidization reactions can play an important rolepossible, and TiC is the most stable compound among the studied carbides. Its formation in the structural evolution ofis favored during sintering. ${ \mathrm { S o } } ,$ mixture during sintering., not only the solubility of the carbide-forming metal in For comparison, an additional experiment was carried out to sinter a composite withcopper, but also the thermodynamic characteristics of the carbidization reactions can play the same composition by HP. However, the thermal conductivity of the obtaan important role in the structural evolution of the mixture during sintering. + +(178 W m−1 K−1) was significantly lower than that of the composite obtained by SPSFor comparison, an additional experiment was carried out to sinter a composite with (420 W m−1 K−1). It is important to note that, although the relative densities of these com-the same composition by HP. However, the thermal conductivity of the obtained composite $( 1 7 8 ~ \mathrm { W ~ m ^ { - 1 } ~ \mathrm { \dot { K } ^ { - 1 } ) } }$ was significantly lower than that of the composite obtained by SPS $( 4 2 0 \mathrm { ~ W ~ m } ^ { - 1 } \mathrm { ~ K } ^ { - 1 } )$ . It is important to note that, although the relative densities of these composites are close, the thermal conductivity of the SPS-ed composite is much higher. The microstructure of these composites is shown in Figure 5. It appears possible that + +the SPS conditions of consolidation can promote a homogeneous distribution of Ti in the copper matrix, which can facilitate uniform wetting of diamonds by the copper matrix. Formatrix, which can facilitate uniform wetting of diamonds by the copper matrix. For examexample, in the study of the formation of Mo-containing coatings on diamond [29], it wasple, in the study of the formation of Mo-containing coatings on diamond [29], it was found found that, under the SPS conditions, a more homogeneous coating was deposited on thethat, under the SPS conditions, a more homogeneous coating was deposited on the surface surface of diamond particles than under the HP conditions. This can be influenced both byof diamond particles than under the HP conditions. This can be influenced both by the the passage of electric current and by the atmosphere (vacuum in the SPS and argon in thepassage of electric current and by the atmosphere (vacuum in the SPS and argon in the HP).HP).). + +![](images/10adeb45d34a42c9c135028f72a998d5ed7c56eb360eb11caa8cc3c2fe2b872c.jpg) + +![](images/4aaa5cdad1843e0f66751879e3f6a04d20d1201bb6997da6917f789d0c49d972.jpg) +(b) +Figure 5. Microstructure of copper–diamond composite obtained with titanium Ti additive 0.7 vol.% by SPS (a) and HP (b). Black areas are diamond, gray areas are copper. + +When the concentration of titanium was 0.35 vol.%, the thermal conductivity was 265 When the concentration of titanium was 0.35 vol.%, the thermal conductivity was−1 for the HP composite and 362 W m−1 K−1 for the SPS composite. When the Ti W m−1 K−1 for the HP composite and 362 W m−1 K−1 for the SPS composite. When the Ti 265 W m−1 K−1 for the HP composite and 362 W m−1 K−1 for the SPS composite. When thecentration was increased up to 2 vol.%, the thermal conductivity of the composites concentration was increased up to 2 vol.%, the thermal conductivity of the composites Ti concentration was increased up to 2 vol.%, the thermal conductivity of the compositesrply decreased (Figure 6). The SEM data show that, at this concentration of titanium, sharply decreased (Figure 6). The SEM data show that, at this concentration of titanium, sharply decreased (Figure 6). The SEM data show that, at this concentration of titanium,oating is formed on the surface of the diamond particles, which probably leads to a a coating is formed on the surface of the diamond particles, which probably leads to a a coating is formed on the surface of the diamond particles, which probably leads to arease in the thermal conductivity of these composites (Figure 7). decrease in the thermal conductivity of these composites (Figure 7).decrease in the thermal conductivity of these composites (Figure 7). + +![](images/2e8555bd651b8035ff95996fac52cdf7165737223cf2a08558c3e107f01784df.jpg) +Figure 6. Dependence of thermal conductivity of the composites on the titanium concentration. + +![](images/2eb398c05f9531a18cd7c4a59d85fa7bb561968382442ecb343ac3b7f07e9845.jpg) +(a) + +![](images/7e4b7de0f05212744c6cf480247bb59ba26792bbdc6981da89b2ce524b1016ca.jpg) + +![](images/29f1a5d5f62729077970feb64a9c4d83aa5dc524398e549b5adbcffc34e06329.jpg) +(c) + +![](images/43f2e5ca40c449b57bf7fd5fa2f5acd2bce297912815b96b443ff106d2675f6c.jpg) +(d) + +![](images/17b5d7ce81345d8e72ffb70d3d7a8348d1b45cf28fa7d71fc3fbb01cdb89b313.jpg) +(e) + +![](images/5bc32b9e62f22803d506f818afa4af46be5c7d48134c55d6bed021fb7c960843.jpg) +(f) +Figure 7. Microstructure of the composites with 2 vol.% Ti additive obtained by HP (a,c,e) and SPSFigure 7. Microstructure of the composites with 2 vol.% Ti additive obtained by HP (a,c,e) and SPS $^ { ( \mathbf { b } , \mathbf { d } , \mathbf { f } ) }$ ), different magnifications. Dark areas are diamond, and light areas are coppe, different magnifications. Dark areas are diamond, and light areas are copper. + +The XRD patterns of composites with the Ti titanium-modified copper matrix areThe XRD patterns of composites with the Ti titanium-modified copper matrix are shown in Figure 8. No peaks other than those corresponding to copper and diamond wereshown in Figure 8. No peaks other than those corresponding to copper and diamond were + +detected, even at a titanium concentration of 2 vol.%. It can be seen that the copper peaks shift towards smaller 2θ angles as the concentration of titanium in the samples increases.shift towards smaller 2θ angles as the concentration of titanium in the samples increases. This shift indicates an increase in the lattice parameter of copper due to the incorporationThis shift indicates an increase in the lattice parameter of copper due to the incorporation of titanium into its crystal structure (Table 2).of titanium into its crystal structure (Table 2). + +![](images/ae3535b5196b6cf36c41dd62fc97eed4bc66cdfe7fd24a4689b74c63dd381994.jpg) + +![](images/f259219635c1560bc1b677d0f031cca31cf0add178e2f4fbd0cee1164fda22c7.jpg) +Figure 8. XRD patterns of composites with a Ti-modified copper matrix. + +Table 2. Titanium concentration, the sintering method and the lattice parameter of Cu in the compospositeites. + +
Sample NumberConcentrations of the Ti, Vol.%Sintering MethodCu Lattice Parameter, Å
10.35HP3.6172 ± 0.0014
20.35SPS3.6169 ± 0.0018
30.7HP3.6176 ± 0.0016
40.7SPS3.6181 ± 0.0012
52HP3.623 ± 0.002
62SPS3.6221 ± 0.0016
+ +d. The sample was mechanically ground in a mortar, after which the diamond par-In addition, a more detailed analysis of the sample containing 2 vol.%Ti was conducted. icles were separated and placed on a single-crystal Si-cuvette. The obtained XRD patterns The sample was mechanically ground in a mortar, after which the diamond particles were are shown in Figure 9. In the 2Ɵ range of 10–70°, only the diamond peak is present (Figure separated and placed on a single-crystal Si-cuvette. The obtained XRD patterns are shown 9a). Due to its high intensity, a 2Ɵ range of 30–44° was analyzed separately (Figure 9b). In in Figure 9. In the 2Trange of 10–70◦, only the diamond peak is present (Figure 9a). Due his case, we can see a peak corresponding to copper and two small peaks that may corre-to its high intensity, a 2Trange of 30–44◦ was analyzed separately (Figure 9b). In this case, spond to titanium carbide. These two peaks are broadened, which indicates the small size we can see a peak corresponding to copper and two small peaks that may correspond of crystallites. Their position is slightly shifted towards larger 2Ɵ angles, suggesting a to titanium carbide. These two peaks are broadened, which indicates the small size of smaller lattice parameter as compared with stoichiometric TiC. Based on the obtained crystallites. Their position is slightly shifted towards larger 2Tangles, suggesting a smaller data, we can assume that the coating consists of titanium carbide TiC . The presence of lattice parameter as compared with stoichiometric TiC. Based on the obtained data, we copper is probably due to the fact that some copper parcan assume that the coating consists of titanium carbide $\operatorname { T i C } _ { \mathbf { x } } .$ remained adhered to the . The presence of copper is surface of the diamond particles.probably due to the fact that some copper particles remained adhered to the surface of the diamond particles. + +![](images/b98674e9e89724bc40b47ab98b230d44ce63caeb0f7ea20fceecef87d839bdf8.jpg) +(a) + +![](images/d39dc6f1f2170e8961983b33dac951f2f90b03de84758ba2cc5b5c63ca4ed769.jpg) + +Figure 9. XRD patterns of composites with a 2 vol.% Ti-modified copper matrix, 2θ ranges from 20° Figure 9. XRD patterns of composites with a 2 vol.% Ti-modified copper matrix, 2θ ranges from to 70° (a) and from 30° to 44° (b).20◦ to 70◦ (a) and from 30◦ to 44◦ (b). + +Thus, when composites were sintered by SPS, the thermal conductivity of the sam-Thus, when composites were sintered by SPS, the thermal conductivity of the samples ples first increased with increasing titanium concentration and then sharply decreased. first increased with increasing titanium concentration and then sharply decreased. When When composites were processed by HP, there is a continuous decrease in the thermal composites were processed by HP, there is a continuous decrease in the thermal conductivity conductivity with increasing titanium concentration (Figure 7a). The observed effect can-with increasing titanium concentration (Figure 7a). The observed effect cannot be related not be related to the features of dissolution of titanium in the copper matrix as, in both to the features of dissolution of titanium in the copper matrix as, in both cases, the lattice cases, the lattice parameter of copper increased with increasing titanium concentration parameter of copper increased with increasing titanium concentration (Table 2). It is (Table 2). It is possible that, during SPS, diffusion of titanium through the copper matrix possible that, during SPS, diffusion of titanium through the copper matrix to the diamond to the diamond surface occurs faster than during HP owing to a higher real temperature surface occurs faster than during HP owing to a higher real temperature of the sample of the sample (50–100° higher than the measured temperature) and local overheat-(50–100◦ higher than the measured temperature) and local overheating/melting of the ing/melting of the metals in the former [30]. When Ti-Cu alloys are sintered in contact metals in the former [30]. When Ti-Cu alloys are sintered in contact with carbon materials, with carbon materials, titanium diffuses towards the surface of the carbon particles titanium diffuses towards the surface of the carbon particles [31,32]. The peculiarities of [31,32]. The peculiarities of heating by electric current can lead to the formation of a Ti-Cu heating by electric current can lead to the formation of a Ti-Cu alloy with a more uniform alloy with a more uniform structure and a more uniform wetting of the diamond crystals. structure and a more uniform wetting of the diamond crystals. The formation of local The formation of local high-temperature regions at the inter-particle contacts, including high-temperature regions at the inter-particle contacts, including those between particles of those between particles of different metals, during SPS can cause the development of tem-different metals, during SPS can cause the development of temperature gradients at the perature gradients at the particle size scale, accelerating diffusion and alloying between particle size scale, accelerating diffusion and alloying between the metals [33]. A direct the metals [33]. A direct effect of electric current on diffusion in the Cu-Ti system cannot effect of electric current on diffusion in the Cu-Ti system cannot be ruled out and requires be ruled out and requires further investigation. It is interesting to note that, when tungsten further investigation. It is interesting to note that, when tungsten was added to modify was added to modify the matrix, the thermal conductivity of the composite was higher the matrix, the thermal conductivity of the composite was higher when HP was used for when HP was used for consolidating the powders instead of SPS. Thus, depending on the consolidating the powders instead of SPS. Thus, depending on the grade of solubility of grade of solubility of the metal in the copper matrix, different approaches may be instru-the metal in the copper matrix, different approaches may be instrumental for increasing the mental for increasing thermal conductivity. + +# 4. Conclusions + +Based on results of the present study, it can be concluded that a critical factor for achieving a high thermal conductivity of copper–diamond composites with a modified copper matrix is the solubility of the additive metal in copper. Since Ti has the highest copper matrix is the solubility of the additive metal in copper. Since Ti has the highest solubility among the studied metals (W, Mo, Cr, Ti), Ti-containing composites possess the highest thermal conductivity. Titanium diffuses through the copper matrix to the diamond highest thermal conductivity. Titanium diffuses through the copper matrix to the dia-surface, which promotes better wetting of the diamond surface by the copper matrix. The mond surface, which promotes better wetting of the diamond surface by the copper ma-composite with a titanium concentration of 0.7 vol.% obtained by the SPS has the highest trix. The composite with a titanium concentration of 0.7 vol.% obtained by the SPS has the thermal conductivity (420 W m−1 K−1) in the studied series. Another important factor highest thermal conductivity (420 W m K ) in the studied series. Another important fac-is the sintering method—even with the same (measured) sintering temperature and Ti + +concentration, the samples obtained by the SPS have higher thermal conductivity, which can be explained by the difference in the heating mechanisms and overheating of local areas of the sample due to influence of electric current passing through the sample. + +Author Contributions: Conceptualization, A.V.U. and D.V.D.; methodology, A.V.U. and D.V.D.; investigation, A.V.U., D.V.D., M.A.E., D.A.S. and S.V.S.; writing—original draft preparation, A.V.U. and D.V.D.; writing—review and editing, A.V.U. and D.V.D. All authors have read and agreed to the published version of the manuscript. + +Funding: This work was supported by the Ministry of Science and Higher Education of the Russian Federation within State Assignment to ISSCM SB RAS, project #121032500065-5. + +Data Availability Statement: Not applicable. + +Conflicts of Interest: The authors declare no conflict of interest. + +# References + +1. Miracle, D.B. Metal matrix composites—From science to technological significance. Compos. Sci. Technol. 2005, 65, 2526–2540. [CrossRef] +2. Yamamoto, Y.; Imai, T.; Tanabe, K.; Tsuno, T.; Kumazawa, Y.; Fujimori, N. The measurement of thermal properties of diamond. Diam. Relat. Mater. 1997, 6, 1057–1061. [CrossRef] +3. Dai Sh Li, J.; Lu, N. Research progress of diamond/copper composites with high thermal conductivity. Diam. Relat. Mater. 2020, 108, 107993. +4. Zhou, H.; Ran, M.; Li, Y.; Yin, Z.; Tang, Y.; Zhang, W.; Zheng, W.; Liu, J. 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b/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/16_Scalable monolayer-functionalized nanointerface for thermal conductivity enhancement in copper-diamond composite/16_Scalable monolayer-functionalized nanointerface for thermal conductivity enhancement in copper-diamond composite.md @@ -0,0 +1,284 @@ +# Scalable monolayer-functionalized nanointerface for thermal conductivity enhancement in copper/diamond composite + +![](images/ad438b48a4820634edb1f6135e35de2cfaf36930d074c1348b0aa359d3287143.jpg) + +Bin Xu a , Shih-Wei Hung b , Shiqian Hu a , Cheng Shao a , Rulei Guo a , Junho Choi a , Takashi Kodama a , Fu-Rong Chen b , Junichiro Shiomi a, * + +a Department of Mechanical Engineering, The University of Tokyo, Tokyo, Japan +b Department of Materials Science and Engineering, City University of Hong Kong, Hong Kong, China + +# a r t i c l e i n f o + +Article history: + +Received 30 October 2020 + +Received in revised form + +3 January 2021 + +Accepted 5 January 2021 + +Available online 8 January 2021 + +Keywords: + +High thermal conductivity + +Composite + +Self-assembled monolayer + +Time-domain thermoreflectance + +Molecule dynamics + +Plasma sintering + +# a b s t r a c t + +Aiming at developing high thermal conductivity copper/diamond composite, an unconventional approach applying self-assembled monolayer (SAM) prior to the high-temperature sintering of copper/ diamond composite was utilized to enhance the thermal boundary conductance (TBC) between copper and diamond. The enhancement was first systematically confirmed on a model interface system by detailed SAM morphology characterization and TBC measurements. TBC significantly depends on the SAM coverage and ordering, and the formation of high-quality SAM promoted the TBC to 73 MW/m2 -K from 27 MW/m2 -K, the value without SAM. With the help of molecular dynamics simulations, the TBC enhancement was identified to be determined by the number of SAM bridges and the overlap of vibrational density of states. The diamond particles of 210 mm in size were simultaneously functionalized by SAM with the condition giving the highest TBC in the model system and sintered together with the copper to fabricate isotropic copper/diamond composite of 50% volume fraction. The measured thermal conductivity marked 711 W/m-K at room temperature, the highest value among the ones with similar diamond-particles volume fraction and size. This work demonstrates a novel strategy to enhance the thermal conductivity of composite materials by SAM functionalization. + +© 2021 Elsevier Ltd. All rights reserved. + +# 1. Introduction + +With the advance in high-performance power electronic devices and the accompanying increase in heat generation density, there is a growing need for technology that realizes higher heat dissipation. There, the development of materials with a thermal conductivity higher than metals like copper (Cu, 400 W/m-K at room temperature) is essential. A way to realize this without losing the merit of metals, such as surface treatability and mechanical compliance that makes them connectable to the surrounding components, is to form a composite by adding higher thermal conductivity particles. Cu/diamond composite is a promising candidate in this context because of the extremely high intrinsic thermal conductivity of diamond. However, according to the effective medium theory (EMT) [1], thermal boundary conductance (TBC) [2] at the interface between Cu matrix and diamond particles plays a crucial role in + +determining the thermal conductivity of the composite, especially when the particle size is on the order of micrometers, which is realistic considering the material cost that increases with the size of the diamond particles. + +Many studies have been carried out to enhance the TBC of the pristine Cu/diamond interface, and a few have directly probed the TBC using the time-domain thermoreflectance (TDTR) method. Approaches like the precise control over the chemical component [3], pressure [4], and the introduction of carbide buffer layers (like TiC [5,6]) was found effective to enhance the TBC. Among the above approaches, the formation of carbide buffer layers between diamond and Cu has been most commonly applied to the actual fabrication of Cu/diamond composite. Li et al. obtained the composite with an enhanced thermal conductivity of 716 W/m-K by coating Ti on diamond particles (diamond size 70 mm, 65 vol%) [7]. Shen et al. fabricated the composite with an enhanced thermal conductivity of 726 W/m-K by forming MoC2 layer (diamond size 100 mm, 65 vol%) [8]. Abyzov et al. enhanced the thermal conductivity of the composite by forming CrC or WC and obtained thermal conductivity of 770 W/m-K (diamond size 180 mm, 61 vol%) [9]. + +Formation of other kinds of carbide buffer layers like B4C [10] and ZrC [11] also help to fabricate high thermal conductivity Cu/diamond composite. Note that the values of thermal conductivity introduced above were measured at room temperature. + +Although an appropriate carbide buffer layer can improve the thermal conductivity of the Cu/diamond composite, the thickness and component of these buffer layers are hard to control during the high-temperature fabrication process, where the diamond surface reacts with the buffer layer [12]. This would give rise to alloying of the buffer layer and diamond with a thickness corresponding to their inter-diffusion length. Although a thicker alloyed layer improves the binding of diamond and Cu, since the thermal conductivity of the alloyed layer is relatively low, the partially alloyed buffer layer suppresses the effective TBC [13]. + +More recently, the self-assembled monolayer (SAM) technique [14], which can form a uniform organic layer (thickness 1e2 nm) of alkyl chain molecules, has been applied for TBC enhancement. Many molecular dynamics (MD) simulations studies have verified the function of SAM in improving the TBC. Luo et al. indicated that thermal transport across gold (Au)-SAM-Au [15,16], and GaAs-SAM-GaAs [17] could be promoted by bridging the mismatch of the vibrational density of states (vDOS) of the interface by SAM. Hung et al. studied the thickness of the SAM layer on TBC determination at Cu/SAM/diamond interface, discovering that the thicknessdependent spectral transmission of low-frequency phonons predominate the TBC [18]. Experimental studies using the TDTR method also clarified the effect of SAM on TBC enhancement. Sun et al. [19] enhanced the TBC by six-fold (from 28 $\mathsf { M } \mathsf { W } / \mathsf { m } ^ { 2 } { \cdot } \mathsf { K }$ to 169 $\mathsf { M W } / \mathsf { m } ^ { 2 } { \ - \mathrm { K } } )$ through functionalizing the Au/polymer interface with proper SAM to bridge the vibrational mismatch of the pristine interface. Losego et al. [20]and O’Brien et al. [21] achieved a doubled (from 36 $\mathsf { M W } / \mathsf { m } ^ { 2 } .$ -K to 65 $\mathsf { M W } / \mathsf { m } ^ { 2 } { \cdot } \mathsf { K } ,$ Au/quartz interface) and 4.7 times (from 30 $\mathsf { M } \mathsf { W } / \mathsf { m } ^ { 2 } { - } \mathsf { K }$ to 430 $\mathsf { M W } / \mathsf { m } ^ { 2 } { \cdot } \mathsf { K } ,$ , Cu/SAM/SiO2) enhanced TBC, respectively, by changing the end group of SAM from methyl to thiol, which increases the interfacial adhesion and thus the phonon transmission. + +The progress in the SAM functionalization to enhance TBC, as well as the stability of SAM structure at high temperature [22] motivate applying well-defined SAM to the fabrication of high thermal conductivity Cu/diamond composite. However, despite the progress made in understanding the influence of end group and chain length of SAM on TBC, there is a lack of experimental studies of the effect of SAM morphology (i.e., coverage, thickness, and chain-molecules ordering of SAM), which is sensitive to the formation process especially in scalable production that is essential for industrial application. Therefore, issues concerning the SAM morphology in fabricating high thermal conductivity composite materials still remain to be uncovered. + +Here, a systematic study combining the well-controlled model experiment, MD simulation, and composite fabrication/characterization is performed, aiming at developing high thermal + +![](images/d22d99af9077a681ad8e944e38235daf56f9ab14ecba50acd89fb09da880642b.jpg) +Fig. 1. Schematics of (a) copper/SAM/diamond composite and the zoomed up morphology of the SAM-functionalized interface; (b) Chemical structure of copper/ SAM/diamond interface. (A colour version of this figure can be viewed online.) + +conductivity Cu/diamond composite with the SAM functionalization at the interface (Fig. 1 (a)). To understand the effect of SAM to promote the thermal conductivity, we adopt a model system of a Cu deposition layer on a SAM modified diamond substrate to provide detailed characterizations of SAM morphology and the resulting TBC. By performing SAM functionalization on the diamond substrate (for the model system) and particles (for the composite fabrication) concurrently in the same reaction container, morphologies of SAM and corresponding TBC are ensured to be comparable for both cases. Several SAM formation methods are adopted to form SAM with different morphologies. By verifying the correlation between the detailed SAM morphology and the resulting TBC with the assistance of MD simulations, we elucidate the heat conduction mechanism for different SAM morphologies. Consequently, the formation of high-quality SAM gives rise to considerable enhancement of TBC in the model experiment. Applying the best SAM formation condition realizing the highest TBC onto the diamond particles facilitates the fabrication of actual composite with pronounced improvement in thermal conductivity + +# 2. Method + +# 2.1. Formation of SAM + +Samples #1~#6 were prepared with parameters shown in Table S1. The formation of SAM was carried out by two steps. For the 1st step, the diamond sample was immersed into the piranha solution (30% $\mathrm { H } _ { 2 } \mathrm { O } _ { 2 } + 7 0 \% \mathrm { H } _ { 2 } \mathrm { S O } _ { 4 } )$ for 2 h following by ultrasonicate washing using deionized water, acetone solution in sequence to obtain hydroxyl-end surface. For the 2nd step, based on our separate experimental study on the effect of end group and alkyl-chain length on TBC of Cu/diamond interface, which will be presented elsewhere soon, suggesting the superiority of octadecyltrimethoxysilane (chain length ~21 $\mathring { \mathsf { A } } , \ - \mathsf { C H } _ { 3 }$ end group) in enhancing TBC, octadecyltrimethoxysilane was used for SAM formation in the present study (Fig. 1 (b)). Both gas-phase and liquidphase process was used for SAM formation. For the gas-phase method, 30 mL silane, and diamond sample (1 g diamond particles/diamond substrate) were sealed inside a nitrogen-filled container and then heated to 150 -C for 2 h. For the liquid-phase method, the diamond sample was first immersed in a 30 mL toluene solution (99.5%, Wako) with 30 mL silane and 45 mL triethylamine for 2 or 48 h in a sealed container. Afterward, the SAMfunctionalized diamond samples prepared by both methods were washed in toluene and acetone solutions in sequence using ultrasonic to eliminate the physically adsorbed silane. + +# 2.2. Morphology of SAM + +To characterize the morphology of SAM, deionized water contact angle measurement, ellipsometry measurements, and X-ray photon spectroscopy (XPS) were carried out. The thickness and contact angle of SAM was evaluated by using the substrate sample prepared in the same manner and the same reaction vessel as diamond particles. The contact angle, which is decided by the surface energy of the substrate, strongly depends on the morphology of the surface. Therefore, it is commonly used to evaluate SAM morphology. MARY-102 (Five lab) ellipsometer, which is equipped with a 633 nm laser, and a fixed incident angle of 70-, was used to measure the SAM thickness. 25 spots were measured for individual samples to obtain the thickness and confirm the uniformity of SAM layer. The SAM thickness was obtained by fitting the measured psi and delta to the theoretical curve calculated with the air/SAM/substrate (diamond or silicon) structure (Fig. S1). The refractive index and extinction coefficient of each + +layer was summarized in Table S3. Herein, the thickness measurement was also performed for SAM concurrently formed on silicon substrates to ensure the validity of the result, since the psi varies with the thickness of SAM slightly on diamond (Fig. S1 (b)) but significantly on the silicon substrate (Fig. S1 (a)). + +The XPS measurement (Ulvac, PHI-5000 Versaprobe iii) was used to evaluate the coverage of SAM on the surface of diamond particles. Here, we focused on the XPS spectrum around the carbon 1s (C-1s) and silicon 2p (Si-2p) peaks. Since C-1s peaks mostly originate from diamond particles, with Si-2p peaks originating from the alkyl chain molecules, it is possible to evaluate the SAM coverage by the intensity ratio of Si-2p to C-1s peaks $( { \mathrm { I } } _ { \mathrm { S i } } / { \mathrm { I } } _ { \mathrm { C } } ) .$ . By comparing the ISi/IC of 6 different samples, we can obtain the relative coverage of SAM. + +# 2.3. TDTR measurement + +TDTR was used to measure the TBC of the SAM modified Cu/ diamond interface. TDTR is a pump-probe method to measure heat conduction properties at the nanoscale [23]. In this measurement, pulse laser (frequency~80 MHz, pulse width~140 fs) with a wavelength of 800 nm is split into two beams, with one followed by a wavelength conversion to 400 nm for pumping and the other for probing. The pulse laser signal is modulated with specific frequencies (11.05 MHz in the present study) to avoid the noise as well as to control the thermal penetration depth. When the aluminum (Al) transducer layer on top of the sample absorbs the pump beam, there is a temperature decay on the top of the transducer that can be probed by the probe laser. By varying the interval time between the pump and probe laser using a light path delay stage for probe laser, we can obtain the continuous temperature decay profile and hence the targeted thermal transport properties. Here, we used the sample with Al/Cu/diamond structure to measure the TBC of Cu/ diamond interface $\left( \mathsf { G } _ { \mathsf { C u - D I } } \right)$ . The Al (~60 nm) and Cu (~20 nm) layers were deposited by EB vacuum evaporation, where the Al layer acts as the transducer for the thermal reflectance signal. For this structure, because the TBC of Al/Cu interface $\left( \mathsf { G } _ { \mathsf { A l - C u } } \right)$ is extremely small, making the corresponding sensitivity ignorable (Fig. S2 (a)), the Al/Cu layers can be regarded as single layer [24]. This can simplify the heat transfer model to a two-layer structure (Table S2 summarizes the parameters of the structure). For this structure, the sensitivity calculation was also carried out (Fig. S2 (b)), showing a high sensitivity for the TBC of Cu/diamond interface $\left( \mathsf { G } _ { \mathsf { C u - D I } } \right)$ , thus the high reliability of the TBC measurement. + +# 2.4. MD simulation + +Non-equilibrium molecular dynamics (NEMD) [25] simulations were carried out to study the TBC across Cu/SAM/diamond interfaces. The generation of Cu and diamond surfaces follows the procedure of our previous study [18] to form a $\phantom { - } 1 5 . 0 \times 5 . 0 \times 3 0 . 0 \mathrm { n m } ^ { 3 }$ slab for each surface. The diamond surface was fully hydroxylated initially. The hydroxyl groups were then used as the starting points to graft the alkylsilane $( \mathrm { - } \mathrm { O S i } ( \mathrm { O H } ) _ { 2 } ( \mathrm { C H } _ { 3 } ) _ { 1 7 } \mathrm { C H } _ { 3 } )$ . Seven coverages, 0, 0.72, 0.96, 1.40, 2.24, 3.12, and 3.92 molecules/nm2 , were examined to understand the effect of SAM coverage on interfacial heat transfer. + +The interactions of Cu atoms were described by using the embedded atom method (EAM) potential [26]. The interactions of carbon atoms within the diamond surface were described by using the Tersoff potential [27]. The interactions between Cu and diamond surfaces were adopted from the previous study [28]. The force field parameters of hydroxyl groups and alkyl chains on the diamond surface were obtained from Summers et al. [29], which is based on the OPLS all-atom model [30]. The parameters of the force + +field are summarized in Tables S4eS7. The geometric combination rule was applied for the parameters of van der Waals interactions between different atoms. The Coulomb interaction was treated using the particle-particle-particle mesh (PPPM) Ewald summation method [31] for slab geometries [32]. The equations of motion were integrated with a time step of 0.5 fs. All the simulations were performed using LAMMPS [33] molecular dynamics package and visualized with the PyMOL [34] software. + +The system was initially relaxed in the canonical ensemble (NVT) at 300 K using Nose-Hoover [35,36] thermostat to ensure the system temperature. After that, two 0.5 nm thick layers at the two ends in the z-direction were fixed to stabilize the system. The system size in the z-direction was adjusted gradually according to the normal stress to remove the normal stress imposed at the interface. After the systems reached equilibrium, the NEMD simulations with constant heat exchange algorithm [37] were performed in the microcanonical (NVE) ensemble. The heat transfer was established by adding specified amount of energy to the heat source region and subtracting the same amount of energy from the heat sink region. The heat source and heat sink regions were defined as a 1.0^nm^thick layer next to the fixed layers. The temperature difference at the interface, DT, can be obtained from the temperature profile of the intermediate region. The value of TBC, G, can be evaluated by $G = J / \Delta T ,$ , where J is the magnitude of heat flux, which was $0 . 6 ~ \mathrm { G W } / \mathrm { m } ^ { 2 } . \mathrm { ~ A ~ } 5$ ns production run was performed to obtain a steady temperature profile. A linear regression of the temperature profile of each surface was applied to determine the value of DT. + +# 2.5. Composite fabrication + +Cu/diamond composite with a 50% diamond volume fraction was fabricated by plasma sintering (ELENIX, ED-PAS). The particle size of Cu particles (Sigma-Aldrich) is around 45 mm; diamond particle size is around 210 mm with thermal conductivity of 1900 W/m-K (Henan Famous Industrial Diamond Co. Ltd., China). Sintering starts from a short pulse plasma treatment (150 s) in order to eliminate the oxide layer on Cu particles, followed by the primary sintering process using direct current. The sintering pressure was fixed to be 70 MPa, with the vacuum pressure of the chamber below 5 Pa. Thermal conductivity of the composite was measured by laser flash (NETZSCH, LFA447), with density measured based on Archimedes’ principle, heat capacity calculated from the theoretical value of Cu and diamond. + +# 3. Result and discussion + +Here, we briefly describe the correlation of the coverage, thickness, and chain-molecular ordering of SAM, which becomes important in the following discussion. In the case of close-packed alkyl chain molecules (high SAM coverage), as illustrated in Fig. 1 (b), owing to the strong interchain repulsive force, the alkyl chain molecules are straight and stand perpendicular to the diamond surface spontaneously, forming an ultra-thin SAM layer, whose thickness approximates to the length of a single octadecyltrimethoxysilane chain molecule. In contrast, when the alkyl chain molecules are loose-packed (low SAM coverage), the morphology of SAM becomes relatively complicated. Because of the larger interchain distance, the repulsive force interaction becomes weaker, making the chain molecules bent and tangled. Although there is no experimental study investigating the influence of such poor ordering and the resultant small thickness of the SAM layer [38] on TBC, some simulation studies [39] have pointed out possible negative consequences on the heat transfer through the SAM bridges. + +Therefore, for the composite of Cu and SAM-functionalized diamond particles, whose interface area is extensive and hard to control, it is a critical issue to elucidate the correlation between TBC and SAM morphology experimentally. To evaluate the morphology of SAM, we conducted detailed characterizations of the morphology indicators like the coverage, thickness, and ordering. To make the TBC measured on the model system comparable to that in the composite, we performed SAM functionalization for diamond substrate (in the model system used for TDTR, contact angle, and ellipsometer measurements) and diamond particles (in the actual composite used for XPS and Fourier transform infrared spectroscopy (FT-IR) measurements) at the same time. To actively control the morphology of SAM, we performed SAM formation by two commonly used gas-phase [40] and liquid-phase [14] methods under different detailed experimental conditions (sample #1~#6). The SAM coverage, which is defined as the number of alkyl chain molecules per surface area on diamond, was conducted by XPS measurement. Fig. 2 (a) shows the SAM coverage indicated by the intensity ratio of Si-2p to C-1s peaks (ISi/IC). By comparing the ISi/IC of 6 different samples (Fig. 2 (b)), we found that in the case of sample #6 prepared by the combination of piranha-solution pretreatment and liquid-phase method with long process time (48h), ISi/IC is significantly enhanced by four times compared with that of sample #2 prepared by the gas-phase method with short process time. The FT-IR was used to characterize the SAM on sample #6, where sharp peaks representing the CeH vibration mode (2850 cm1 and 2920 cm1 ) (Fig. 2 (c)) was observed, also indicating the existence of SAM on the diamond surface. + +The water contact angle and ellipsometer measurements were conducted for the representative diamond-substrate samples #1, #4, and #6 to confirm the coverage and to evaluate the ordering of SAM (Fig. 2 (d)). Compared with the small contact angle around 45- in the case of the pristine sample (sample #1), the contact angle increases to 60- in the low SAM coverage case (sample #4), and further exceeds $1 0 0 ^ { \circ }$ in the high SAM coverage case (sample #6). Since the methyl-ends functional groups of octadecyltrimethoxysilane are hydrophobic, the increasing contact angle indicates the increase of + +the coverage and the improvement of the ordering [41]. This result agrees well with the coverage measured by XPS. However, since the alkyl chain is also hydrophobic, extra physically adsorbed alkyl chain molecules may also lead to an increase of the contact angle. + +To further clarify the morphology of SAM, eliminating the possibility of the physical adsorption, the thickness of the SAM was measured using the ellipsometer (Fig. 2 (d)). Sample #6 with high SAM coverage shows a thickness of around 19 ${ \mathring { \mathsf { A } } } ,$ similar to that of the close-packed model described above, whose thickness is around 21 Å. Combined with the measured large contact angle around 100-, SAM in sample #6 was verified to be monolayer with good molecular ordering. On the other hand, the measurement of the low SAM coverage (sample #4) case gives a smaller thickness around 15 Å and, together with the smaller contact angle, reveals the relatively poor ordering of SAM. These results rule out the possibility of SAM physical adsorption and identify the difference between SAM morphologies in the samples, which enables us to discuss their effect on TBC. + +The TDTR measurement was carried out for samples #1~#6. The inset figure in Fig. 2 (e) shows the schematic of the model experimental system. Fig. 2 (e) shows the temperature decay profiles for samples with (sample #6) and without (sample #1) SAM functionalization (The temperature decay profiles for samples #2~#5 are shown in Fig. S3). There is a distinct difference in the temporal profiles where the decay is considerably faster in the SAMfunctionalized case, indicating that bridging the Cu/diamond interface with SAM can significantly enhance TBC. Moreover, by plotting the TDTR-measured TBC for sample #1~#6 as a function of relative coverage represented by I /I (Fig. 2 (f)), we found that TBC can be enhanced by increasing the SAM coverage. Besides, the trend is nonlinear, with the slope of TBC increasing as the SAM coverage increases. As the previous study [18] has demonstrated, the SAM can bridge the phonon vibration spectra of the two materials with distinctly different vibrational properties [42], like Cu and diamond, it is possible to attribute the TBC enhancement to the increasing number of SAM bridges. + +One issue here is whether to view this bridging effect as one + +![](images/b9d2c6ed9e120263cddf9cab352e5cd08a992c30d3e668c836cf7b2d740a4526.jpg) +(a) + +![](images/2148bb315509aa5fcf55c4b80c1c905e3c0a11d2568610592b0149ac5fb22a08.jpg) + +![](images/d2d38eef5a7c97352e153016a313b46282447cd37c6f8d370d1462b67cab6237.jpg) +(c) + +![](images/a379845dcf810bd913b6753f067ee0968a54ac9a2194b068a6b503f888621294.jpg) +(d) + +![](images/50043caf784981cdc1892bee92ae2775589eca5507d3117bd0c01825192c95ef.jpg) + +(f) + +![](images/4aa03e164b7992c589aec93fea8ac55eec14b233e05cc4d647b27a9be1cfbe2a.jpg) +Fig. 2. Results of the model experiments. (a) XPS spectrum of C 1s and Si 2p and (b) intensity ratio of Si and C1s peaks (ISi/IC) for samples by different SAM functionalization methods (sample #1~#6); (c) FT-IR spectrum of sample #6; (d) Contact angle and SAM thickness of samples #1, #4 and #6; (e) Temperature decay profiles for the samples with (sample #6) and without SAM (sample #1) measured by TDTR, markers (squares and circles) are experimental data, and curves correspond to the fitting results to the physical (heat conduction) model; (f) Thermal boundary conductance of samples with different SAM coverages (Dashed curve is the guide for the eye). (A colour version of this figure can be viewed online.) + +SAM molecule giving rise to one phonon transport channel as in the often discussed ‘parallel-channel model’ [16]. In this model, each SAM molecule is assumed to act as an isolated channel due to the weak van der Waals interchain interactions, ignoring the effects of interchain thermal energy transport, interchain phonon scattering, and collective interchain vibrational modes. However, this would result in a linear increase of TBC with increasing SAM coverage and contradicts with our experimental result. This gap between the parallel-channel model and the present experimental result may result from the difference in the coverage ranges. The previous theoretical studies mainly focused on the relatively high coverage cases, where molecular ordering did not explicitly change with coverage, while the present experimental work studied a broader range of coverage, including the low coverage cases. In such an extended range of coverage, the variation of molecular ordering becomes much more significant. This assumption was supported by the MD study on the morphology variation with coverage [38], which suggested a significant molecular ordering variation in the case of low coverage, whereas slight change under high coverage. Besides, according to the studies of heat transfer through onedimensional systems like single polymer nanofiber [43] and polyethylene single-molecule [39], the conformation of the chain molecules was demonstrated to critically influence the intrachain phonon transport and the vibration modes. These factors may be the cause of the experimentally observed nonlinear dependence of TBC on the coverage. + +To further explore the cause of the coverage dependence of TBC, we carried out NEMD simulations. The schematics and corresponding temperature profiles of the representative simulation system are shown in Fig. S4. The resultant values of TBC for different SAM coverages are summarized in Fig. 3(a). Compared with the interface without SAM, the enhancement in TBC, from 24.8 to $6 8 . 9 ~ \mathrm { \ M W / m ^ { 2 } { - K } } ,$ can be achieved even at the lowest SAM coverage, 0.72 molecule/nm2 . With further increase of the SAM coverage, TBC keeps increasing at an elevating rate (Fig. 3(a)). The MD simulations capture the experimentally observed nonlinear increasing trend with the coverage for SAM-functionalized cases. + +Using the above MD simulations, we evaluated the phonon transport in terms of matching of the vibrational spectra, which determines the effect of SAM bridging. For this purpose, the vDOS, computed by the Fourier transform of normalized velocity autocorrelation functions of atoms [44], were calculated for different SAM coverages. Fig. S5 shows the calculated vDOS for different components with different SAM coverages. For the case with only hydroxyl groups (0 coverage), the large vDOS mismatch between + +the diamond and Cu results in the low TBC. With the addition of SAM, whose vDOS covers that of both Cu and diamond, the effect of bridging over the vibrational mismatch can be enhanced. Herein, since the temperature jump at the Cu/SAM interface is much larger than that inside the SAM and at the diamond/SAM interface (Fig. S4 (b)), the Cu/SAM interface predominately determines the TBC. Besides, we observed that the vDOS of SAM shifts to the lowfrequency regime (0e5 THz) with the increasing SAM coverage (Fig. 3(b)). The increase of low-frequency phonon modes inside SAM enhances the vibrational matching between SAM and Cu, whose vDOS mainly distributes within 0e8 THz. This improvement of spectral overlap between Cu and SAM [45] with increasing coverage (Fig. 3(c)) causes the enhancement of TBC. Note here, as the vDOS of SAM were calculated from the normalized velocity autocorrelation function, which means that the vDOS of SAM is not a function of the number of SAM molecules, the increasing vDOS overlap between Cu and SAM indicates the enhancement of phonon transport along individual SAM molecules. This promoted phonon transport should be responsible for the increasing slope of TBC with SAM coverage. + +Herein, we give a possible explanation for the variation of vDOS in terms of the SAM morphology under different coverages since the vibration of alkyl chain molecules is sensitive to the interaction with surrounding molecules [46]. We calculated the root-meansquare displacement (RMSD) of all SAM molecules to obtain the average mobility, which was found enhanced with the increase of coverage (Fig. S6). This result matches well with the morphology variation of SAM. Under low coverage conditions, the movement of chain molecules is strongly restricted as the backbones of some alkyl chain molecules attach on the surface of diamond or Cu (insertion (a) in Fig. S6). A model proposed by F. Sun et al. [19] in estimating the vDOS can be implemented to describe the variation in such cases. In this model, considering the whole SAM layer as a harmonic oscillator connecting the Cu and diamond, the vibrational frequency of the SAM was indicated roughly proportional to Fp , where F is the force constant taking the SAM layer as a whole [19]. Because the force constant of the whole SAM layer is partially decided by the nonbonded (like the van der Waals force) interaction, the direct interaction between the backbone of SAM and the diamond or Cu, whose Young’s modulus is typically more than two orders of magnitude higher than that of SAM (~1 GPa) [47], may increase the force constant of SAM. This would lead to the upshift of the vDOS and hence cause the decrease of the population of lowfrequency modes, in agreement with that in the previous study on the water/SAM/gold system [48]. Following this, in the case of + +![](images/abb73a5977eb71de24515876d18cd75886dda7364fdf423931a7d32d8ae4516c.jpg) + +![](images/73e85799161861607cd717ac4ea4d2740c5944c837cd29b175205e1362a08ccb.jpg) + +![](images/d0cc90d1ee8cd3637758bf1abe51c972dc315a497b992102177f05244583a69f.jpg) +Fig. 3. Molecular dynamics simulation results. (a) Thermal boundary conductance (the dashed curve is the guide for the eye), (b) vibrational density of states (0e15 THz) of SAM, and (c) corresponding vibrational spectrum overlap of copper and SAM with different coverages. (A colour version of this figure can be viewed online.) + +high coverage, the small force constant induced by the interchain interaction between individual SAM molecules [47] (insertion (b) in Fig. S5) may, in turn, induce the downshift of vDOS spectra and hence high population of the low-frequency vibration modes. It should be noted that proof of the explanation for the downshift of vDOS under higher SAM coverage requires more exploration and clarification considering the complexity of interaction in the SAM system and remains open for future studies. + +Based on the model-experiment and MD simulation studies, the high SAM coverage diamond particles (sample #6) were chosen for the fabrication of high thermal conductivity Cu/diamond composite by the plasma sintering method. Different sintering temperatures and time were used to fabricate the composite samples. Fig. 4 (a) shows the dependence of thermal conductivity on the sintering temperature that can be divided into two regions. With increasing sintering temperature from $7 7 0 ^ { \circ } \mathrm { C }$ , the thermal conductivity is low and remained almost unchanged below $7 8 5 ~ ^ { \circ } \mathrm { C }$ [region (i)]. However, the thermal conductivity is discontinuously elevated at $7 8 5 ^ { \circ } \mathrm { C } ,$ above which the value is steadily high [region (ii)]. The effect of sintering time is shown in Fig. 4 (b). Similar to the trend of temperature dependence, the thermal conductivity is low for short sintering time [region (i)], and it steeply increases when sintering time is 60 min and then saturates quickly. With high-temperature and long-time sintering, thermal conductivity as high as 711 W/ m-K was achieved. + +As a comparison, the composite was fabricated with pristine diamond particles treated only by piranha solution (pristine-diamond). Unlike that of SAM-functionalized-diamond, for pristinediamond, with increasing sintering temperature, the thermal conductivity increases continuously at low temperatures $( < 8 0 0 ^ { \circ } \mathsf { C } )$ and saturates at around $8 0 0 ^ { \circ } C \left( \mathrm { F i g . } S 7 \left( \mathsf { a } \right) \right)$ . This difference in sinteringparameter dependences can be explained by the difference in the diamond surface chemical state. For pristine-diamond, whose surface is terminated by hydroxy, is highly hydrophilic. Hence, the affinity between intenerated Cu and the pristine-diamond surface is much higher than that of SAM-functionalized-diamond, which + +![](images/78be278f72131c7fe492b68ae73a7c71fdf286ff8bcfdd33fe90177c062eb9b5.jpg) + +![](images/8405bc67467f424a7f04f2d52bb4527370099afc86749455aefaec0c885fead8.jpg) + +![](images/9b047d9ab351ea11911125595fefbfec92bc8bbd88741e01564ce1e6fa1e26fe.jpg) + +![](images/a456eed3e38f3868323c0ae2ccb1123edd916fc94198bf809cee1191923d7906.jpg) + +![](images/d64b4b980cb433dd9dd680be246e9edeb4b62eea429d9352a4614390f4b67fbb.jpg) + +![](images/ee41dd262184e4ca2bb60d203171bc2a81a4702721393e6ec52387832b36f0e8.jpg) + +![](images/724b07f50a396da83f81abe42496678dffc66a2e33509382acf1cd56aa712140.jpg) + +![](images/243ad78043b74646d401a8f06b9a982c23596d80dc72dd5378ada75f3dff9664.jpg) +Fig. 4. Thermal conductivity of copper/SAM-functionalized-diamond composite with different (a) sintering temperature and (b) sintering time, SEM image of composite sintered under (c) (d) low temperature $( 7 7 0 ^ { \circ } \mathsf C ,$ , 60 min), (e) (f) short sintering time $( 8 1 0 ^ { \circ } \mathsf C ,$ 15 min), and (g) (h) high temperature, long sintering time $( 7 9 5 ^ { \circ } \mathsf C ,$ 60 min). The arrows on the SEM images are to indicate the gaps between copper and diamond. The arrows in (h) denote some locations of copper on diamond particles. (A colour version of this figure can be viewed online.) + +has a hydrophobic surface that originates from the methyl end group. Due to the difference in affinity, which is a function of temperature [49], the formation of a tightly adhered interface using SAM-functionalized-diamond requires higher temperature and longer reaction time. + +This explanation can be further confirmed by the tensile fracture surface image of the composite measured by scanning electron microscope (SEM). For samples fabricated under low temperature and short sintering time [region (i) in Fig. 4 (a)(b), respectively] using SAM-functionalized-diamond, many large gaps between diamond and Cu matrix were observed because of the weak binding (Fig. 4(c)e(f)). When the sintering temperature and time were increased to above $8 0 0 ^ { \circ } \mathsf C$ and 1 h [region (ii) in Fig. 4 (a)(b)], the binding between Cu and diamond is well promoted and results in fewer gaps generated by the tensile fracture. Meanwhile, we observed some Cu on the diamond surface originated from the transcrystalline fracture of the Cu matrix, which only happens when the interfacial bonding is more robust than the fracture stress of Cu matrix $( \operatorname { F i g } . 4 ( \mathbf { g } ) – ( \mathbf { h } ) )$ . + +Furthermore, despite the reported high thermal stability [22], since the difference in the detailed heating parameters like heating time and vacuum degree may influence the thermal stability of SAM [50]. It is necessary to confirm the existence of SAM after the high-temperature process used in this study. The XPS (Fig. S8 (a) (b)) results show no perceptible difference in the Si-2p peaks for SAM-functionalized-diamond particles before and after the hightemperature process (800 -C, 1 h), with a negligible difference observed in the ISi/IC (Fig. S8 (c)). This result indicates the survival of SAM after the high-temperature sintering process. + +To estimate the TBC in the composite, further analysis was performed using the effective medium theory (EMT) [51], which was proven to be effective in modeling the thermal conduction in alkyl chain cross-linking nanocrystal composite materials [52,53]. Many theoretical models have been developed to describe the impact of the TBC on the thermal conductivity of composites [54]. Among those, the model proposed by Nan. et al. [1] is usually used for the composite with high thermal conductivity and shows good accordance with the experiment [55]. Moreover, to account for the impact of pores on the effective thermal conductivity of the matrix, we combine this model with a two-step approach proposed by Chu. et al. [56](Details of the model and calculation are shown in Supporting Information). The TBC of diamond composite fabricated by SAM-functionalized-diamond and pristine-diamond were calculated respectively. For the SAM-functionalized-diamond case, the TBC is 82 $\mathrm { \bar { M } W / m ^ { 2 } { - } K }$ , while that of the pristine-diamond case is 30.5 $\mathsf { M W } / \mathsf { m } ^ { 2 } { \mathrm { - K } } .$ These results are slightly higher than that measured by TDTR, whose value is 27.5 $\mathsf { M } \mathsf { W } / \mathsf { m } ^ { 2 } { \cdot } \mathsf { K }$ and 72.9 MW/ $\mathrm { m } ^ { 2 } { \cdot } \mathrm { K } ,$ , respectively, but still remain to be in reasonable agreement (Fig. 5(a)), indicating that the approach to scaling the nanointerface in Cu/SAM/diamond model system to the composite using plasma sintering is effective. + +By applying SAM functionalization on Cu/diamond composite fabrication, we succeeded in enhancing the thermal conductivity to as high as 711 W/m-K. Here, we compared our results with the previous studies $( \mathrm { F i g . ~ } 5 ( \mathbf { b } ) ) ~ [ 7 - 1 3 , 5 7 - 6 5 ]$ . Since the thermal conductivity is particularly sensitive to the diamond size and volume fraction, which are critical for deciding the material cost, although the highest experimentally achieved thermal conductivity is around 900 W/m-K [11], it is only meaningful to compare the performance of the composite fabricated with similar diamond volume fraction and size within a similar range. Here, we extracted the data from studies with the same volume fraction (50 vol%) and similar diamond sizes $( 2 0 0 - 2 5 0 \mu \mathrm { m } )$ , as summarized in the inset of Fig. 5 (b). This illustrates that the thermal conductivity of the composite fabricated in this study is higher than the previous + +![](images/1e07e7eb2a1a9c1f5a8c4cb201c89320c7ba6a495c9c4af225ed348361614a6a.jpg) + +![](images/978f4770ef5f327d0d8356c8cdfb85b6e2f105b0acbc24b560532e0a4c8f6f5d.jpg) +Fig. 5. (a) Thermal boundary conductance of pristine- and SAM-functionalized copper/ diamond interface measured by time-domain thermoreflectance (TDTR) and calculated by effective medium theory (EMT) using thermal conductivity measured by laser flash, respectively. (b) Summary of the high thermal conductivity copper/diamond composite with different diameter and volume fraction of diamond in studies thus far (different marks represent diamonds with different sizes, B: ~100 mm, ,: 100e150 mm, ⋄: 150e230 mm, △:230e250 mm ◎: 400 mm ~). + +studies, revealing the advantage of SAM functionalization on composite fabrication. + +# 4. Conclusion + +We developed high thermal conductivity Cu/diamond composite using SAM functionalization by systematically performing the model experiment and MD simulation. The influence of SAM morphology on determining TBC of Cu/diamond interface was identified by detailed morphology characterization and TDTR measurement on the model experimental system. We found a nonlinear increase of TBC with increasing SAM coverage, where the slope grows for high coverage. Assisted by MD simulation, we conclude that it is because the effect of SAM to bridge the vibrational mismatch varies with the SAM morphology. The number of phonon channels and the phonon transmission via individual channels increases with increasing SAM coverage, contributing to elevate the increasing rate of TBC. This controlled functionalization of Cu/diamond nanointerface can give TBC as high as $7 3 \mathrm { M W / m } ^ { 2 } { \mathrm { - K } } .$ . By scaling this nanointerface via a well-controlled plasma sintering process, we obtained a Cu/diamond composite with high thermal conductivity of 711 W/m-K. This study not only provides a scientific perspective on the mechanism of heat conduction across SAMfunctionalized nanointerface but also reveals the advantages of the SAM functionalization for the application of high thermalconductivity composite materials. + +# Author contributions + +J. S. conceived the idea and supervised the whole project. B. X. fabricated the sample. B. X. and R. G. performed the experimental measurements. T. K. and J. C. support with the silane coupling and characterization. SeW. H and S. H. performed the simulations with support from C. S and F-R. C.. B. X., SeW. H., and J. S. wrote the manuscript. All authors have participated in the discussion. + +# CRediT authorship contribution statement + +Bin Xu: Writing - original draft, Investigation, Methodology. Shiqian Hu: Software. Cheng Shao: Software. Rulei Guo: Investigation. Junho Choi: Investigation. Takashi Kodama: Methodology. Fu-Rong Chen: Software. Junichiro Shiomi: Project administration, Writing - review & editing, Supervision. + +# Declaration of competing interest + +The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. + +# Acknowledgements + +Some of the equipments used in this study were supported by JSPS KAKENHI (19H00744) and JST CREST(JPMJCR16Q5, JPMJCR19I2). + +# Appendix A. Supplementary data + +Supplementary data to this article can be found online at https://doi.org/10.1016/j.carbon.2021.01.018. + +# References + +[1] C.W. Nan, R. Birringer, D.R. Clarke, H. Gleiter, Effective thermal conductivity of particulate composites with interfacial thermal resistance, J. Appl. Phys. 81 (1997) 6692e6699, https://doi.org/10.1063/1.365209. +[2] H.K. Lyeo, D.G. 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thermal management of SiC power devices/17_Diamond-Cu composites microchannel heat sink for effective thermal management of SiC power devices.md b/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/17_Diamond-Cu composites microchannel heat sink for effective thermal management of SiC power devices/17_Diamond-Cu composites microchannel heat sink for effective thermal management of SiC power devices.md new file mode 100644 index 0000000..ac195e8 --- /dev/null +++ b/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/17_Diamond-Cu composites microchannel heat sink for effective thermal management of SiC power devices/17_Diamond-Cu composites microchannel heat sink for effective thermal management of SiC power devices.md @@ -0,0 +1,314 @@ +Research Paper + +# Diamond/Cu composites microchannel heat sink for effective thermal management of SiC power devices + +![](images/880ae8bea689495a477a4fba653a888115d5e31eb3b8aa2ecd095d22b575fb67.jpg) + +Kangyong Li a , Rong Zhang a , Kai Yang b , Haoran Shen b , Jialiang Chen c , FanFan Wang a , Jian Huang a , Zexin Liu d , Yue Yue d , Zhiqiang Wang b , Guoqing Xin a,* + +a Wuhan National High Magnetic Field Center and School of Materials Science & Engineering and State Key Laboratory of High Density Electrical Energy Conversion, Huazhong University of Science and Technology, Wuhan 430074, China +b School of Electrical and Electronic Engineering, Huazhong University of Science and Technology, Wuhan 430074, China +c School of Physics, Huazhong University of Science and Technology, Wuhan 430074, China +d School of Chemistry and Chemical Engineering, Huazhong University of Science and Technology, Wuhan 430074, China + +# A R T I C L E I N F O + +Keywords: + +Diamond/Cu composites + +Microchannel + +Heat sinks + +Thermal management + +Power device + +# A B S T R A C T + +SiC power devices offer higher efficiency and power density than their Si counterparts but suffer from severe heat-dissipation challenges that conventional Cu or Al heat sinks cannot meet. Diamond/Cu composites provide ultrahigh thermal conductivity, yet their application is limited by interfacial thermal resistance and the difficulty of machining into precise cooling structures. Here, we simultaneously overcome these two bottlenecks by fabricating a diamond/Cu composite microchannel heat sink: a high-conductivity diamond/Cu composite baseplate (807 W/m⋅K) was achieved via W-coated diamond particles and vacuum annealing to minimize interfacial resistance, and precise Cu microchannels were integrated by active-metal brazing to realize efficient convective cooling. Infrared thermography and thermal-resistance measurements reveal a 39 % reduction in thermal resistance compared with an all-Cu counterpart, while SiC module testing shows a 12.3 ◦C lower peak junction temperature at 200 W. This work demonstrates a practical integration of interface-engineered diamond/ Cu composites with microchannel cooling, offering a scalable route toward advanced thermal management of high-power electronics. + +# 1. Introduction + +Wide-bandgap devices like silicon carbide (SiC) have attracted significant attention as they are able to work at higher breakdown voltage, lower switching losses and higher operating frequency compared with silicon (Si)-based devices [1]. The reduced SiC chip size enables higher power density for the packaged modules, however, leading to higher heat flux and elevated junction temperature [2]. In SiC power devices, long-term operation at high temperatures leads to substantial failures, mainly due to the malfunction of packaging components [3]. Examples of such failures include bond wire lift-off [4], substrate delamination [5], die-attach cracks [6] and degradation of encapsulation materials [7]. Furthermore, multi-chips in parallel connection become an inevitable approach to increase the device’s current capacity due to the limited flow of the individual chips [8]. The asymmetric layouts and mismatched chip parameters result in unbalanced currents, leading to mismatched power losses and non-uniform junction temperatures of the + +parallel chips [9]. The mismatched thermomechanical stress of the parallel chips further degrades the reliability and stability of the SiC devices [10]. Consequently, the thermal management challenges associated with the high-power density SiC devices have emerged as a critical factor restricting device reliability and lifespan [11]. + +Several cooling methods have been explored to address the overheating issues of high power density devices, including jet impingement cooling [12,13], spray cooling [14,15], microchannel cooling [16–18] and thermoelectric cooler cooling [19,20]. Among these cooling techniques, microchannel cooling shows promising performances due to its compact structure and high heat transfer coefficient [21]. By optimizing the microchannel structure, the temperature of electronics with high heat flux has been well-controlled [22]. For instance, Sarvey et al. [23] optimized the shape and arrangement of the fins in the microchannel heat sink to enhance the cooling performance for an integrated circuit. Similarly, Drummond et al. [24] applied hierarchical manifold microchannels to successfully control the temperature of a 0.2 × 0.2 mm2 + +hotspot below 130 ◦C through optimizing the aspect ratio. These literatures reports indicate the great potential of microchannel cooling in dealing with high heat flux in SiC devices [25]. In the realm of power device thermal management, Cu and Al are the commonly employed materials for heat sinks, with thermal conductivities of 398 W/m⋅K and 270 W/m⋅K, respectively. Given the soft nature of Cu and Al, microchannel heat sinks are predominantly fabricated by skiving thick Cu or Al baseplates, with the channel width capable of scaling down to approximately 200 μm. However, the through-plane thermal resistance of the thick Cu and Al baseplate becomes a bottleneck, especially as SiC chip areas shrink and device-level heat fluxes can reach 500 $\mathrm { W } / \mathrm { c m } ^ { 2 }$ [26]. As a result, the heat conduction capabilities of traditional Cu and Al heat sinks are no longer adequate to meet the heat dissipation requirements of SiC devices, motivating higher-thermal-conductivity thermal management materials. + +Diamond/Cu composites combine high thermal conductivity with a low coefficient of thermal expansion (CTE), showing great promise in addressing the overheating problems of SiC power devices [27–29]. However, the raw-diamond/pure-Cu composite owns weak interfacial bonding, and the phonon scattering at the interface is pretty serious, resulting in the actual thermal conductivity being far inferior to the theoretical value [30]. For example, Ma et al. acquired the commercial Ti-coated diamond particles (prepared by magnetron sputtering) and hot-forged into a pancake Cu/diamond composite, but the diamond/Cu composites was revealed that the thermal conductivity of the composite was only 350 W/m⋅K [31]. In other reports, Ishida et al. [32] have utilized electroplating to achieve 595 W/m⋅K and Constantin et al. [33] employing 3D printing technology to achieve 330 W/m⋅K. According to the interfacial phonon mismatch theory model (AMM) and the differential medium effective model (DEM), the acoustic properties of the interfacial layer determine the thermal conductivity of the composites by affecting the interfacial thermal conductivity [34]. Besides, Yang et al. [35] used Molecular Dynamics to study the interfacial thermal conductance of diamond copper composites. They reported that the acoustic differences between diamond and copper are the main reason for low interfacial thermal conductance. Without careful interface engineering, the effective thermal conductivity of diamond/Cu composites can be significantly limited relative to that of their constituents [36]. Thus, the engineering of interfaces becomes crucial in the development of diamond/Cu composite thermal management materials [37]. One approach is to add reactive elements(e.g., Zr [38,39], Cr [40–42], B [43,44], or Ti [45,46]) into the Cu matrix so that they diffuse to the diamond surface and form a carbide interlayer that improves interfacial bonding and interfacial thermal conductivity [47]. Exceptional thermal conductivities exceeding 900 W/m⋅K have been reported by gas pressure infiltration with Cu alloy to achieve a better interface. However, high diamond loadings (60–81 vol%)[48,49] and melted copper at high temperature are typically empolyed, which entail higher costs and complexity. Besides, it should be noted that the dissolution of the alloying elements into the Cu matrix inevitably degrades its thermal conductivity [50]. Therefore, researchers prefer to use diamond surface metallization to introduce the interface layer [51]. Coating diamond with a carbide layer or pure metals can further improve the thermal properties of composites by mitigating the negative effects of alien elements in the Cu matrix [52]. Ti [53], Cr [54], W [55–57], Mo [58], and Zr [59] layers have been coated onto the diamond surface to enhance the interfacial bonding effectively. Several coating methods have been reported, such as ion plating [60], magnetron sputter deposition [36], and the molten salt method [61]. Compared to other diamond coating techniques, direct current (DC) magnetron sputtering technology was an ambient-temperature coating approach that did not harm the diamond during the deposition process [62]. However, additional heat treatment is generally required to increase their interface bonding force, leading to more complicated steps in diamond pretreatment [63]. Besides, reported thermal conductivities of diamond/Cu composites still deviate substantially from theoretical predictions, largely because the coating + +layers are difficult to control in terms of composition, thickness, and stability [64]. Consequently, achieving a coating that is both wellbonded and thermally efficient remains a bottleneck for realizing the full potential of diamond/Cu composites. + +Furthermore, diamond’s remarkable hardness brings significant challenges to machine diamond/Cu composites into a precise-shaped structure. Conventional machining methods such as laser engraving and CNC struggle to mill diamond/Cu composites into efficient heat dissipation structures such as microchannels. Therefore, other more flexible processing techniques need to be combined to prepare diamond/Cu composite materials into fine heat dissipation structures. For example, Wu et al. [65] fabricated micro pin fins structure with diamond/Cu composites (450 W/m⋅K) via electroplated method, and Zhang et al. [66] three-dimensionally printed diamond/Cu composites via selective laser melting (336 W/m⋅K). However, these methods were unable to fully exploit the advantages of diamond/Cu composites’ high thermal conductivity. This manufacturability bottleneck remains the main obstacle preventing diamond/Cu composites from being widely deployed in high-power electronics. + +In this study, diamond/Cu composites microchannel heat sinks have been fabricated by actively brazing Cu microchannels and diamond/Cu plates, and applied for SiC devices cooling. A diamond/Cu baseplate with high thermal conductivity up to 807 W/m⋅K has been prepared by the optimization of the interface between diamond particles and Cu. Skived Cu microchannels were then integrated with the diamond/Cu baseplate using active metal brazing to obtain the final heatsink. Benefited from the high thermal conductivity, diamond/Cu microchannel heatsinks exhibited a 39 % reduction in thermal resistance compared to the traditional Cu microchannel heatsinks with the same dimensions. The diamond/Cu microchannel heatsinks have been employed as the direct cooling baseplate for the SiC power devices. Under the same coolant flow and power input, the peak junction temperature of the SiC device with diamond/Cu microchannel heatsink shows a $1 2 . 3 ~ ^ { \circ } \mathrm { C }$ reduction compared to that of the SiC device with Cu microchannel heatsink. This research thus offers a promising solution for the effective thermal management of SiC power devices. + +# 2. Experimental + +# 2.1. Preparation of the diamond/Cu composites + +High-purity synthetic single-crystal diamond particles (HWD-40 grade, 70/80 mesh; Henan Huanghe Whirlwind Co., Ltd., China) and gas-atomized Cu powder (99.99 % purity, 25–50 µm particle size distribution; Zhongye Xindun Alloy Co., Ltd., China) were employed as starting materials. Owing to the intrinsically low thermal boundary conductance (TBC) between diamond and Cu, which is arising from atomic-structure and phonon-spectrum mismatches [67], surface pretreatment of diamond particles is carried out by a two-step metallization process: first, metallic tungsten (W) was deposited on the diamond surface by magnetron sputtering, and then vacuum annealed in a tube furnace at a temperature of 1000 ◦C for 1 h. This process can induce the formation of a uniform WC transition layer on the diamond surface, thereby enhancing the interfacial bonding strength between the dia mond particles and Cu [68]. + +The diamond/Cu composites microchannel heat sink was fabricated through a three-step process, as illustrated in Fig. 1(a). First, W-coated diamond particles were uniformly blended with Cu powder. Second, the mixture was consolidated by vacuum hot-press sintering in a graphite mold at 1000 ◦C, 50 MPa, for 3 h [53], and the specific t profile for the sintering process is shown in Fig. 1(b). This parameter set was chosen based on literature for high-volume-fraction diamond/metal composites [37,69,70], where $1 0 0 0 ~ ^ { \circ } \mathrm { C }$ ensures sufficient Cu-matrix softening for densification and a 40–60 MPa window provides strong interparticle contact without damaging the mold. In our trials, 50 MPa consistently achieved > 95 % theoretical density and high thermal conductivity, + +![](images/af343479daa4c630416948ccbaf6389d997d1124a9cb739e96f8aa452bcbcda8.jpg) +(a) + +![](images/0d6911ff47575fed920b48fc80ac166afe98c51700f777b821618cba7bcc6394.jpg) +(b) + +![](images/3fb15ff50cbc6a72b8ea8d971e8b0afba99545e0a1c54847740f245e611ffd6b.jpg) +(c) +Fig. 1. Schematic illustration of the fabrication process for the diamond/Cu baseplate microchannel heatsink. (a) Three-step manufacturing protocol including mixing W-coated diamonds and Cu powder, vacuum hot pressing, and active metal brazing; (b) Temperature–pressure profile used for vacuum hot pressing; (c) Heating profile applied during active metal brazing with AgCuTi. + +whereas pressures < 40 MPa left residual porosity and > 60 MPa often fractured graphite molds, increasing experimental and production costs. A 3 h dwell allowed adequate interfacial reaction and near-full densification. Finally, the sintered diamond/Cu baseplate was joined to precision-machined Cu microchannels (200 μm width) via active-metal brazing using an AgCuTi filler. The brazing parameters (heating curve in Fig. 1(c)) were selected to balance filler wettability, minimal thermal damage to the composite, and void-free bonding—based on prior vacuum-bra2zing studies for dissimilar materials [71,72] and verified by preliminary bonding trials in this work. This method yielded a highthermal-conductivity microchannel heat sink with a robust metallurgical bond between the diamond/Cu baseplate and the Cu microchannels, ensuring efficient thermal management performance. + +# 2.2. Characterization + +The surface morphology of diamond particles and interfacial bonding characteristics of diamond/Cu composites were analyzed using scanning electron microscopy (SEM, Thermo Fisher Scientific, Scios 2). Elemental distribution of W-coated diamond was examined by energy dispersive X-ray spectroscopy (EDS, Oxford Instruments, Ultim Max 40). Composites density was measured via a high-precision solid densitometer (DK-600A, DECCA, China) based on Archimedes’ principle. The relative density (D) of the composites can be calculated by the following Eq. (1) [60]: + +$$ +D = \frac {\rho_ {c}}{\rho_ {l}} \times 100 \% \tag{1} +$$ + +where $\rho _ { c }$ represents the measured density of diamond/Cu composites, + +and $\rho _ { l }$ denotes the theoretical density derived from rule-of-mixtures calculations. + +The thermal diffusivity of diamond/Cu composites at room temperature was determined through laser flash analysis (LFA467 Hyper Flash, NETZSCH, Germany). Specimens with standardized dimensions (3 mm thickness × 12.7 mm diameter) underwent triplicate measurements under identical conditions. The thermal diffusivity coefficient was measured three times and obtained by averaging the three values. To ensure the consistency of the thermal radiation emissivity on the surface, the samples were carbon-sprayed before the test. The density $\rho _ { c }$ of the composites material was measured by a densitometer, and the specific heat capacity $C _ { p }$ of the sample was calculated based on the mass fractions of diamond and Cu in the composites. The thermal conductivity (λ) of the composites material can be calculated by the following Eq. (2) [73]: + +$$ +\lambda = \alpha \times C _ {p} \times \rho_ {c} \tag {2} +$$ + +where α represents the measured thermal diffusivity. + +# 3. Results and discussion + +# 3.1. Microstructural characterization + +Low thermal boundary conductance (TBC) exhibits between Cu and diamond, which arises from the fundamental mismatches in their atomic structures and vibrational properties and hinders efficient phononmediated heat transfer across the interface [67]. To address this interfacial incompatibility, a strategic approach involving diamond surface metallization was implemented to establish W-based interfacial + +architectures [74]. Fig. 2 presents the surface morphology of diamond particles before and after depositing W coating. The uncoated diamond particles have sharp edges and smooth surfaces, which are typical of their natural crystallographic structure (Fig. 2(a)). Fig. 2(b) shows SEM photographs of diamond particles uniformly covered with a layer of W on the surface after magnetron sputtering. Fig. 2(c) presents a highmagnification SEM image of the W-coated diamond surface. The coating appears uniform and continuous, with an estimated average thickness of approximately 122.4 nm, as measured from the crosssectional profile. The EDS elemental maps in Fig. 2(d-f) offer additional information about the chemical composition of the coated particles. Specifically, the uniform distribution of W in Fig. 2(e) confirms an even deposition of the W layer on all crystallographic planes of the diamond particles. + +To further enhance interfacial bonding, high-temperature annealing was carried out to promote carbide formation. XRD analysis (Fig. 2(g)) shows that before annealing, the coating is primarily metallic W. After annealing at 1000 ◦C for 1 h, the intensity of W peaks decreases, while reflections corresponding to WC emerge. The diffraction peaks can be indexed to standard cards: diamond (JCPDS 06–0675) and WC (JCPDS 25–1047), confirming the generation of a thin carbide interlayer [57]. XPS analysis (Fig. 2(h)) provides further evidence of this phase evolution. The deconvoluted W 4f spectrum reveals characteristic peaks of metallic W, WC, and W C. These findings confirm that part of the W coating transforms into carbides during annealing. The coexistence of W and carbide phases indicates that a well-bonded interfacial layer is + +established. Importantly, the carbide thickness remains limited (~122.4 nm), which provides strong adhesion without introducing excessive interfacial thermal resistance [75]. Overall, the combined SEM, EDS, XRD, and XPS results demonstrate that controlled W metallization and annealing lead to a uniform W/WC/W C interfacial architecture. This structure effectively balances interfacial bonding strength and phonon transport, thereby facilitating high thermal conductivity in the diamond/Cu composites. + +Fig. 3 presents the fractured characteristics of W-coated diamond/Cu composites. The selected diamond/Cu was synthesized under optimized processing parameters of 1000 ◦C for 3 h under 50 MPa uniaxial pressure with a volume ratio of 55 % diamond and 45 % Cu powder. The SEM image demonstrates homogeneous dispersion of diamond reinforcements within the Cu matrix (Fig. 3(a)). Multi-scale analysis (Fig. 3(a-b)), combining low-magnification imaging and high-resolution microscopy, reveals the critical interfacial feature of the absence of debonding cracks at diamond/Cu interfaces. This microstructural evolution confirms the effectiveness of the tungsten carbide transition layer in establishing metallurgical bonding at the diamond-Cu interface. The observed pore-free interface morphology, particularly the intimate contact region shown in Fig. 3(b), significantly reduces phonon scattering at phase boundaries − a critical prerequisite for achieving enhanced thermal transport properties in particle-reinforced metal matrix composites [76]. + +![](images/e90a3cea2ed5dfa98e9628e08beef23e3bc634dab887f6b5eea7571e9774afb9.jpg) + +![](images/cb9715b65695d77bdc9e812d6e36cfe9a31fcc3ef608d51e8cdee6b1930cbd30.jpg) + +![](images/642c890e6fcba30ea23afa3174c1e1539823cf67875de5bde9aeb81d6842176a.jpg) + +![](images/77e36858cce44800bbab8a3a27d5d6fdabd780eb6c63ffacf0f04d3eefae2b15.jpg) + +![](images/a8b741ced5439fc5eeadd11123ee7f92d85b4d4e4a6a2ea26555337ca7b73b7c.jpg) + +![](images/be7663b8341f10564bd181527ca8ddf0a81233fcf4cd35ac3147784d1efdaca5.jpg) + +![](images/72a2a603b55ed7aa3a0f116b958dbb8724ac4bb10245462cf468246371195f66.jpg) + +![](images/20c6445981a42f69911d5b71983933d5e0bc7bb5f888a9849e2dde2c495e58ed.jpg) +Fig. 2. SEM images and EDS elemental mapping, XRD and XPS analysis of diamond particles with and without W coating. (a) Surface morphology of pristine diamond particles, (b) Surface morphology of W-coated diamond particles, (c) Highly magnified image of W-coating; (d) EDS analysis of the magnified area of Wcoated diamond particle, (e, f) EDS mapping distribution of W (e) and C (f). (g) XRD patterns of W-coated diamond before and after annealing, (h) XPS spectra of W 4f. + +![](images/a531ff9fe888333002b8cf6300808a0f766decd3add2df7c0e27e72ef7d2c297.jpg) + +![](images/f332734ecae167be0ff29fbadfe3f57ce379436b4d55e8f612c0bbe9c1acbc96.jpg) +Fig. 3. Fractured morphology of the diamond/Cu composites prepared by vacuum hot-press sintering at different magnifications. (a) The cross-sectional view of the diamond/Cu plate, (b) Enlarged view of diamond particles tightly embedded in Cu without pin-holes. + +# 3.2. Thermal performance of the diamond/Cu composites + +The thermal conductivity and heat spreading performance of diamond/Cu composites were thoroughly evaluated and compared with pure Cu baseplates, as shown in Fig. 4. The thermal conductivity and relative density of diamond/Cu composites with different diamond volume fractions are given in Fig. 4(a). It can be seen that the thermal conductivity reaches a maximum value of 807 W/m⋅K when the diamond volume fraction is increased to 55 %. The thermal conductivity decreases slightly at a higher diamond filling fraction of 60 %, which + +may be caused by the incomplete sintering densification due to the increase in the contact interface between the matrix and the reinforcing phase, at which point the relative density decreases to 97 % [77]. Fig. 4 (b) gives photographs of a diamond/Cu composites baseplate and a Cu baseplate, both with nickel-plated surfaces. A heat source was placed on the lower surface of the baseplate, and an infrared thermographic camera was used to detect the temperature distribution on the upper surface of the heatsink baseplate, as shown in Fig. 4(c), and the temperature measurement curves are shown in Fig. 4(d). The heating rate of the diamond/Cu baseplate is higher than that of the Cu baseplate under + +![](images/d79a554831e9c72298c4bb4ebce853a3c18f88e9def5f45475ee7f243789bae7.jpg) +(a) + +![](images/ae2f240e3b6d640a9173b9a92653608686ce8ae3d2cd8022caf8ec311bd21a49.jpg) + +![](images/c4692291591b6b6f10797428a4b8923340ab1e41ad293b7a4fb7bc0d6927db24.jpg) +(c) + +![](images/c65d2b900e2e822890b90587e1e342b2a4f39dd02afa8f0e902025f80cbee5c8.jpg) + +![](images/67faf445d4f27257f46d09af8bce6b034a04677140ffb3c31754b49dc95f426c.jpg) +(e +Fig. 4. Thermal performance of the diamond/Cu composites. (a)Thermal conductivities and relative densities of diamond/Cu composites with different diamond volume fractions; (b) Photograph of the diamond/Cu baseplate and the Cu baseplate with electroplated nickel finishing; (c) Infrared thermal images of surface of different baseplate; (d) Average surface temperature of two types of baseplates versus time; (e) Temperature profiles along the long axis centerline of the baseplates at 240 s. + +the same heating conditions. When the heating time reaches 240 s, the steady state temperature of the diamond/Cu baseplate is $1 4 5 . 8 \ ^ { \circ } \mathrm { C } ,$ which is 16 ◦C higher than that of the Cu baseplate, indicating that the heat transfer performance of the diamond/Cu baseplate is superior to that of the Cu baseplate along the vertical direction. Fig. 4(e) shows the temperature distribution curves along the centerline of the diamond/Cu and Cu heatsink baseplates at steady state. After statistical analysis, we can get the maximum temperature difference $( \mathrm { T } _ { \mathrm { m a x } } - \mathrm { T } _ { \mathrm { m i n } } )$ along the centerline of the diamond/Cu baseplate is $4 . 5 ^ { \circ } \mathrm { C } ,$ and the standard deviation is 0.9622, which is lower than that of the pure Cu baseplate by 29.69 % and 41.85 %, respectively, indicating that the diamond/Cu baseplate has a better heat dispersion capability along the lateral direction. + +# 3.3. Microchannel heat sink thermal resistance measurements + +A diamond/Cu microchannel heatsink was assembled by joining the highly thermally conductive diamond/Cu composites with a skived Cu microchannel heatsink. To assess the quality and effectiveness of the brazing process, the diamond/Cu microchannel heatsink was inspected and analyzed in three dimensions using X-ray imaging. It can be seen that the surface of the assembled diamond/Cu microchannel heatsink is clean, smooth and well-defined, with no visible solder swell or accumulation. Moreover, no significant voids were found inside the heatsink, indicating that the diamond/Cu microchannel heatsink is tightly brazed. Fig. 5(b) illustrates the schematic diagram of the experimental setup used to measure the thermal resistance of the microchannel heatsinks. This setup was designed to ensure uniformity during thermal performance testing, with the thermocouple positioned 2 mm from the centerline of the heat source’s long axis. Fig. 5(c) compares the surface temperature rise of the diamond/Cu heatsink and the pure Cu heatsink at different heating power levels. The results demonstrate that the surface temperature rise of the diamond/Cu heatsink is significantly lower than that of the Cu heatsink at all tested power levels. Specifically, when the heating power increased to 200 W, the surface temperature rise of the Cu heatsink is approximately 11.4 ◦C, while the surface temperature rise of the diamond/Cu heatsink was only 7.1 ◦C. + +The above results comprehensively demonstrate that the diamond/ Cu heatsink has a more efficient heat dissipation capability compared to the conventional Cu heatsink. The following equation can be used to calculate the thermal resistance of the microchannel heatsink to further quantify this heat dissipation performance [78]: + +$$ +R = \frac {T _ {h} - T c}{P} \tag {3} +$$ + +where $T _ { h }$ is the maximum temperature of heatsink surface[79], $T _ { c }$ is the temperature of water temperature at the inlet[80], P is the heating power [81]. + +Fig. 5(d) presents a comparison of thermal resistance between the two heatsinks under varying heating powers. The Diamond/Cu heatsink consistently exhibits lower thermal resistance than the Cu counterpart, demonstrating the enhanced heat dissipation efficiency of the composites material. At a heating power of 200 W, the diamond/Cu heatsink achieves thermal resistance of 0.035 K/W, which is 39 % lower than the 0.057 K/W measured for the Cu heatsink. This significant reduction in thermal resistance highlights the superior thermal conductivity of the diamond/Cu composites, thereby optimizing the overall heat dissipation performance of the microchannel heatsink. + +# 3.4. Thermal performance of SiC power device with diamond/Cu heat sink + +The SiC power device was assembled using the developed diamond/ Cu composites microchannel heatsink. Nano-silver paste is employed as the solder material to integrate the DBC substrate, SiC MOSFET, and diamond/Cu microchannel heatsink, as illustrated in Fig. 6(a). Fig. 6(b) displays the X-ray transmission image of the SiC device, demonstrating robust bonding between the diamond/Cu microchannel heatsink, SiC MOSFET, and DBC substrate, with no visible voids or defects. Fig. 6(c) shows a photograph of the complete SiC power device. Finite element simulations are further performed to compare the thermal performance of the SiC power device integrated with the diamond/Cu heatsink (Fig. 6 (d)) and the standard Cu heatsink (Fig. 6(e)). Under identical cooling + +![](images/9fd019ca0767329e392a21fd79b5ac300f02ddedffc79662b397bd84825eacd0.jpg) +(a) + +![](images/dbfafabf1a1c35342c858d51073f1987e51eb17ad443d725fa91c4529c4c9397.jpg) + +![](images/7019eaf860417976577fef9eb17ab59f5e4f005e110c6739f8d31878b3e45961.jpg) + +![](images/405cb49c357e8cc437d342bde60730fd0fd9c9bb0ec420173bffea74118b9130.jpg) +(c) + +![](images/80ef3344d6281979c2916486d4f1d56f1e85870a68a722e1205e8ffdffbb696a.jpg) +(d) +Fig. 5. Thermal resistance measurements of microchannel heatsink. (a) Photograph and X-ray transmission image of diamond/Cu heatsink; (b) Schematic diagram of the thermal resistance measurement setup for the microchannel heatsink; (c) Comparison of the temperature rise on the heatsink surface; (d) Comparison of the thermal resistance. + +![](images/82d5a230c7cd04c5296e39db02fa4f9a05faf595b35a57e5197f0a1635b913a4.jpg) + +![](images/f83b6739184ddfd5c5350cb9aff085ddc56553b482af5d8f06d3ae63252da81b.jpg) + +![](images/715158d1862ba8bfa749eb0cde4926b476c26fe0ebfdb4172cafb1424bef3c70.jpg) + +![](images/febc83c4f7b81823da528b265a4e1ce498e5212b2850ac718673f37cc9d82135.jpg) + +![](images/c4994214aedd4c3ec827478b4293ad93d4ac8bbd7d12a0682ca922da8eb6a82d.jpg) + +![](images/7df80b60a76913dd51662b7f7fa10969373cde7ea8478644fbec88d12673c0a2.jpg) + +![](images/91998a9e6d79b61822f81e69fc8abae9349fa26391387baccc007ed8495694d3.jpg) + +![](images/92f0902877179d99f90048cfaec327bc6cdad8357fdb24ac5e777f76510cade6.jpg) +Fig. 6. Thermal performance evaluation of SiC power device integrated with the diamond/Cu composites heatsink. (a) Photograph of the SiC MOSFET chip sintered onto the diamond/Cu heatsink using nanosilver paste; (b) X-ray image after chip sintering; (c) Photograph of the fully assembled SiC power device package; (d) Simulated temperature distribution of the power device using the diamond/Cu heatsink; (e) Simulated temperature distribution of the power device using a conventional Cu heatsink (f) Schematic diagram of electrode layout and wire bonding configuration; (g) Infrared thermal images of the power device using different heatsinks under identical operating conditions; (h) Junction temperature of the power device as a function of time for different heatsink materials. + +conditions (chip power: 200 W, coolant flow rate: 4 L/min), the maximum junction temperature of a device equipped with the diamond/ Cu heatsink reaches $7 0 . 9 \ ^ { \circ } \mathrm { C } ,$ significantly lower than that of the Cubased SiC device (80.8 ◦C). The results confirm that the diamond/Cu microchannel heatsink significantly enhances the thermal management capability of SiC power devices, enabling their stable operation under high-power-density conditions. In order to capture the surface temperature of SiC chips, no sealant was applied to the SiC device (Fig. 6(f)). Under this configuration, the gate and drain are short-circuited and the two Cu rows are connected to the DC power supply in order to generate heat. Fig. 6(g) presents infrared thermal images of the power devices equipped with two different heatsinks. The surface temperature of the device equipped with a diamond/Cu microchannel heatsink reaches only $7 1 . 4 ^ { \circ } \mathrm { C } ,$ while the device with the Cu heatsink reaches a higher temperature of 84.8 ◦C during the initial operation. This indicates that the diamond/Cu microchannel heatsink provides superior heatdissipating capability. The junction temperature rise curves in Fig. 6 (h) further highlight the heat removal advantages of the diamond/Cu heatsink. The steady-state junction temperature of the diamond/Cu heatsink is approximately 74.1 ◦C, significantly lower than the 86.4 ◦C recorded for the device with Cu heatsink. Additionally, during the initial heating phase, the surface temperature of the SiC power device with the Cu heatsink rises more rapidly, while the device with the diamond/Cu + +heatsink experiences a more gradual increase. This difference helps to effectively reduce the thermal stress on the SiC power device during operation. + +The improved thermal performance of the SiC power device with the diamond/Cu microchannel heatsink can be attributed to the superior thermal conductivity and uniform heat distribution of the diamond/Cu composites. The lower junction temperature and more even temperature distribution enhance the reliability and efficiency of the SiC power device. These findings highlight the potential of diamond/Cu composites for high-power electronic applications, where effective thermal management is essential. + +# 4. Conclusion + +In summary, this study addresses two long-standing bottlenecks that have limited the practical use of diamond/Cu composites for thermal management: weak interfacial bonding that suppresses conductivity, and poor machinability that prevents integration into efficient cooling structures. By combining W-coated diamond particles with vacuum annealing, we obtained a diamond/Cu baseplate of 807 W/m⋅K demonstrating effective interface engineering. Through active-metal brazing of skived Cu microchannels, we further overcame the manufacturability challenge and realized a composite microchannel heat sink + +capable of device-level operation. Infrared thermography and thermalresistance testing confirmed a 39 % reduction in thermal resistance compared with an all-Cu counterpart, while SiC module integration achieved a 12.3 ◦C reduction in peak junction temperature at 200 W. These results not only validate the dual approach of interface optimization and structural integration but also establish a scalable route to exploit the intrinsic advantages of diamond/Cu composites for advanced thermal management in SiC devices and beyond. + +# CRediT authorship contribution statement + +Kangyong Li: Writing – review & editing, Writing – original draft, Methodology, Investigation, Formal analysis, Data curation, Conceptualization. Rong Zhang: Writing – review & editing. Kai Yang: Formal analysis. Haoran Shen: Investigation. Jialiang Chen: Investigation. FanFan Wang: Data curation. Jian Huang: Investigation. Zexin Liu: Data curation. Yue Yue: Investigation. Zhiqiang Wang: Supervision. Guoqing Xin: Writing – review & editing, Methodology, Funding acquisition, Conceptualization. + +# Declaration of competing interest + +The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. + +# Acknowledgments + +This work was supported by the Key Program of National Natural Science Foundation of China (No. U24B20100). The authors appreciate the Analytical and Testing Center of Huazhong University of Science and Technology (HUST) for SEM, EDS and LFA testing. + +# Appendix A. Supplementary data + +Supplementary data to this article can be found online at https://doi. org/10.1016/j.applthermaleng.2025.129116. + +# Data availability + +The data that support the findings of this study are available from the corresponding author upon reasonable request. + +# References + +[1] Y. Zhang, T. Palacios, (Ultra)Wide-bandgap vertical power FinFETs, IEEE Trans. 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via Selective Laser Melting/18_Fabrication of Titanium and Copper-Coated Diamond-Copper Composites via Selective Laser Melting.md new file mode 100644 index 0000000..4740074 --- /dev/null +++ b/金刚石散热项目/项目前期工作_MD/03_改性粉末与CuDiamond复合热沉/18_Fabrication of Titanium and Copper-Coated Diamond-Copper Composites via Selective Laser Melting/18_Fabrication of Titanium and Copper-Coated Diamond-Copper Composites via Selective Laser Melting.md @@ -0,0 +1,291 @@ +Article + +# Fabrication of Titanium and Copper-Coated Diamond/Copper Composites via Selective Laser Melting + +Lu Zhang 1 , Yan Li 1,2 , Simeng Li 1, Ping Gong 1, Qiaoyu Chen 1, Haoze Geng 1, Minxi Sun 1, Qinglei Sun 1,* and Liang Hao 1,2,* + +Gemmological Institute, China University of Geosciences, Wuhan 430074, China; luzhang18468235771@163.com (L.Z.); yanli@cug.edu.cn (Y.L.); simengli@cug.edu.cn (S.L.); 2201710163@cug.edu.cn (P.G.); jojo_chen128@163.com (Q.C.); 13194594977@sina.cn (H.G.); smx8632@163.com (M.S.) +2 Hubei Gem & Jewelry Engineering Technology Research Center, Wuhan 430074, China +* Correspondence: sunqinglei@cug.edu.cn (Q.S.); haoliang@cug.edu.cn (L.H.) + +![](images/7a1264dab1ca0ca6bb316c4d5693fd4e9698a5f25d152768541d9213ad1a43ea.jpg) + +# check for updates + +Citation: Zhang, L.; Li, Y.; Li, S.; Gong, P.; Chen, Q.; Geng, H.; Sun, M.; Sun, Q.; Hao, L. Fabrication of Titanium and Copper-Coated Diamond/Copper Composites via Selective Laser Melting. Micromachines 2022, 13, 724. https:// doi.org/10.3390/mi13050724 + +Academic Editor: Yee Cheong Lam + +Received: 8 April 2022 + +Accepted: 28 April 2022 + +Published: 30 April 2022 + +Publisher’s Note: MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations. + +![](images/f83f3d501b95369f6a45a6b8f4b76fca877cf85f3905f4e4f46fd99b19211923.jpg) + +Copyright: © 2022 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/). + +Abstract: The poor wettability and weak interfacial bonding of diamond/copper composites are due to the incompatibility between diamond and copper which are inorganic nonmetallic and metallic material, respectively, which limit their further application in next-generation heat management materials. Coating copper and titanium on the diamond particle surface could effectively modify and improve the wettability of the diamond/copper interface via electroless plating and evaporation methods, respectively. Here, these dense and complex composites were successfully three-dimensionally printed via selective laser melting. A high thermal conductivity (TC, 336 W/mK) was produced by 3D printing 1 vol.% copper-coated diamond/copper mixed powders at an energy density of 300 J/mm3 (laser power = 180 W and scanning rate = 200 mm/s). 1 and 3 vol.% copper-coated diamond/copper composites had lower coefficients of thermal expansions and higher TCs. They also had stronger bending strengths than the corresponding titanium-coated diamond/copper composites. The interface between copper matrix and diamond reinforcement was well bonded, and there was no cracking in the 1 vol.% copper-coated diamond/copper composite sample. The optimization of the printing parameters and strategy herein is beneficial to develop new approaches for the further construction of a wider range of micro-sized diamond particles reinforced metal matrix composites. + +Keywords: selective laser melting; thermal management materials; titanium-coated diamond/copper composites; copper-coated diamond/copper composites + +# 1. Introduction + +Metal matrix composites (MMCs) are one of the most advanced potential materials in thermal management applications. The addition of nano/micro-sized fillers into metal matrix can reinforce the mechanical, thermal, and electrical performance of MMCs. To date, only relatively simple MMC designs with limited functionality have been produced by various solid (e.g., powder metallurgy) and liquid (e.g., stirring and squeeze casting) methods [1]. + +Diamond/copper composites are the excellent examples of MMCs. Diamond is a nextgeneration thermal management material, and the density of these composites is below that of pure copper. Many researchers employed diamond particles to tailor copper’s coefficient of thermal expansion (CTE) and enhance the thermal conductivity (TC) [2]. Although these diamond-particle-reinforced MMCs have great prospects in heat dissipation, it’s not easy to fabricate complex structure due to the complexity of processing diamond-based materials. With high energy density, small laser spot, and high cooling rate, selective laser melting (SLM) is a unique additive manufacturing (AM) technology that is used to construct complex designs [3], which has drawn the attention from industry and academia [4]. The method adopts a high-energy laser beam for melting metal powders layer by layer in + +3D computer aided design (CAD) data [5,6], after which the micro-melted metal pool cools to form precise metal parts [7–11]. Thus, combining SLM’s ability to create complex designs with an MMC’s excellent performance could amplify the capabilities and reshape the application [12,13]. + +However, diamond and copper have extremely poor wettability, causing weak interfacial bonding and large interfacial thermal resistance [14,15]. There are two general solutions applied for this issue [16]. The first is matrix alloying, which leads the alloying element, for example, titanium [15], boron [17,18], zirconium [19,20], or chromium [21], into the copper matrix. However, it is difficult to precisely control the additive amount, thus resulting in excess elements residual in matrix, thereby decreasing the TC of MMC. The second is surface modification, which uses strong carbide-forming elements, like titanium [22], chromium [23,24], tungsten [25], or molybdenum [26], onto the diamond particle surface. The abovementioned four metal layers can react with diamond and partly dissolve in a copper matrix. Since the formation free energy of titanium carbide (TiC) is sufficiently low, it can easily form carbide in a short time and chemically bond with diamond. The interface layer formed serves as a bridge connecting copper and diamond, thereby increasing the composite’s TC. Sun et al. [27] reported that the copper/diamond composites with titanium-coated diamond particle composites were synthesized by mechanical alloying, after annealing at $8 0 0 ^ { \circ } \mathrm { C } ,$ , the TC of a 40 vol.% diamond/copper composite reached 409 W/mK. The titanium-coating combination to the surface metallization of diamond particles helps to improve the copper/diamond interface and obtains a high TC within the composite material. Zhang et al. [28] reported that copper coated diamond powder in a variable mass ratio of diamond and copper was fabricated through an electroless plating process, before the coated powder was sintered through spark plasma sintering, the TC of a 70 vol.% diamond/copper composite reached 404 W/mK. In conclusion, they all have these disadvantages of the single sample shape and the relatively terrible interface bonding. + +This research investigates the comparison for the thermal and mechanical performance between copper-coated diamond/copper combined materials manufactured by SLM and those with titanium-coated diamond/copper combined materials. In the current research, dense complex diamond/copper combined materials with titanium and copper-coated diamond particles were successfully 3D printed by implementing different forming processes and 3D printing strategies. A series of experiments were carried out via SLM single-line scanning to determine the appropriate processing parameters, and a cubic sample was prepared to characterize the composites’ morphology and microstructure. The optimization of 3D printing parameters and strategy herein can be useful to develop new approaches for the further construction of a wider range of diamond-particle-reinforced MMCs [29]. In this work, the dense 1 vol.% copper-coated diamond/copper composite manufactured via SLM displays good interfacial bonding of the copper matrix and diamond reinforcement, and excellent thermal and mechanical performance were firstly revealed, which would serve as a promising candidate for thermal management material. + +# 2. Materials and Methods + +# 2.1. Materials + +The micron-level pure gas atomized copper powders were purchased from China Metallurgical Research Institute (purity of 99.99%), the particles of which were primarily spherical, allowing them to improve the fluidity and bulk density in comparison with polyhedral particles [30]. The average powder particle size of $1 8 . 8 5 6 \pm 1 5$ µm was below that of the laser spot (30 µm). The diamond particles were purchased from Henan Yuxing Sino-crystal Micro-diamond Co., Ltd., with an average size of approximately 25 µm. + +In this experiment, copper and titanium were directly deposited on the diamond particle surface via electroless plating (Figure 1a) and evaporation methods (Figure 1b), respectively. Table 1 lists the mix ratios of copper powder and coated diamond particles, combined in a ball mill at 100 rpm for 3 h, before which were dried for 3 h at 60 $^ \circ C$ and then sifted through a 400 mesh. + +![](images/ea5ebd7bf51386ccdbf464849bc1f0ee5b8f2b7c7fcccaddf496b1f1f2b33fb2.jpg) + +![](images/c5f36de4a13f42081206a7bdbc256ca4232196c67c7b00e8e7230d40861dd342.jpg) + +![](images/b04e8a2f516c25009081bc21ad5be25d9f9591751e7e548d2237a92fff4edc56.jpg) +atic diagram of depositing copper and titanium on the diamond particle surface via Figure 1. Schematic diagram of depositing copper and titanium on the diamond particle surface ating and (b) evaporation process, respectively; (c) The preparation process for rec-via (a) electroless plating and (b) evaporation process, respectively; (c) The preparation process for r and cubic samples. rectangular contour and cubic samples. + +Table 1. The coated diamond/copper composite compositions. + +
Diamond vol.%Diamond wt.%Total Mass (g)Coated Diamond Quality (g)Copper Quality (g)
10.40500.2049.80
31.20500.6049.40
52.03501.0148.99
+ +# 2.2. Preparation of the Titanium and Copper-Coated Diamond/Copper Components through SLM + +In order to obtain the high-quality titanium and copper-coated diamond/copper samples, the SISMA MYSINT100 system with a neodymium-doped yttrium aluminum garnet fiber laser (wavelength: 1060 nm; maximal output laser power, 180 W; laser spot size, 30 µm) was used to investigate the detailed parameters from singlet to cubic formation in high-purity N2 atmosphere (residual oxygen content $< 0 . 5$ vol.%). The titanium and copper-coated diamond/copper powders were protected from oxidation. Rectangular contour $( 1 \times 3 \mathrm { m m } ^ { 2 } )$ and cubic $( 5 ^ { ^ { - } } \times 5 \times 5 \mathrm { m m } ^ { 3 } )$ samples were used for the monorail and block experiments, separately (Figure 1c). The chessboard laser scanning strategy was used for the cubic samples, each layer was separated into four squares, and the scanning direction in each square was perpendicular to the adjacent square. The no-hatch laser scanning strategy was used for the rectangular contour samples. A single-line scan was performed via SLM to determine the appropriate processing parameters, which were then used to prepare the cubic samples for the morphology and microstructure characterization. + +The hatch distance indicated the degree to which the titanium and copper-coated diamond/copper powders were repeatedly scanned by the laser. At the same laser power and scanning rate, the smaller the hatch distance, the greater the laser’s influence on the composites. A variety of volumetric laser energy density (D) was adopted to determine the finished part parameters. D is computed through Equation (1): + +$$ +D = \frac {P}{h \times t \times v} \tag {1} +$$ + +in which $P , v ,$ h and t refer to the laser power, scanning rate, hatch distance and layer thickness, respectively. + +# 2.3. Characterization Techniques + +The surface morphology and microstructure were observed via optical microscopy (OM, Leica A205, Wetzlar, Germany) and scan electron microscopy with energy dispersive spectroscopy (SEM-EDS, Hitachi-Su8010, Hitachi High-Tech, Clarksburg, MA, USA). Prior to the microscopic observation, the specimens were polished and etched (10 s within a mixed solution of 100 mL distilled water, 5 mL HCl and 5 g FeCl3). The width of a single weld pool within the printed composites was assessed via image analysis with ImageJ. Titanium-coating and copper-coating thickness on the diamond particle surface were determined through focused ion beam (FIB) and SEM. The coated diamond cross-section specimens were prepared using FIB milling. X-ray diffraction (XRD, Bruker D8 Advance, Rheinstetten, Germany) employing Cu Kα radiation at 40 KV and 40 mA within the scope of $2 \theta = 5 ^ { \circ } - 9 0 ^ { \circ }$ was used to test the phase structure with a step size of 0.02◦. The composite density was determined adopting Archimedes’ law. A Netzsch LFA 427 Transient Laser Flash machine and the calorimetric technology were used for measuring the thermal diffusivity and specific heat, respectively. The TC was obtained by multiplying the composite density, thermal diffusivity, and specific heat. A carbon spray was used to coat the top and bottom surfaces of the specimens (Ø 12.7 × 3 mm) with graphite to improve their capacity for absorbing the applied energy. The mean TC of each specimen was determined by measuring three parallel positions. Finally, samples $( 3 \times 4 \times 2 5 \mathrm { m m } ^ { 3 } )$ were used to test the bending strength and CTE. The bending strength test was performed applying an Instron 5569 Universal Testing machine, and a dilatometer (DIL 402C, Netzsch, Selb, Germany) was used to examine the CTE of the composite materials from room temperature + +to 400 ◦C. The roughness of the top surface was measured by a laser scanning confocal microscope (OLYMPUS OLS4100) with Gaussian filtering to evaluate the quality of 1 vol.% copper-coated diamond/copper combined materials surfaces, the OLS4100 can correctly identify a measuring position and easily perform roughness measurement of a target micro area, the accuracy of height measurement was less than 0.2µm error per 100µm. + +# 3. Results and Discussion + +# 3.1. Characterization and Analysis + +The vacuum evaporation technique was used to coat the diamond particle surface with titanium [31]. Figure 2a displays the titanium-rich regions on the diamond sites via EDS element mapping, showing the preference of titanium for diffusion bonding with the diamond. The EDS elemental map of the titanium-coating layer is exhibited in Figure 2b, demonstrating the enrichment of Ti. The protective deposition on the surface edge caused the Pt signals. The titanium layer thickness on the diamond particle surface ranged between 93.04 and 122.8 nm. The coating layer was closely attached to the diamond particles, and the clear stepped surfaces illustrated that the titanium and diamond were strongly chemically bonded. Diamond particles were coated employing copper through electroless plating for enhancing its wettability with molten copper, the copper-coating layer was thick enough to melt with the copper powder that only copper region on the diamond site was displayed by EDS elements mapping (Figure 2c). And the copper layer thickness on the diamond particle surface ranged between 0.99 and 1.77 µm (Figure 2d). + +The series of diffraction peaks marked in black plum blossom were indexed to the TiC (PDF#32-1383), and the peaks with rhombus were indexed to the diamond (PDF#06-0675). No impurities were characterized by the XRD (Figure 2e). As the natural non-wettability of diamond and copper made it difficult to form intermediate products and impossible to achieve perfect interface bonding [8], active element titanium was coated on the diamond particle surface for clearly generating carbide TiC at the interface. As a result, the metal matrix covered the diamond surface well, improving the poor interface bonding of the diamond and metal. Existing research confirmed these phenomena [15,32–36]. The series of diffraction peaks marked with heart symbols were indexed to copper (PDF#04-0836), and those peaks with rhombus were indexed to diamond (PDF#06-0675). No impurities were characterized by the XRD (Figure 2f). The copper plating on the diamond particle surface improved the poor interface adhesion between the diamond and copper. + +# 3.2. Formation of the SLM Titanium and Copper-Coated Diamond/Copper Composites + +# 3.2.1. SLM Manufacturing of Titanium-Coated Diamond/Copper Composites + +The rectangular contour samples were printed in the single-track experiment, and their morphology was characterized and correlated with the process parameters. Due to scanning a single-track layer-by-layer along the Z-axis to form each rectangular contour sample, the width of a single wall was the width of a single weld pool. Samples with different process parameters were observed via OM to estimate the range of process parameters in the cubic experiment. The SLM laser parameters selected for the single-track formation test of the titanium-coated diamond/copper combined materials were as follows. The laser power was improved from 140 to 160 and 180 W, and the scanning rate was improved from 200 to 300 and 400 mm/s at a fixed powder layer thickness of 0.025 mm. It was important to observe the continuity of the molten pool, as any discontinuity would produce defects, for example, porosity, delamination and surface roughness. Figure 3a,b exhibited that the superior process parameters of the rectangular contour samples with 3 and 5 vol.% titanium-coated diamond/copper composites comprised a laser power of 140 W and a scanning rate of 200 mm/s, that with 1 vol.% titanium-coated diamond/copper composites comprised a laser power of 180 W and a scanning rate of 200 mm/s. In the single-track experiment, when the layer thickness was below the median powder particle size (0.025 mm), a significant amount of powder particles larger than the layer thickness + +would rub the previous layer in the melting region as the scraper moved, leading to failureEVIEW 6 of the print formation. Therefore, a layer thickness of 0.025 mm was optimal. + +![](images/ccff64f39af885f425142c94c7366103ef7533a5967bf49a077d2ca2a4d54448.jpg) + +![](images/cd8571aa0e8bb6c80aca9d4fd3612d408164fa97130f528b32e6720904180e05.jpg) + +![](images/481f642e0df9289010f631f53c1a7ca776cf32dd54d7590bd0b7bacc5246f5a5.jpg) + +![](images/2327cd0e1b3cfc51224a7b66e375f675c31c0dfe232c2190e74e174deec43cfe.jpg) + +![](images/abd007f1a81f31986563a3ab3a8d68104e3c46fad316975230013d6396a6f7b2.jpg) + +![](images/3406a32b5f426c773f998341f2211b9470e553cb4b8ce934256ed0ed7193d3f7.jpg) + +![](images/961f894c24fd3759b667eaa5d07ca1d5979933ffbd538cb2ca379678afd7eafb.jpg) + +![](images/a6bc39b6e7243ebd5f0afb63c5e4a323c736489f5f3055344ccd10c1453567bc.jpg) + +![](images/0a00bbdf026b381a7429bea3e63ff4110862f131aa567c730ba2121011fc76ed.jpg) + +![](images/74ced158468e9428f9a9ea5cd32aa0ba406b19b4363de5160d307cd8a9744d95.jpg) + +![](images/f03fdbc98690f7437786005da28dbcfde0423eb70657f824141d4ec83c07b56d.jpg) + +![](images/765ef077ec9d44c5b5eebe5758fd163ed875ce90e514c60c5f5cdf53cf3a4dd0.jpg) +gure 2. (a) The morphologies of the titanium-coated diamond particles and the corresponFigure 2. (a) The morphologies of the titanium-coated diamond particles and the corresponding EDS DS element mappings; (b) The thickness of the titanium-coating layer on the diamond parelement mappings; (b) The thickness of the titanium-coating layer on the diamond particle surface urface and the corresponding EDS element mappings; (c) The morphologies of the copper-coand the corresponding EDS element mappings; (c) The morphologies of the copper-coated diamond iamond particles and the corresponding EDS element mappings; (d) The thickness of the copparticles and the corresponding EDS element mappings; (d) The thickness of the copper-coating layer ating layer on the diamond particle surface and the corresponding EDS element mappings;on the diamond particle surface and the corresponding EDS element mappings; The XRD pattern of RD pattern of (e) the titanium-coated and (f) the copper-coated (e) the titanium-coated and (f) the copper-coated diamond particles. + +![](images/88c51a28b31d0aefdcdfdf682d481c08d435a297bf0f472674b390bbd6ab819b.jpg) +b + +![](images/fd1c78df305e90571f5feb8dd574c76544c0ad26aea08da662e32b58b8bb4cd7.jpg) + +![](images/98f1cb3f9c0c812a115b005bb0aa98e5b18988f67eb4ba8db7816eeeda94d50b.jpg) + +![](images/ec56eab9601c0a1f132ef28606b66c86724aa602e159591a3c39aaf9c65a698b.jpg) +Excessive melting Goodmelting Unstable melting Weak sintering + +![](images/b54d4d555658147dca057f7095a58d5bfb0fd4a0e1b9804cdda6431682d872c9.jpg) +c + +![](images/f17df5202221535615841ff0e06676745be4f20d5e06816daa2389d1b0654ad7.jpg) + +![](images/0762e7ed29a39a003fd8d025c9e69278fce1494f77fcc098e03e845fcb126115.jpg) +gure 3. (a) Rectangular contour samples of the 1, 3 and 5 vol.% titanium-coated diamondFigure 3. (a) Rectangular contour samples of the 1, 3 and 5 vol.% titanium-coated diamond/copper mposites; (b) the corresponding processing window of the laser power and scanning rate;composites; (b) the corresponding processing window of the laser power and scanning rate; (c) The M manufactured morphology: top and front view of the 1, 3 and 5 vol.% titanium-coatSLM manufactured morphology: top and front view of the 1, 3 and 5 vol.% titanium-coated diand/copper composite, respectivelymond/copper composite, respectively. + +ompared with the single-track experiment involving only three process paraCompared with the single-track experiment involving only three process parameters (layer thickness, scanning rate and laser power), the cubic experiment further considered the hatch distance and the scanning strategy. The hatch distance (h) was calculated as follows: + +$$ +h = (1 - H r) \times w \tag {2} +$$ + +where Hr and w indicate the overlap rate and melt pool width, respectively. The higher the overlap rate, the higher the density and the lower the porosity. The highest densities and lowest porosities were obtained for each parameter group when the melt pool overlap rate reached 60%. The center of the next scan track overlapped the edge of the adjacent single-track melt pool and was partially remelted to achieve good metallurgical bonding between the pools and reduce the porosity. However, when the overlap was greater than 60%, the porosity increased slightly. If the laser energy density increased further upon a rise in the laser power or a decrease in the hatch distance, the porosity increased slightly when the density reached its peak value, that is, when the porosity was at its lowest. Existing research confirmed these phenomena [37–39]. + +The data of the single-track formation test of the 1, 3 and 5 vol.% titanium-coated diamond/copper combined materials could be obtained through the ImageJ software, the melt pool width was found to be approximately 250 µm, and the h was calculated as approximately 100 µm. Figure 3c shows the unpolished samples of the SLM-printed 1, 3 and 5 vol.% titanium-coated diamond/copper combined materials. The 1 and 3 vol.% composites had higher relative density than that of the 5 vol.% composite, the relative density of 1, 3 and 5 vol.% titanium-coated diamond/copper composites were 96%, 90% and 81%, respectively. A higher relative density is related to better specimen performance [40]. Hence, 1 and 3 vol.% coated diamond/copper composites featuring high relative density were chosen for the investigation below. + +Additionally, with the emergence of a metal vapor above the molten pool, a recoil pressure was induced onto the surface of pool. Meanwhile with the quick development of a thermal gradient inside the liquid, a flow of molten metal was produced from the hot region to the coldest one, referred to as the Marangoni role. Both roles brought particle ejection, known as recoil-induced ejection, as described in Figure 4 [41,42]. Besides, a denudation area was built on the sides of the molten track, where we could sweep powders under the flow of the metal vapor and the gas shield. This kind of particle ejections was known as entrainment particle (Figure 4) [43]. Metal vapor-induced entrainment caused powderspattering behavior [44]. To be intuitive, micrometric particles changed the surface tension, molten pool rheology, and metal vapor density according to declaration [45–47]. Also, it was believed that the imbalance of the molten pool, hot spatter collision during flight and entrained particles gathered on the solidified area were among the major driving forces for big spatter production. A pronounced spattering situation was found while mixing 5 vol.% coated diamond with copper by comparing with a pure copper bed condition, thus determining the negative role of the coated diamond in the molten pool. The hot particles partially fused with the solid materials after falling back into the powder bed, cooled downR REVIEW 9 of 16 after ejection [29]. + +![](images/ead683561630823d794fded357a2b872bf07ba94d26660eed8b0fe26f9d7a974.jpg) +Figure 4. Spatter ejection phenomenon in SLM of the coated diamond/copper combined materials. Figure 4. Spatter ejection phenomenon in SLM of the coated diamond/copper combined materials. + +# 3.2.2. The Formation of SLM Copper-Coated Diamond/Copper Composites + +The preparation of dense copper/diamond composite materials with a high TC required the shaping of a strong interphase layer between the diamond and copper due to their incompatibility [48,49]. For improving the wettability of metal matrix and diamond, many researchers adopted electroless plating for coating the copper layer on diamond powders [50]. The superior SLM process parameters of the rectangular contour samples with 1 vol.% copper-coated diamond/copper composites comprised a laser power of 180 W and a scanning rate of 200 mm/s (Figure 5a), that with 3 vol.% copper-coated diamond/copper composites comprised a laser power of 160 W and a scanning rate of 100 mm/s (Figure 5c,d). According to the quality of the melted rails, the processing window was separated into four different zones (Figure 5d,e), that is, a weak sintering zone (52.8%; A), an unstable melting zone (22.2%; B), a continuous track zone (2.8%; C), and an over-melting zone (22.2%; D), in reference to their different line-energy densities (LEDs, J/m). The less continuous track area suggested a formation difficulty under the constrained process parameters. It was not easy to form a scanning track due to the insufficient LED within zone A. The powder failed to form a stable width track within zone B, and a large amount of unmelted powder adhered to the surface due to insufficient energy. In zone C, the melt flow in the molten pool stabilized and had a sufficient penetration depth into the previous layer, thus obtaining a relatively smooth trajectory when the input LED was 1600 J/m. When the LED was too high (>1700 J/m) in zone D, a track featuring a width of approximately 517 µm occurred, and micro-cracks appeared due to the accumulation of excess heat related to the high power and low scanning rate caused by the high residual stress [51]. + +The printing parameters were optimized to form the dense copper/diamond composites. Scanning speeds ranging from 50 to 300 mm/s and laser power levels ranging between 130 and 180 W caused numerous printing defects (e.g., balling and pores). To obtain parts with minimal printing defects, the narrow processing window was determined, with a scanning rate of 100 mm/s and a high laser power (160 W). With a high laser power (170–180 W) and a low scanning rate (50 mm/s), the printed parts exhibited superfusion (Figure 5c,e). These parameters increased the molten pool size, thereby increasing the height and width of the powder tracks. + +Additionally, the values of roughness gradually decreased with increasing laser power (Figure 6). When the laser power reached the maximum value of 180 W, the surface roughness value Sa reached a minimum value of 5.751 µm at the scanning rate of 200 mm/s. While the scanning rate was varied between 50 and 200 mm/s, the values of roughness gradually decreased with increasing scanning rate. When the scanning rate was 200 mm/s, the surface roughness value Sa reached a minimum. While the scanning rate was varied between 200 and 300 mm/s, the values of roughness gradually increased with increasing scanning rate. Laser power and scanning rate played important roles in the roughness of the composite. The lifetime of the molten pool is the key parameter that influence the flatness and surface roughness, which will be discussed in the future [52]. + +# 3.3. Comparison of the SLM Titanium and Copper-Coated Diamond/Copper Composites + +Interfacial bonding holds the key to determining the thermal and mechanical performance of composites [53]. 1 vol.% copper-coated diamond/copper composite sample showed relatively better interface bonding of the copper matrix and diamond reinforcement (Figure 7c), no remarkable defects like flaws or cracks were seen at the interface. The pull-out of diamond particle was merely discovered in the polished surface, implying strong interfacial bonding between the copper matrix and diamond particles. However, 1 vol.% titanium-coated diamond/copper combined material sample showed the pull-out of diamond particle that could be discovered in the polished surface (Figure 7a). For the 3 vol.% titanium-coated diamond/copper combined material sample, the copper matrix and diamond reinforcement bonding were extremely poor and showed obvious cracking (Figure 7b). Resulting from poor interface bonding between diamond and copper, 3 vol.% titanium-coated diamond/copper composite showed a low TC of 57 W/mK at + +an energy density ofstress $2 8 0 \mathrm { J } / \mathrm { m m } ^ { 3 }$ (140 W, 200 mm/s). There were small cracks in the edge area between copper matrix and diamond reinforcement for the 3 vol.% copper-coatedThe printing parameters were optimized to form the dense copper/diam diamond/copper composite sample (Figure 7d). Good interfacial bonding was exhibitedsites. Scanning speeds ranging from 50 to 300 mm/s and laser power levels in the 3 vol.% copper-coated diamond/copper composite sample for comparison withtween 130 and 180 W caused numerous printing defects (e.g., balling and pore the 3 vol.% titanium-coated diamond/copper composite sample. The reason why theparts with minimal printing defects, the narrow processing window was deter interfacial bonding of titanium-coated diamond/copper composite was worse than that ofa scanning rate of 100 mm/s and a high laser power (160 W). With a high laser copper-coated diamond/copper composite was that introducing electroless copper plating180 W) and a low scanning rate (50 mm/s), the printed parts exhibited superfu process could avoid essentially the particles gathering and it could improve the relative5c,e). These parameters increased the molten pool size, thereby increasing the density, the interfacial bonding and the TC of the diamond/copper composites [28].width of the powder tracks. + +![](images/b84ac5ddcbe9dc406d9fd9f563a82cec2f04141d8d48ec54f335c5efdba976cf.jpg) + +![](images/55e7fe2af343caad39bba471dafb3d35f27720b30bb09ea9015c93d4115af9c6.jpg) + +![](images/e1d5e5fd735f2505d2a25fb6dd2ced3218446c5361ecd722727dfec1ce7c8b24.jpg) + +![](images/40e63f65d1854ffb7878306fd83f3fce66289c5f809fcbd5a63b72276ebceaac.jpg) + +![](images/d05d5ca0a91b5b0d91bbeda48e7227132b4cc3600d3f4ef27e36647496707b64.jpg) + +![](images/04e0cf819abe26bdfb132498cb7eec349b792e63ea652f10458c8b347d010075.jpg) + +![](images/f48cccf8ab8d63da9a6b234761b374596b176ad85527e39af02dab607007bd3e.jpg) + +![](images/3956c4d675150268cce5e37d0b00479368334daab9db504ece4cbb4642984d04.jpg) + +![](images/7b28f21b86d4dd2621ae4bd1ba67b3f479e8b7db2640ddc56af90c437fc8f7f8.jpg) + +![](images/d1e80aa938c68483aac03865017e7c0bcd6cf5b971d1dd6e6290cdbf0e99177f.jpg) +Figure 5. 1, 3 vol.% copper-coated diamond/copper combined materials and pure copper featuring various process parameters within the XY plane and surface morphologies: (a) 1 vol.% copper-coated diamond/copper composites; (b) Pure copper; (c,d) 3 vol.% copper-coated diamond composites and process window of laser power and scanning rate; (e) Typical track types of zones A, B, C, and D. + +![](images/42eae12f220dbce9a8375da2b3b8bae2967e15f243b414a7358bd818ee0c6146.jpg) + +![](images/6956055c5ac2e3394c443cfec89a1a9f0de68ffc693195546c16df3b924f90ba.jpg) +Figure 6. 1 vol.% copper-coated diamond/copper combined materials relationship between the sur-Figure 6. 1 vol.% copper-coated diamond/copper combined materials relationship between theR REVIEW 13 of 16 surface roughness and (a) laser power and (b) scanning rate. + +![](images/bf655769bc893999fbed6cf8cfcf4ffa885bac1f979050fdea89ff30198daae0.jpg) + +![](images/f8c8679dbe7ee1aa15648b3aca55f84059ff663f43f5f425ee8a8ed6b08422a6.jpg) + +![](images/bf549a8c5494d78c4770c428f782ba00c1873e7076952663bbd23983af2a3915.jpg) + +![](images/479a3cb1dfcf7e5d36ff797f88deb7c40c8c1eebf950045df80a40dbe7923890.jpg) + +![](images/a85fbac140b4102d877f0c9a3401233c2dd986ff2d5292e0cf89320cab33aca8.jpg) + +![](images/e25f862ec3f1f36314b693053343e353803fb0cac05f33c263578f313a8c1915.jpg) +Figure 7. SEM images of the copper matrix and diamond bonding of (a) 1 vol.% and (b) 3 vol.% Figure 7. SEM images of the copper matrix and diamond bonding of (a) 1 vol.% and (b) 3 vol.% titanium-coated diamond/copper combined materials; (c) 1 vol.% and (d) 3 vol.% copper-coated di-titanium-coated diamond/copper combined materials; (c) 1 vol.% and (d) 3 vol.% copper-coated amond/copper combined materials; (e) The bending stress and (f) CTE values of titanium and cop-diamond/copper combined materials; (e) The bending stress and (f) CTE values of titanium and copper-coated diamond/copper combined materials and pure copper. + +The bending strength of 1 and 3 vol.% copper-coated diamond/copper composites significantly exceeded that of corresponding titanium-coated diamond/copper composites. The maximum bending strength of 3 vol.% copper-coated combined materials was 108 MPa, while that of the 3 vol.% titanium-coated combined materials was only 36 MPa. The maximum bending strength of the 1 vol.% copper-coated combined materials was 150 MPa, that of the 1 vol.% titanium-coated composites was 148 MPa. And the printed composites with 1 vol.% copper-coated diamond were approximately three times stronger than the printed copper (≤58 MPa) for the bending strength. When the diamond concentration rose from 1 to 3 vol.%, the bending strength decreased as the relative density decreased (Figure 7e), the viscosity of melt increases obviously as the diamond particle content increases at a constant size, resulting in the decrease in the fluidity and deterioration in the sample surface quality, leading to the decrease in relative density [54]. These experiments indicated that a moderate TC (174 W/mK) was produced by printing 1 vol.% titanium-coated diamond/copper mixed powders at an energy density of 360 J/mm3 (180 W, 200 mm/s), a maximum TC (336 W/mK) was produced by printing 1 vol.% copper-coated diamond/copper mixed powders at an energy density of 300 J/mm3 (180 W, 200 mm/s) and a moderate TC (162 W/mK) was produced by printing 3 vol.% copper-coated diamond/copper mixed powders at an energy density of 533 J/mm3 (160 W, 100 mm/s). Additionally, the printed composites with 3 vol.% copper-coated diamond were approximately three times larger than those with 3 vol.% titanium-coated diamond for the TC. A moderate TC (183 W/mK) was produced by printing the copper powders at an energy density of 171 $\mathrm { J } / \mathrm { m m } ^ { 3 }$ (180 W, 350 mm/s) as shown in Figure 5b. This difference was likely caused by the dissimilarity of those interfacial bonding between the copper matrix and diamond reinforcement. The higher energy density required to print diamond/copper composites compared to pure copper was caused by solid coated diamond particles into the molten copper pool. The solid coated diamond particles improved the molten metal viscosity and restricted its ability to flow and fuse. And as the coated diamond particle content further increased, the porosity and the laser absorptivity of the mixed powder further increased, the powder fluidity and the thermal conductivity further decreased [6]. + +Figure 7f shows the CTE curves of the 1 and 3 vol.% titanium and copper-coated diamond/copper composites, as well as copper upon raising the temperature from 30 to 400 ◦C. The CTE of 1 and 3 vol.% copper-coated diamond/copper composites was significantly below that of corresponding titanium-coated diamond/copper composites, the minimum CTE of 1 vol.% copper-coated diamond/copper combined materials was very close to that of the copper. The shear stress had little effect on the interior of the coppercoated diamond/copper composite during heating, demonstrating the interfacial bonding of the copper matrix and diamond was better in the 1 vol.% copper-coated diamond/copper composites. During the heating process, the shear stress along the interface of the copper matrix and diamond had little effect on the CTE and led to a better thermal stability. This highlighted the advantages of copper plating on diamond particle surfaces, while the copper-coated diamond/copper composite properties reached high levels in comparison. + +# 4. Conclusions + +SLM technology was used to form titanium and copper-coated diamond/copper composites. The microstructure, roughness, interface bonding, thermal and mechanical performance were studied. + +(1) The values of roughness gradually decreased with increasing laser power. When the laser power reached the maximum value of 180 W, the surface roughness (Sa) reached a minimum value of 5.751 µm. The surface roughness Sa reached a minimum at the scanning rate of 200 mm/s. +(2) 1 vol.% copper-coated diamond/copper composite sample showed relatively best interface bonding of the copper matrix and diamond reinforcement, corresponding the lowest CTE and the strongest bending strength. + +(3) 1 vol.% copper-coated diamond/copper composites had the highest TC (336 W/mK) at an energy density of 300 J/mm3 (180 W, 200 mm/s). 3 vol.% copper-coated diamond/copper composites had the moderate TC (162 W/mK, 533 J/mm3, 160 W, 100 mm/s). 1 vol.% titanium-coated diamond/copper composites had the moderate TC (174 W/mK, 360 J/mm3, 180 W, 200 mm/s). 3 vol.% titanium-coated diamond/copper composites had the lowest TC (57 W/mK, 280 J/mm3, 140 W, 200 mm/s). The copper powders had the moderate TC (183 W/mK, 171 J/mm3, 180 W, 350 mm/s). + +The article offered electroless plating and evaporation methods for SLM to coat copper and titanium on the diamond particle surface for modifying and improving the wettability of diamond/copper interface, which opened up a new way for laser 3D printing technology to print a broad range of diamond-particle-reinforced MMCs. Thereby, unleashing their full potential for electronic package and thermal management applications. + +Author Contributions: L.Z. and Q.S. designed the experiments and wrote manuscript; L.H. provided the initial idea of this paper and financial support; Y.L. revised the manuscript; M.S. and H.G. performed the experiments; S.L. improved the language; Q.C. and P.G. analyzed the data. All authors have read and agreed to the published version of the manuscript. + +Funding: This research is funded by Wuhan Applied Foundational Frontier Project from Wuhan Science and Technology Bureau Project, China (No. 2020010601012172), the National Natural Science Foundation of China (No. 61805095, No. 51675496, No. 51902295), and the Fundamental Research Funds for the Central Universities, China University of Geosciences (Wuhan) (No. CUG2021234). + +Data Availability Statement: Not applicable. + +Conflicts of Interest: The authors declare no conflict of interest. + +# References + +1. Sharma, D.K.; Mahant, D.; Upadhyay, G. Manufacturing of metal matrix composites: A state of review. Mater. Today Proc. 2020, 26, 506–519. [CrossRef] +2. Silvain, J.F.; Heintz, J.M.; Veillere, A.; Constantin, L.; Lu, Y. A review of processing of Cu/C base plate composites for interfacial control and improved properties. Int. J. Extrem. Manuf. 2020, 2, 012002. [CrossRef] +3. 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等文件。 + +转换状态:18 / 18 完成。 + +## 01_CVD金刚石薄膜与热扩散片 + +对应路线:CVD 金刚石薄膜、独立膜、热扩散片、大面积基板。 + +- `1_Diamond as the heat spreader for the thermal dissipation of GaN-based electronic devices` +- `2_Heat-spreading diamond films for GaN-based high-power transistor devices` +- `3_State-of-the-art synthesis and post-deposition processing of large area CVD diamond substrates for thermal management` +- `4_Fabrication of free-standing diamond membranes` +- `5_Multilayer diamond heat spreaders for electronic power devices` + +可用于 PPT 的证据: + +- CVD diamond / diamond films are a recognized route for high-power electronic thermal management. +- Free-standing diamond membranes have been demonstrated from diamond-on-silicon films; converted MD records a 15 mm diameter membrane with about 12 um thickness. +- This category supports the proposal's Route B: thin diamond heat spreader in a constrained thermal stack. + +## 02_金属化金刚石载板与器件封装 + +对应路线:金属化 CVD 金刚石、激光器 submount、器件封装、可靠性。 + +- `6_Au-(Ti-W) and Au-Cr metallization of chemically vapor-deposited diamond substrates for multichip module applications` +- `7_Advanced metallization schemes for bonding of InP-based laser devices to CVD-diamond heatsinks` +- `8_Thermal stability of metallized CVD diamond.` +- `9_Increasing the output power of single 808-nm laser diodes using diamond submounts produced by microwave plasma chemical vapour deposition` + +可用于 PPT 的证据: + +- Au/Cr and Au/(Ti-W) metallization systems on CVD diamond have published adhesion and annealing evidence; one converted paper reports adhesion above 9 klbf/in2 for Au/(Ti-W) and 11.8 klbf/in2 for Au/Cr under optimized conditions. +- Metallized CVD diamond stability is a known package-level issue, supporting the pitch deck's claim that interface engineering is essential. +- A laser diode submount paper reports stable 8 W operation for more than 150 h on diamond with 700 W/mK thermal conductivity, and 12 W stable cw operation for 24 h on 1600 W/mK diamond. +- This category supports Route A: chip-on-diamond submount for XR optical engine, laser/LED, sensor and high-power module evaluation. + +## 03_改性粉末与CuDiamond复合热沉 + +对应路线:Ti/W/Cr 等界面改性、Cu/diamond 复合材料、复合热沉、微通道热沉。 + +- `10_High thermal conductive copper-diamond composites:state of the art` +- `11_Microstructure and thermal conductivity of Cu-diamond composites with Ti-coated diamond particles produced by gas pressure infiltration` +- `12_Enhanced thermal conductivity in diamond0copper composites with tungsten coatings on diamond particles prepared by magnetron sputtering method` +- `13_Optimized thermal properties in diamond particles reinforced copper-titanium matrix composites produced by gas pressure infiltration` +- `14_Effect of coating on the microstructure and thermal cond` +- `15_The Influence of the Carbide-Forming Metallic Additives (W, Mo, Cr, Ti) on the Microstructure and Thermal Conductivity of Copper-Diamond Composites` +- `16_Scalable monolayer-functionalized nanointerface for thermal conductivity enhancement in copper-diamond composite` +- `17_Diamond-Cu composites microchannel heat sink for effective thermal management of SiC power devices` +- `18_Fabrication of Titanium and Copper-Coated Diamond-Copper Composites via Selective Laser Melting` + +可用于 PPT 的证据: + +- Cu/diamond composites are a mature research route for combining high thermal conductivity with lower CTE than pure copper. +- Converted SLM paper reports 336 W/mK for 1 vol.% copper-coated diamond/copper composite under optimized parameters. +- Interface modification, coating and carbide-forming additives are recurring evidence themes; this supports the proposal's focus on interfacial thermal resistance, not only bulk thermal conductivity. +- This category supports the PoC's diamond-metal submount / heat spreader comparison and future shaped or microchannel heat sink options. + +## PPT 调用建议 + +优先把论文证据用于“可行性”和“机制支撑”,避免把文献数据写成博志金钻产品保证值。建议表达: + +- Literature supports the feasibility of CVD diamond heat spreaders, metallized diamond submounts and Cu/diamond composites. +- The proposed PoC will validate system-level thermal resistance and hotspot reduction in a Samsung-defined mobile stack. +- Final metrics should be measured on Bozhi Jinzhuan samples under agreed Samsung test conditions.