Microjoining and nanojoining
i WPNL2204
ii Related titles: Advanced welding processes (ISBN 978-1-84569-130-1) This book introduces the range of advanced welding techniques currently in use. It covers gas tungsten arc welding (GTAW), gas metal arc welding (GMAW), high-energy density processes such as laser welding, and narrow-gap welding methods. The book reviews general issues such as power sources, filler materials and shielding gases. Particular attention is given to monitoring and process control as well as to automation and robotics. New developments in advanced welding (ISBN 978-1-85573-970-3) Recent developments in high-technology areas have significantly transformed the welding industry where automation, computers, process control, sophisticated scientific instruments and advanced processing methods are all common. Today’s engineers and technologists have to support complex systems and apply sophisticated welding technologies. This comprehensive book discusses the changes in advanced welding technologies preparing the reader for the modern industry. Adhesive bonding (ISBN 978-1-85573-741-9) This important collection reviews key research on adhesive behaviour and applications in sectors as diverse as construction and automotive engineering. The book is divided into three main parts: fundamentals, mechanical properties and applications. Part I focuses on the basic properties of adhesives, surface assessment and treatment. Part II concentrates on understanding how adhesives perform under stress and the factors affecting fatigue and failure. The final part of the book reviews industry-specific applications in areas such as building and construction, transport and electrical engineering. Details of these and other Woodhead Publishing materials books, as well as materials books from Maney Publishing, can be obtained by: • visiting our web site at www.woodheadpublishing.com • contacting Customer Services (e-mail:
[email protected]; fax: +44 (0) 1223 893694; tel.: +44 (0) 1223 891358 ext. 130; address: Woodhead Publishing Limited, Abington Hall, Abington, Cambridge CB21 6AH, England) If you would like to receive information on forthcoming titles, please send your address details to: Francis Dodds (address, tel. and fax as above; e-mail:
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Microjoining and nanojoining Edited by Y. Zhou
Woodhead Publishing and Maney Publishing on behalf of The Institute of Materials, Minerals & Mining WPNL2204
CRC Press Boca Raton Boston New York Washington, DC
WOODHEAD
PUBLISHING LIMITED
Cambridge, England
WPNL2204
iv Woodhead Publishing Limited and Maney Publishing Limited on behalf of The Institute of Materials, Minerals & Mining Woodhead Publishing Limited, Abington Hall, Abington Cambridge CB21 6AH, England www.woodheadpublishing.com Published in North America by CRC Press LLC, 6000 Broken Sound Parkway, NW, Suite 300, Boca Raton, FL 33487, USA First published 2008, Woodhead Publishing Limited and CRC Press LLC © 2008, Woodhead Publishing Limited The authors have asserted their moral rights. This book contains information obtained from authentic and highly regarded sources. Reprinted material is quoted with permission, and sources are indicated. Reasonable efforts have been made to publish reliable data and information, but the authors and the publishers cannot assume responsibility for the validity of all materials. Neither the authors nor the publishers, nor anyone else associated with this publication, shall be liable for any loss, damage or liability directly or indirectly caused or alleged to be caused by this book. Neither this book nor any part may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, microfilming and recording, or by any information storage or retrieval system, without permission in writing from Woodhead Publishing Limited. The consent of Woodhead Publishing Limited does not extend to copying for general distribution, for promotion, for creating new works, or for resale. Specific permission must be obtained in writing from Woodhead Publishing Limited for such copying. Trademark notice: Product or corporate names may be trademarks or registered trademarks, and are used only for identification and explanation, without intent to infringe. British Library Cataloguing in Publication Data A catalogue record for this book is available from the British Library. Library of Congress Cataloging in Publication Data A catalog record for this book is available from the Library of Congress. Woodhead Publishing ISBN 978-1-84569-179-0 (book) Woodhead Publishing ISBN 978-1-84569-404-3 (e-book) CRC Press ISBN 978-1-4200-7083-5 CRC Press order number: WP7083 The publishers’ policy is to use permanent paper from mills that operate a sustainable forestry policy, and which has been manufactured from pulp which is processed using acid-free and elementary chlorine-free practices. Furthermore, the publishers ensure that the text paper and cover board used have met acceptable environmental accreditation standards. Project managed by Macfarlane Book Production Services, Dunstable, Bedfordshire, England (
[email protected]) Typeset by Replika Press Pvt Ltd, India Printed by TJ International Limited, Padstow, Cornwall, England
WPNL2204
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Contents
Contributor contact details
xiii
Introduction
xix
Y ZHOU, University of Waterloo, Canada
Part I Basics of microjoining 1
Mechanisms of solid-state bonding processes
3
J E GOULD, Edison Welding Institute, USA
1.1 1.2 1.3 1.4 1.5 1.6 1.7 1.8 1.9 1.10
Introduction Mechanisms of solid-state bonding Contaminant displacement/interatomic bonding Extension of the contacting surfaces Separation of the contaminated areas Realignment of the grain structures for bonding Thermal dissolution of oxides/contaminants Breakdown of the interfacial structure Summary References
3 4 7 7 8 14 14 18 21 23
2
Mechanisms of soldering and brazing
25
T TAKEMOTO, Osaka University, Japan
2.1 2.2 2.3 2.4 2.5 2.6 2.7
Introduction Definition of soldering and brazing Basic metallurgical reaction during soldering and brazing Materials for soldering and brazing Soldering and brazing process Summary and future trends References
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vi
Contents
3
Fundamentals of fusion microwelding
51
G A KNOROVSKY, Sandia National Laboratories, USA and V V SEMAK, Pennsylvania State University, USA
3.1 3.2 3.3 3.4 3.5 3.6 3.7 3.8 3.9
Introduction Fundamental aspects Forces acting on the pool Practical aspects Some problems in applying fusion microjoining Examples: laser, arc and resistance welding Summary and conclusions Acknowledgements References
51 52 54 73 76 81 87 88 88
4
Modeling of solid state bonding
91
Y TAKAHASHI, Osaka University, Japan
4.1 4.2 4.3 4.4 4.5 4.6 4.7
Introduction Viscoplastic deformation model and interfacial deformation Thermocompression bonding Thermosonic bonding Numerical simulation of lap resistance welding Concluding remarks References and further reading
91 91 93 107 115 117 119
5
Modeling of fusion microwelding
121
B H CHANG, Tsinghua University, P. R. China
5.1 5.2 5.3 5.4 5.5 5.6 5.7 5.8
Introduction Features of thermal process in fusion microwelding Modeling of conductive heat transfer Modeling of convective heat transfer Modeling of resistance microwelding Summary and future trends Acknowledgements References
121 122 124 128 142 154 156 156
6
Sensing, monitoring and control
159
M MAYER, University of Waterloo, Canada
6.1 6.2 6.3 6.4
Introduction Definitions and methods Signals from joining processes Applications
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Contents
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6.5 6.6
Future trends References
169 171
7
Assembly process automation and materials handling
174
Y M CHEUNG and D LIU, ASM Assembly Automation Ltd, Hong Kong
7.1 7.2 7.3 7.4 7.5
7.6 7.7
Introduction Assembly equipments for an in-line process Material handling in assembly equipment Control of process parameters for assembly processes Case study: design of customized assembly equipment – an alignment laser welder for a fibre pigtailed TO-can laser diode package Future trends References
174 175 178 183
187 200 201
Part II Microjoining and nanojoining processes 8
Microelectronics wire bonding
205
I LUM, M MAYER and Y ZHOU, University of Waterloo, Canada
8.1 8.2 8.3 8.4 8.5 8.6 8.7 8.8
Introduction Wire bonding process Process parameters Mechanisms of bond formation Quality control Equipment and fixturing Current directions References
205 206 212 216 222 224 226 229
9
Solid-state diffusion bonding
234
A SHIRZADI, University of Cambridge, UK
9.1 9.2 9.3 9.4 9.5 9.6 9.7 9.8
Basic definition Solid-state bonding in comparison with other joining methods Main bonding parameters and apparatus Theoretical aspects and modelling of solid-state diffusion bonding Various approaches to overcome surface oxide problem Non-chemical surface oxide removal using liquid gallium Summary References
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Contents
10
Bonding using nanoparticles
250
A HIROSE, Osaka University, Japan and K F KOBAYASHI, Fukui University of Technology, Japan
10.1 10.2
10.5 10.6 10.7
Introduction Structure and thermal characteristics of composite Ag nanoparticle Bonding of various metals using composite Ag nanoparticle Effects of bonding conditions on bondability of Cu-to-Cu joints Bonding of Si chip Conclusion and future trends References
11
Diffusion soldering and brazing
10.3 10.4
250 251 255 259 265 267 267 269
M L KUNTZ and Y ZHOU, University of Waterloo, Canada
11.1 11.2 11.3 11.4 11.5 11.6 11.7
Introduction to diffusion soldering/brazing Process description of diffusion soldering/brazing Diffusion soldering/brazing process mechanics Evaluating joint properties Modeling of diffusion soldering/brazing Summary and future trends References
269 273 278 285 289 296 296
12
Laser soldering
299
Y H TIAN and C Q WANG Harbin Institute of Technology, P. R. China
12.1 12.2 12.3 12.4 12.5 12.6 12.7 12.8
Introduction Overview of laser soldering Fundamentals of laser soldering Fluxless laser soldering Reliability of solder joint Summary and future trends Acknowledgments References
299 299 301 312 314 322 324 324
13
Fluxless soldering
327
J P JUNG, University of Seoul, Korea and S M HONG, Samsung Electronics Co. Ltd., Korea
13.1 13.2 13.3
Flux in soldering Demand for fluxless soldering Fluxless soldering process
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Contents
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13.4 13.5
Summary References
342 342
14
Laser microwelding
345
I MIYAMOTO, Osaka University Japan and G A KNOROVSKY, Sandia National Laboratories, USA
14.1 14.2 14.3 14.4 14.5 14.6 14.7 14.8 14.9 14.10 14.11 14.12
Introduction Fundamentals of laser microwelding Thermal modeling of laser microwelding Weld defects Laser microwelding technologies Evaluation of microweld joints Applications of laser microwelding Novel laser microwelding technologies Future trends Acknowledgements Equation annex References
345 347 363 368 377 387 391 396 408 410 410 412
15
Electron beam microwelding
418
G A KNOROVSKY, Sandia National Laboratories, USA, T DORFMÜLLER, U DILTHEY and K WOESTE, RWTH – Aachen University, Germany
15.1 15.2 15.3 15.4 15.5 15.6 15.7
Introduction Technology Microwelding process Examples of joining and applications Summary Acknowledgements References
418 419 445 459 471 471 471
16
Resistance microwelding
473
S Fukumoto, University of Hyogo, Japan, Y ZHOU, University of Waterloo, Canada and W TAN Medtronic Inc., USA
16.1 16.2 16.3 16.4 16.5 16.6 16.7 16.8 16.9
Introduction Fundamentals Process variations Process conditions: welding parameters Process conditions: surface Dynamic resistance and process control Equipment Summary and future trends References
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Contents
17
Adhesive bonding
500
S. BÖHM, E STAMMEN, G HEMKEN and M WAGNER, Technical University Braunschweig, Germany
17.1 17.2 17.3 17.4 17.5 17.6 17.7 17.8 17.9 17.10 17.11 17.12
Introduction Adhesive bonding as joining technology Adhesive bonding process: pre-treatment Dispensing and mixing of adhesives Curing and setting of adhesives Quality control Principles of adhesive selection and examples of applications Microelectronic interconnections and packaging Other applications of adhesive bonding Future trends Sources of further information References
500 500 512 514 519 521 522 524 535 537 538 543
18
Introduction to nanojoining
545
S SAHIN, Celal Bayar University, Turkey and M YAVUZ and Y ZHOU, University of Waterloo, Canada
18.1 18.2 18.3 18.4
Introduction Nanojoining methods Summary and future trends References
545 546 576 576
Part III Microjoining of materials and applications of microjoining 19
Joining of high temperature superconductors
583
G ZOU, Tsinghua University, P. R. China
19.1 19.2 19.3 19.4 19.5 19.6 19.7 19.8 19.9 19.10 19.11
Introduction Superconducting materials Processing technologies of high temperature superconductors Needs for joining high temperature superconductors Joining of BSCCO bulks Soldering of BSCCO tapes Diffusion bonding of BSCCO tapes Joining of YBCO bulks Conclusion and future trends Acknowledgements References
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Contents
20
Joining of shape memory alloys
xi
620
K UENISHI and K F KOBAYASHI, Osaka University, Japan
20.1 20.2 20.3 20.4 20.5 20.6 20.7 20.8
Introduction Basics of shape memory alloys Application of SMAs Background on welding and laser processing of SMA Brazing of nithinol Laser microwelding Future trends References
620 620 623 624 626 628 630 630
21
Wafer bonding
633
J WEI and Z SUN, Singapore Institute of Manufacturing Technology, Singapore
21.1 21.2 21.3 21.4 21.5 21.6 21.7
Introduction Direct wafer bonding Anodic wafer bonding Wafer bonding via intermediate layers Bonding of dissimilar materials Concluding remarks References
633 635 641 648 655 658 659
22
Plastics microwelding
665
I JONES, TWI Ltd, UK
22.1 22.2 22.3 22.4 22.5 22.6 22.7 22.8
Introduction Theory of welding plastics Introduction to welding processes Processes for microwelding of plastics Ultrasonic welding Laser welding Future trends References
665 665 671 673 673 678 689 689
23
Microjoining in medical components and devices
691
K J ELY, Boston Scientific, USA and Y ZHOU, University of Waterloo, Canada
23.1 23.2 23.3 23.4 23.5
Introduction Materials challenges Medical components and devices Joint design and process selection Testing and verification
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Contents
23.6 23.7
Summary and future trends References
715 716
24
Hermetic sealing of solid oxide fuel cells
718
M BROCHU, McGill University, Canada and R E LOEHMAN, Sandia National Laboratories, USA
24.1 24.2 24.3 24.4 24.5
Introduction Materials involved in the fabrication of a SOFC Different approaches for sealing SOFCs Summary References
718 720 722 736 737
25
Joining of bulk nanostructured materials
741
M BROCHU, McGill University, Canada
25.1 25.2 25.3 25.4 25.5 25.6 25.7
Introduction Fabrication processes Non-equilibrium state of microstructure Attempts in joining bulk nanomaterials Metallic glasses Summary References
741 742 745 747 752 754 754
26
Ceramic/metal bonding
758
A WU, Tsinghua University, P.R. China
26.1 26.2 26.3 26.4 26.5 26.6 26.7 26.8
Introduction Characteristics of typical ceramic materials General difficulties in joining of ceramic/metal Brazing ceramics to metal Vacuum diffusion bonding Typical applications Summary and future trends References
758 758 761 767 771 780 784 784
Index
786
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Contributor contact details
(* = main contact)
Editor
Chapter 2
Professor Y. (Norman) Zhou Microjoining Laboratory Centre for Advanced Materials Joining University of Waterloo 200 University Avenue West Waterloo Ontario Canada N2L 3G1
Professor Tadashi Takemoto Joining and Welding Research Institute Osaka University 11-1 Mihogaoka Ibaraki-shi Osaka 567-0047 Japan E-mail: takemoto@jwri. osakau.ac.jp
E-mail:
[email protected]
Chapter 3 Chapter 1 Dr Jerry E. Gould Resistance and Solid State Welding Edison Welding Institute 1250 Arthur E. Adams Drive Columbus, OH 43221 USA
Dr Gerald A. Knorovsky* Sandia National Laboratories Joining & Coating Department Albuquerque New Mexico USA E-mail:
[email protected]
E-mail:
[email protected] Dr Vladimir V. Semak Electro-Optics Center Pennsylvania State University 222 Northpointe blvd. Freeport, PA 116229 USA E-mail:
[email protected]
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Contributor contact details
Chapter 4
Chapter 7
Professor Yasuo Takahashi Osaka University Center for Advanced Science and Innovation (CASI) 2-1, Yamadaoka Suita Osaka 565-0871 Japan
Dr Y. M. Cheung* and Dr Deming Liu ASM Assembly Automation Ltd 4/F Watson Centre 16 Kung Yip Street Kwai Chung Kowloon Hong Kong
E-mail:
[email protected]
E-mail:
[email protected] [email protected]
Chapter 5
Chapter 8
Professor Baohua Chang Department of Mechanical Engineering Tsinghua University Beijing 100084 P.R. China
Ivan Lum*, Professor Michael Mayer and Professor Y. (Norman) Zhou Microjoining Laboratory Centre for Advanced Materials Joining University of Waterloo 200 University Avenue West Waterloo Ontario Canada N2L 3G1
E-mail:
[email protected]
Chapter 6
E-mail:
[email protected] [email protected] [email protected]
Professor Michael Mayer Microjoining Laboratory Centre for Advanced Materials Joining University of Waterloo 200 University Avenue West Waterloo Ontario Canada N2L 3G1
Chapter 9
E-mail:
[email protected]
Dr Amir Shirzadi Phase Transformations & Complex Properties Research Group Department of Materials Science & Metallurgy University of Cambridge Pembroke Street Cambridge CB2 3QZ UK E-mail:
[email protected]
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Chapter 10
Chapter 12
Professor Akio Hirose* Division of Materials and Manufacturing Science Graduate School of Engineering Osaka University 2-1 Yamadaoka Suita Osaka 565-0871 Japan
Professor Yanhong Tian and Professor Chunqing Wang P.O. Box 436 State Key Laboratory of Advanced Welding Production Technology School of Materials Science and Engineering Harbin Institute of Technology Harbin P.R. China 150001
E-mail:
[email protected]
E-mail:
[email protected] Professor Kojiro F. Kobayashi Department of Mechanical Engineering Fukui University of Technology 3-6-1 Gakuen Fukui 910-8505 Japan E-mail:
[email protected]
Chapter 11 Professor Michael L. Kuntz* and Professor Y. (Norman) Zhou Microjoining Laboratory Centre for Advanced Materials Joining University of Waterloo 200 University Avenue West Waterloo Ontario Canada N2L 3G1 E-mail:
[email protected];
[email protected]
Chapter 13 Professor Jae Pil Jung* Department of Material Science and Engineering University of Seoul 90 Junnong-dong Dongdemun-gu Seoul 130-743 Korea E-mail:
[email protected] Dr Soon Min Hong Micro-joining Laboratory Institute of Intelligent System Mechatronics Center Samsung Electronics Co. Ltd Kyungki-Do 442-742 South Korea E-mail:
[email protected]
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Contributor contact details
Chapter 14
Chapter 16
Dr Gerald A. Knorovsky Sandia National Laboratories Joining & Coating Department Albuquerque New Mexico USA
Professor Shinji Fukumoto* Graduate School of Engineering University of Hyogo 2167 Shosha Himeji Hyogo 671-2201 Japan
E-mail:
[email protected] E-mail:
[email protected]
Professor Isamu Miyamoto Osaka University 2-1 Yamada-Oka Suita Osaka 565-0871 Japan E-mail:
[email protected]
Chapter 15 Dr Gerald A. Knorovsky Sandia National Laboratories Joining & Coating Department Albuquerque New Mexico USA
Professor Y. (Norman) Zhou Microjoining Laboratory Centre for Advanced Materials Joining University of Waterloo 200 University Avenue West Waterloo Ontario Canada N2L 3G1 E-mail:
[email protected]
E-mail:
[email protected] Dipl.-Ing Thomas Dorfmüller, Professor Dr-Ing Ulrich Dilthey and Dr-Ing Klaus Woeste, Welding and Joining Institute (ISF) RWTH - Aachen University Germany
Dr Wen Tan Medtronic Inc. Medtronic Energy and Component Center 6800 Shingle Creek Parkway Brooklyn Center Minneapolis, MN 55430 USA E-mail:
[email protected]
E-mail:
[email protected] [email protected] [email protected]
WPNL2204
Contributor contact details
Chapter 17 Professor Dr-Ing Stefan Böhm*, Dipl.-Chem Elisabeth Stammen, Dipl.-Ing Gregor Hemken and Dipl.-Ing Mario Wagner Technical University Braunschweig Institute of Joining and Welding (IFS) Langer Kamp 8 38106 Braunschweig Germany
xvii
Professor Mustafa Yavuz Nano- and Micro-systems Research Laboratory Department of Mechanical and Mechatronics Engineering University of Waterloo 200 University Avenue West Waterloo Ontario Canada N2L 3G1
E-mail:
[email protected]
E-mail:
[email protected]
Chapter 18
Chapter 19
Dr Salim Sahin* Department of Mechanical Engineering Celal Bayar University 45140 Manisa Turkey
Professor Guisheng Zou Department of Mechanical Engineering Tsinghua University Beijing 100084 P.R. China
E-mail:
[email protected] Professor Y. (Norman) Zhou Microjoining Laboratory Centre for Advanced Materials Joining University of Waterloo 200 University Avenue West Waterloo Ontario Canada N2L 3G1 E-mail:
[email protected]
E-mail:
[email protected]
Chapter 20 Professor Keisuke Uenishi* and Professor Kojiro F. Kobayashi Department of Management of Industry and Technology Graduate School of Engineering Osaka University 2-1 Yamadaoka Suita Osaka 565-0871 Japan E-mail:
[email protected]
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Contributor contact details
Chapter 21
Chapter 24
Dr Jun Wei* and Dr Zheng Sun Singapore Institute of Manufacturing Technology 71 Nanyang Drive Singapore 638075
Professor M. Brochu* McGill University Mining and Materials Engineering Department 3610 University Street Montreal Canada H3A 2B2
E-mail:
[email protected] [email protected]
Chapter 22 Dr Ian Jones TWI Ltd Granta Park Great Abington Cambridge CB21 6AL UK
E-mail:
[email protected] Dr R. E. Loehman Sandia National Laboratories Advanced Materials Laboratories 1001 University SE Albuquerque New Mexico 87106 USA
E-mail:
[email protected]
Chapter 25 Chapter 23 Dr Kevin J. Ely Boston Scientific Corporation 4100 Hamline Avenue St Paul, MN 55112 USA E-mail:
[email protected] Professor Y. (Norman) Zhou University of Waterloo 200 University Avenue West Waterloo Ontario Canada N2L 3G1 E-mail:
[email protected]
Professor M. Brochu McGill University Mining and Materials Engineering Department 3610 University Street Montreal Canada H3A 2B2 E-mail:
[email protected]
Chapter 26 Professor Aiping Wu Department of Mechanical Engineering Tsinghua University Beijing 100084 P.R. China E-mail:
[email protected] WPNL2204
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Introduction
Y Z H O U , University of Waterloo, Canada
Joining, whether at macro-, micro-, or nano-scale, is an essential part of man-made product manufacturing and assembly, providing mechanical coupling and support, electrical connection or insulation, environmental protection, etc. This is especially true with the growing trend of product miniaturization over the past few decades. Effective microjoining has become one of the most critical technical prerequisites for success in manufacturing at ever-smaller scales. Microjoining is an important part of microelectronic packaging and interconnection, but actually covers a much broader area and is crucial for manufacturing many other miniature components, devices and systems, such as medical implants, sensors and transducers, batteries, and optoelectronics (Table 1). Table 1 Examples of micro- and nano-scale components, devices and systems Medical components and devices
Ultrasound catheters, endoscopes, neurostimulators, pacemaker and defibrillator implants, implantable radioactive capsules, cochlear implants
Sensors and transducers
Airbag sensors, strain gauges, automotive deceleration sensors, bi-metal contacts, smoke detectors
Micro- and optoelectronics
Fibre-optic connectors (assemblies), microwave modules, fuel and solar cells, glass-to-metal seals, integrated circuits
Micro-systems
Micro-gears, micro-turbines, micro-pumps, micro-motors
Other microproducts
Battery packs and cells, incandescent and vapour phase lamps, ink-jet cartridges, relays and thermostats
Nanoscale devices and systems
Nano-robots, nano-sensors, nano-electronics, nano-structures
While microjoining is a popular term, it is very loosely defined and mostly used relative to conventional (regular, large- or macro-scale) welding and joining. A first criterion in defining ‘microjoining’ could be based on the dimensions of the parts (building blocks) to be joined. For example, the
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Introduction
joining of parts with a characteristic dimension, e.g., sheet thickness or wire diameter, up to 100 µm, could be called microjoining. Then, anything in between microjoining and macrojoining (or millijoining), with dimensions above 1 mm, could be sub-millimeter (or meso-) joining. Similarly, we can define nanojoining and sub-micrometer joining. But this strict definition is not only sometimes problematic as will be discussed later, but also has the potential to become a little complicated and artificial. There are in fact many other ways to define microjoining, such as based on industrial or application areas, process characteristics, the equipment used (e.g., the precision of such a system), the materials to be joined and even the scales of microwelds. For example, (regular) resistance welding is mostly used in joining steels or aluminium alloys as in autobody and appliance assembly with sheet thickness or wire diameter often above 0.5 mm, while resistance microwelding is used predominantly to join non-ferrous materials in the fabrication of precision components and devices with sheet thickness or wire diameter in an approximate range from 20 to 400 µm, such as in batteries, printed circuit boards, relays, sensors, air-bag diffuser screens, and medical devices. There is sometimes no clear distinction in resistance welding and microwelding in terms of sheet thickness. But each uses a different line of equipment, often provided by two distinct groups of equipment manufacturers. Another example is wire bonding used for semiconductor chip-level interconnection, in which the diameter of bond wire is mostly around 25 µm but can be up to a few hundred micrometers especially for power electronics. It would be difficult to say that a joint with wire diameter less than 100 µm is microjoining and the other is not since both use the same equipment. Similar arguments may be applied to nanojoining. For the most part of this publication, ‘microjoining’ is employed as a flexible term, such as relative to macrojoining, to include joining of sheet thickness or wire diameter up to 0.2–0.5 mm for applications in manufacturing of miniature products. On the other hand, nanojoining is a new and somewhat distinct field with very little published information yet available, such as that on joining of nanotubes and nanowires. Many microjoining or microwelding processes can trace their origins back to the 1950s. For example, a capacitor discharge machine was introduced then for resistance welding of orthodontic supplies, and this development was later adopted in the aerospace industry in the 1960s [1]. Thermocompression wire bonding, a variant of hot pressure welding, was developed by Bell Laboratories in 1957 [2] and later evolved into ultrasonic wire bonding in the 1960s and thermosonic wire bonding, a combination of thermocompression and ultrasonic wire bonding, in 1970 [3]. The wire bonding processes have been one of the key chip-level interconnection technologies in the success of microelectronic products. While resistance microwelding and wire bonding were developed for microjoining applications, electron
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xxi
beam welding, developed in 1958 [4] is suitable for both macro- and microscale joining because of the precision beam quality, and fixturing and controls integrated with the system. There is no doubt that the growth of many microjoining processes and techniques have been tied to the growth of the microelectronics, medical, aerospace, and defence industries. While many microjoining processes and applications are now thought of as mature, microjoining continues to face challenges because of ever advancing miniaturization. For example, novel microjoining processes are urgently needed for packaging and interconnecting in MEMS (MicroElectroMechanical Systems), also called microsystems or micromachines, where individual electrical, mechanical, fluidic, and optical components needs to be connected and coupled to the macroscopic external environment [5]. The latter causes a particular problem in sensor technology where, on the one hand, systems must be protected from mechanical damage and corrosion, and on the other hand, be exposed as widely as possible to the environment to obtain true, undistorted physical and chemical values [5]. For this reason, ‘packaging and interconnecting techniques will play the key role in any further industrial expansion (of microsystems)’ [5]. Further down the road of miniaturization, as in the micro- and macro-world, there is an emerging need to join nanoscale building blocks, such as nanowires and nanotubes, to themselves to form nanoscale devices and systems, and then to join these to the surroundings, to be integrated into micro- and macro-scale devices and systems. In fact, several nanojoining or nanowelding processes have been already attempted (e.g., see ref. [6]). The objective of all joining processes, whether macro-, micro-, or nano, is to produce permanent union of or connections between parts or building blocks to be assembled, through the formation of primary (and occasionally secondary) chemical bonds between faying surfaces. An interlayer or intermediate material may be needed when the parts are not compatible in atomic structure, for example as in ceramic/metal joining, or when a locally reduced melting temperature is needed to facilitate carrying out the joining process. The foregoing definition, and the scope of this volume, excludes mechanical joining or fastening, since in the latter, the original free surfaces of the materials in the joints are retained unaltered. In principle, two ideal solid surfaces, e.g. both perfectly clean and atomically flat, will bond together if brought into intimate contact, because they will be drawn together spontaneously by interatomic forces [7]. This theoretical principle has been confirmed in tightly controlled environments, such as in direct silicon wafer bonding [8]. However, most engineering surfaces are characterized as rough and contaminated, requiring some form of energy, usually heat and/or pressure, to be applied to overcome these surface impediments to make a joint. Roughly, there are four major types of joining processes: solid-state bonding, soldering/brazing, fusion welding, and adhesive
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Introduction
bonding. Solid-state bonding is mainly achieved by macroscopic or microscopic deformation where no melting occurs. Fusion welds and brazed/soldered joints are achieved by melting and epitaxial solidification of molten metal, but localized melting of base metals occurs in fusion welding, and (in initial stages at least) only the melting of filler metals occurs in brazing/soldering. While the melting points in brazing are above 450°C, those in soldering are below 450°C. If liquid is removed during the process, solid-state joints can be also produced by certain fusion welding, soldering and brazing processes, e.g. as in resistance cross-wire microwelding [9]. In adhesive bonding, organic adhesives are used to join parts. Despite the importance of microjoining, there has never been a single, comprehensive book on the topic except for a short pamphlet and later an expanded version (still short at 34 pages on 300 × 210 mm paperback) more than 20 years ago [10, 11]. Descriptions of many microjoining processes and techniques can be found in microelectronics packaging books, e.g. the book by Tummala [12] is one excellent example. But those books generally adopt the package (system) point of view and cover very broad topics, from design, materials and processing, to manufacture, testing, performance and reliability; therefore, insight into microjoining process details themselves are rather limited. This book aims to provide an up-to-date overview of a wide range of microjoining processes and techniques, including related fundamental and technological aspects, e.g., increasing effects of surface properties as the parts continue to downsize, and selections of joining techniques for specific applications or materials. Nanojoining processes are also introduced because of this important, emerging field. This book is divided into three parts covering the most commonly used microjoining and nanojoining processes (Table 2). Part I covers underlying principles of major microjoining groups, i.e., solid-state bonding, soldering/ brazing, and fusion microwelding, and common issues in microjoining, such as, computer modeling, sensing and monitoring, control and automation, and part handling, positioning and fixturing. Part II describes characteristics of various microjoining and nanojoining processes, including bonding mechanisms, important parameters and their effects, metallurgical aspects, engineering mechanics aspects (e.g., thermal stresses especially in joining of dissimilar materials), equipment, quality control, etc. In particular, the possibility of microjoining processes and techniques being extended for nanojoining is discussed in many chapters. Part III explains microjoining of selected materials and specific microjoining application in many different industries. Application examples can also be found in other sections where individual microjoining or nanojoining processes are described. It is intended that this book be a valuable reference for production engineers, designers and researchers using, studying and developing microjoining and nanojoining processes and techniques. Scientific and practical insight is offered
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Table 2 Microjoining processes (with chapter numbers indicated) Ceramics
Solid-state bonding
Anodic bonding (21) Bonding with nanoparticles (10) Cold welding (1) Diffusion bonding (1) Explosive welding (25) Friction welding (25) Friction stir welding (25)
Anodic bonding (21) Chip bonding with nanoparticles (10) Diffusion bonding (1) TAB inner lead bonding (4) Thermocompression wire bonding (1, 4, 8) Thermosonic wire bonding (4, 6, 7, 8) Ultrasonic wire bonding (6, 8) Wafer bonding (21)
Anodic bonding (21) Diffusion bonding (19, 26)
Soldering/ brazing
Diffusion soldering/ brazing (11) Dip soldering/brazing (2) Fluxless soldering/brazing (2, 12, 13) Furnace soldering/brazing (2) Induction soldering/brazing (2) Laser soldering/brazing (2, 12) Reflow soldering (2, 12, 13) Resistance soldering/brazing (16, 20) Wave soldering/brazing (2)
Diffusion soldering (11) Eutectic bonding (2) Flip chip bonding (2) Fluxless soldering, including acid vapour and plasma (2, 12, 13) Laser soldering (2, 12) Lead-free soldering (2) Reflow soldering (2)
Active brazing (23, 24, 26) Air brazing (24) Diffusion brazing (11) Superconductor soldering (19)
Fusion welding
Electron beam microwelding (3, 15, 20, 25) Laser microwelding (3, 5, 7, 14, 20, 23) Plasma microwelding (3, 5, 23) Percussive microwelding (1, 3) Resistance microwelding, including spot, seam, projection, parallel-gap and cross wire (1, 3, 5, 6, 16, 23) TIG microwelding (3, 5, 25)
Adhesive bonding
Adhesive bonding (17)
Adhesive die attachment (7, 17) Adhesive flip chip bonding (17) Adhesive sealing (17)
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Polymers
Gas flame welding (19) Glass sealing (24)
Friction welding, including spin, friction stir, ultrasonic (22, 23) Hot bar and impulse welding (22) Laser welding (22, 23) and other electromagnetic welding processes, such as microwave, induction, resistive implant (22)
Adhesive bonding (17)
Adhesive bonding (17, 23)
xxiii
Semiconductors
Introduction
Metals
xxiv
Introduction
in the areas of medical components and devices (especially implantable), sensors and transducers, micro- and opto-electronics, micro-systems, nanoscale devices and systems, fuel cells and other miniature products. It can also be a useful text for students in materials science and engineering, mechanical engineering, electrical and electronics engineering. A good understanding of joining requires multidisciplinary approaches from various fields, such as physics, chemistry, mathematics, solid and fluid mechanics, materials science (metallurgy), electrical engineering and electronics. A book like this is obviously impossible without the insightful input of all the collaborators. I am very proud of their wonderful contributions. I appreciate also the professional help from the editorial staff at Woodhead Publishing Limited. Finally I would like to dedicate this book to my wife, Wendy, and sons, Antony and Daniel, and my parents, Dad and Mom, for their encouragement, support and love.
References 1. A. Cullison, ‘Welding: A Heavyweight in a Miniature World’, Welding J., 75(5), May 1996, pp. 29–34. 2. O.L. Anderson, H. Christensen and P. Andreatch, ‘Technique for Connecting Electrical Leads to Semiconductors’, Journal of Applied Physics, 28(8), August 1957, p. 923. 3. G.G. Harman, Wire Bonding in Microelectronics – Materials, Processes, Reliability and Yield, 2nd ed. New York: McGraw-Hill, 1997. 4. J.A. Stohr and J. Eriola, ‘Vacuum Welding of Metals’, Welding and Metal Fabrication, 26, October 1958 pp. 366–374. 5. W. Menz, J. Mohr and O. Paul, Microsystem Technology, Weinheim, Germany, Wiley-VCH, 2001. 6. M.S. Fuhrer, J. Nygard, L. Shih, M. Forero, Y.-G. Yoon, M.S.C. Mazzoni, H.J. Choi, J. Ihm, S.G. Louie, A. Zettl and P.L. McEuen, ‘Crossed Nanotube Junctions’, Science, 288, April 21, 2000, pp. 494–497. 7. R.W. Messler, Jr., Principles of Welding: Processes, Physics, Chemistry and Metallurgy, New York: John Wiley & Sons, 1999. 8. U. Gösele, H. Stenzel, T. Martini, J. Steinkirchner, D. Conrad and K. Scheerschmidt, ‘Self-propagating Room Temperature Silicon Wafer Bonding in Ultrahigh Vacuum’, Appl. Phys. Lett., 67(24), 1995, pp. 3614–3616. 9. S. Fukumoto and Y. Zhou, ‘Mechanism of Resistance Microwelding of Crossed Fine Nickel Wires’, Metall. Mater. Trans. A, 35A, 2004, pp. 3165–3176. 10. D.D. Zimmerman and D.H. Lewin (eds), The Fundamentals of Microjoining Processes, Silver Spring, MD: International Society for Hybrid Microelectronics, 1983. 11. K.I. Johnson (ed.), Introduction to Microjoining, TWI: Abington, 1985. 12. R.R. Tummala (ed.), Fundamentals of Microsystems Packaging, New York, McGrawHill: 2001.
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Part I Basics of microjoining
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1 Mechanisms of solid-state bonding processes J E G O U L D, Edison Welding Institute, USA
1.1
Introduction
Solid-state bonding processes are those which accomplish bonding without the requirement of resolidification of liquid metal. Typically, these processes take advantage of applied strain and/or heat for facilitate joining. Joining is largely the result of intimate intermetallic contact in the absence of local protective films. Solid-state bonding processes are the oldest of joining processes, with the official AWS definition of forge welding requiring an anvil and a hammer.1 Solid-state bonding processes have proliferated particularly over the last several decades as new power systems have developed. General classifications of these processes include pressure processes (cold and hot pressure welding, etc.), resistance processes (butt, projection and seam welding, etc.), surface displacement processes (friction and ultrasonic welding, etc.), arc-heated processes (percussive welding, etc.), and diffusion bonding processes. Solidstate bonding has been developed and in some cases adopted for microjoining. For example, ultrasonic wire bonding (including thermosonic wire bonding), a variant of ultrasonic microwelding, is still the dominant technique for chip-level interconnection (Fig. 1.1). Cold pressure welding, and resistance seam and projection welding have been used to seal electronic packages (Fig. 1.2). In this chapter, mechanisms of bonding are described for those processes using both mechanically applied straining and heating. Detailed examinations of bonding mechanisms of the other processes are available in the references. These include the cold-pressure welding processes,2–8 and the diffusion bonding processes.9–12 (also see Chapter 9) and wire bonding (Chapter 8). These processes can be thought of as having two generally separable stages. These include a heating stage, and an upsetting stage. As such, these welding methods can be generally classified as heat and forge processes. Distinctions between these processes then are largely in how heat and forging are applied. Inevitably, however, heat is first applied. Mechanistically, this 3 WPNL2204
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Microjoining and nanojoining
1.1 Wire bonds used as electrical interconnections in microelectronics.
heat is used for two purposes. First, heating the workpieces reduces the yield strengths of the base materials, and permits forging to occur with high degrees of strain at reduced upsetting forces. Second, if heating is properly applied, upsetting creates high degrees of strain over a very localized area (at the bonding surface). Once the appropriate heating has been accomplished, forging (or upsetting) commences. Upsetting also has two major functions. These include collapsing asperities to create intimate contact, and displacing/ dispersing protective oxides and films to facilitate metal-to-metal bonding. Residual heat content/ heating is also considered advantageous, in order to further consolidate/homogenize the joint. This chapter will focus on the specific mechanisms operative when bonding metals using the heat/forge processes defined above. Specific mechanisms for different stages of these processes will be identified, and quantified using best-available theory. Further, these mechanisms will be used to understand the roles of temperature and strain in facilitating bonding with these classes of joining technologies.
1.2
Mechanisms of solid-state bonding
To attempt to define the specific mechanisms of bonding for solid-state welding processes, it is first necessary to have an understanding about the microstructural and surface conditions of the workpieces planned for joining. On a microscopic scale, the surfaces for bonding have been well categorized as irregular and covered with various oxide and contaminant films.2–12 In
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Mechanisms of solid-state bonding processes
(a)
(b)
(c)
1.2 Examples of hermetic seals by (a) cold-pressure welding, (b) resistance project and (c) resistance seam welding.
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5
6
Microjoining and nanojoining
addition, there may be microstructural/compositional irregularities, which further complicate the joining process. A typical representation of the prebond surface condition is presented in Fig. 1.3. This surface is generally characterized as three layers: the base material, a layer of mechanically and/ or chemically affected metal, and surface oxides/contaminant films. There are a number of mechanisms which can proceed to form a bond between such surfaces. The most important of these is that asperities on the surfaces must be collapsed in order to form intimate contact between materials. In forge welding processes, creation of this intimate contact is done mechanically; that is, local yield stresses are exceeded on the contacting surfaces, and surface deformation is used to create the contact. For diffusion bonding processes (not covered extensively in this chapter), such surface collapse is done under relatively low forces, and relies on creep and surface diffusional mechanisms to consolidate the surfaces. Once the surfaces have come under intimate contact, bonding can still not initiate until a number of other criteria are met. The most important of these are how oxide and surface contaminant films can be affected to allow intimate contact of the underlying virgin materials. Generally, there are two mechanisms for this. For forge welding processes, contaminant films can be broken up as a result of mechanical action. In addition, it is also possible to break down particularly metal oxides by dissolution into the matrix. This is a mechanism particularly important in diffusion bonding,13 but also plays a role in other thermally assisted forge processes. Even when base materials are in intimate contact, there are additional changes which must occur to facilitate an adequate joint. First, crystallographic matching across the boundary must occur.14 Obviously, most forge welding applications are between parent materials with randomly oriented grain structures, and so this bond surface must take on the characteristic of a series of high-angle grain boundaries. Generation of this dislocation structure can occur mechanically,6 thermally,9,10 or by a combination of the two. At this stage of the process, intimate solid-state bonding has undoubtedly occurred;
Base metal
Cold-worked layer Contaminant layer
Oxidized layer
1.3 Schematic representation of workpiece surface conditions in the pre-bonded state.
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Mechanisms of solid-state bonding processes
7
however, the localized high-angle grain boundary structure is relatively unstable, and is unlikely to yield an adequate joint. As a result, the last stage of the process is to relieve these local bondline stresses, typically with some sort of thermal assist. Depending on the treatment, this local concentration of strain energy can result in a final bond structure ranging from local recovery to recrystallization.15 As mentioned, these mechanisms collectively permit solid-state bonding between metallic materials, although not all mechanisms are used by all solid-state welding processes. Generally, these mechanisms, particularly as they operate within the group of forge welding processes, can be classified in three general areas. These include surface deformation mechanisms, contaminant dissolution mechanisms, interfacial structure homogenization mechanisms. These are described in detail, including their direct relationship to the forge welding processes, in subsequent sections.
1.3
Contaminant displacement/interatomic bonding
As briefly described above, surface deformation mechanisms have two functions: to collapse surface asperities, and displace surface contaminants. It is of interest that the best information on the role of surface strain and its effect on the extent of solid-state bonding is available in the cold-pressure welding literature.2–8 Several authors have examined the roles of surface condition, mechanisms of interfacial breakdown, and degrees of subsequent bonding for cold-pressure welding applications. Collectively, initial bonding, related to surface straining, appears to progress through the following stages.
1.4
Extension of the contacting surfaces
For any bonding to occur, it is essential that contaminating oxides/films be disrupted. This is, of course, accomplished by application of contact surface strains. It is equally important, however, that surface oxides/films be in a condition in which they can be readily broken up when the surface strain is applied. The ideal underlying uncontaminated material. Table 1.1 lists some hardnesses for some common oxides. Of interest here is the difference between the hardnesses of the Al or Cu oxides. Figure 1.4 shows some fractographs of bond surfaces for cold-pressure welds below critical bonding deformation for Al, Cu, and Ag. It is clear from these results that the Al oxide fails in a brittle manner, while the Cu oxide fails in a shear manner. A more general plot showing the relationship between oxide/metal hardness ratio and required deformation for bonding is presented in Fig. 1.5. One method of improving the characteristics of surface film fracture is to locally cold work the base metal. It has been demonstrated that cold working the surface using scratch
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8
Microjoining and nanojoining Table 1.1 Representative hardnesses of some metal oxides at room temperature (Ref. 6, Table 1) Metal
Al Cu Ag Au
Hardness (HV)
Oxide
Hardness (HV)
15 40 26 20
Al2O3 Cu2O Ag2O –
1800 160 135 –
brushing both minimizes the extent of contaminant films, and creates a local layer of heavily cold-worked material, which on straining can fracture and carry more ductile oxide films.6 For most conventional forge welding processes, extension of the contacting surfaces is done with a combination of heat and force. For this stage of the process, local strain is the most important factor. However, how that strain is distributed is a strong function of how the thermal field is applied. Figure 1.6 shows some results from numerical simulations of the flash-butt welding process.16 Flash-butt welding is a variant of resistance butt welding, commonly used for wire attachment. This plot shows how contact surface strain is affected both by the amount of upset used, and the level of flashing acceleration employed. Flashing acceleration directly controls the heat distribution in the workpiece, with higher flashing accelerations indicating steeper higher conductivity materials, particularly Al and Cu, strain location provided simply by the thermal gradient is difficult.17–19 In such cases, pinch-off dies are recommended.20 Pinch-off dies simply use the constraint of the die (rather than the thermal profile) to create localization of forging. The function of pinch-off dies is shown schematically in Fig. 1.7.
1.5
Separation of the contaminated areas
It is established then that the onset of bonding occurs as with applied surface strain as surface contaminants are separated and virgin base materials are allowed to contact. Considerable work has been done, again largely in the cold-pressure welding area, attempting to quantify the separation of these contaminants and the resulting bond quality. In examining the role of contaminated surfaces, all workers agree that first a critical strain must be achieved that this surface ruptures.2–8 There are discrepancies, however, on how this rupture occurs. Mohamed and Washburn6 suggest that separations of the two contaminated surfaces are unrelated, while Wright et al.7 and later Bay8 suggest that surface contaminants impinge on either side of the joining and, therefore, separate as pairs. Local bonding is then accomplished by extrusion of virgin material into the spaces between the separated contaminated
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Mechanisms of solid-state bonding processes (a)
9
(b)
1000×
1000×
(d)
(c)
1000×
1000× (e)
3000×
1.4 Fractographs showing the faying surfaces of cold-pressurewelded Al, Cu, and Ag at sub-bonding strains (a) Al, 3% deformation, (b) Al, 6% deformation, (c) Cu, 5% deformation, (d) Cu, 23% deformation, (e) Ag, 25% deformation (ref. 6, Figure 1).
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10
Microjoining and nanojoining 6 In
Oxide / metal hardness ratio
5
4
Al Al
3 Sn Fe
2 Cu
Cd Pb
Ni Zn Ag
1
0
20
40 60 Deformation, %
80
100
1.5 Relationship between oxide/metal hardness ratio and the critical deformation for bonding during cold-pressure welding.3
surfaces. Each set of authors has developed models based on their assumption of interfacial breakdown. In each case, the underlying assumption is that the strength of the joint is a direct function of the fraction of the bond area converted by actual base material-base material bonds. The simplest of these models is that developed by Mohamed and Washburn.6 This model assumes a completely brittle contamination layer and no coordination of contaminants on either side of the bondline. The physical representation of this model is presented in Fig. 1.8. The resulting equation for strength is:
f=C
( R R+ 1 )
2
(1)
where f is the ratio of the joint strength to the parent material strength, R is the surface strain, and C is a constant to incorporate contaminant mismatches and contaminant hardness. The model proposed by Wright et al. is slightly more complex. This model was generated for roll-bonding applications, so largely plain-strain conditions exist. The physical representation of this model is presented in Fig. 1.9. This model assumes matchup of contaminants across
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Mechanisms of solid-state bonding processes
11
3500 Acceleration = 0.03 in./sec2 Acceleration = 0.003 in./sec2
3000
Percent strain
2500
2000
1500
1000
500
0
0.00
0.05
0.10 0.15 0.20 Upset distance (in.)
0.25
0.30
0.35
1.6 Thermomechanical modeling results showing the relationship between flashing acceleration, upset distance, and contact surface strain for flash-butt welding mild steel.
Die
Workpiece
Die
Workpiece
Die
Die
Die
Die
Workpiece
Die
Workpiece
Die
1.7 The use of pinch-off dies in upset welding processes to localize strain.
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12
Microjoining and nanojoining Metal Oxide Oxide Metal (a) Original interface
(b) Fracture of brittle oxide film Metal
Metal (c)
First requirement for welding: formation of overlapped oxide-free metallic areas
(d) Second requirement for welding: extrusion of the metal through the gaps created in the oxide and some relative shear displacement at the points of contact of oxide-free metal
1.8 Schematic illustration of interfacial breakup as proposed in the Mohamed and Washburn Model.6
Composite surface
b
b
b
c a
b
b
c a
b
b
c a
a
Uncontaminated welded regions of base metal
Composite surface
1.9 Schematic illustration of interfacial breakup as proposed by Wright et al.7
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Mechanisms of solid-state bonding processes
13
the bondline, and attempts to account for a degree of pre-bonding deformation. The resulting equation for joint strength is: (1 – R f ) 2 f = C 1 – (1 – Rt ) 2
(2)
In this case, C is considered an empirical hardening factor, Rt is the threshold deformation for bonding, and Rf is the total deformation of the process. Eqs 1 and 2 are similar, asymptotically approaching a maximum bond strength as the total deformation (R or Rf) approaches 1. The most complex analysis is provided by Bay. This model includes the effects of contaminant films and sub-surface hardened layers, and is shown in Fig. 1.10. The resulting equation for joint strength is: f = (1 – β ) Y
p – pe p + β Y – Y′ σo 1 – Y ′ σo
(3)
where f is now the ratio of weld tensile strength to base material tensile strength, β is the fraction area covered by contaminant films, p is the applied pressure, p is the threshold pressure for bonding, and σo is the yield strength of the base material. Y and Y ′ are the surface exposure and threshold surface exposure, where the surface exposure is defined by: Y=1–
1 1+X
(4)
Cover layer
Contaminant film (a) Onset of extrusion
Local thinning of contaminant film (b)
Welds (c)
1.10 Schematic illustration of interfacial breakup as proposed by Bay.8
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Microjoining and nanojoining
where X is the degree of expansion of the contact area. These models, of course, show a greater degree of complexity, as a greater number of bonding factors are included. It is important to recall, however, that these models have been developed for cold-pressure processes, and these complexities may be more or less relevant for conventional forge welding processes. One factor of note is the extrusion pressure (described as p in Eq. 3 above). This factor is included either directly or indirectly in each of these models. However, for conventional forge welding processes, extrusion pressures will fall dramatically with temperature, and may be less of a factor. Also to be questioned is the role of sub-surface cold-worked layers which, in conventional forge welding processes, will probably anneal substantially before any macroscopic deformation occurs.
1.6
Realignment of the grain structures for bonding
There is considerable evidence that crystallographic matchup across the bondline is also important at this stage of bonding.3,6,14 Detailed work3 suggests that contact between similarly oriented close packed, or near close packed planes most readily bonds. For Al, (111) to (111) and (110) to (110) were found to bond readily, while (111) to (100) were found difficult to bond. However, most structural materials are polycrystalline, so such ideal crystallographic matchups are relatively uncommon. To accomplish bonding requires some localized crystallographic reorientation. The model here is the one of a series of grain boundaries. Grain boundaries can be thought of as a complex dislocation pileup, accommodating the mis-orientation between grains over a very small distance. The types of macroscopic surface strains, and local inter-contaminant extrusion described here are ideal sources for dislocation generation, and undoubtedly contribute to the generation of this bondline structure. An example of this dislocated structure is presented in Fig. 1.11. This particular example is a resistance butt weld on steel, showing evidence of a residual bondline. This region is characterized by relatively high internal strain energy, and may be a quality concern. Reactions of this region to the applied thermal fields typical of the forge welding processes are described in a subsequent section.
1.7
Thermal dissolution of oxides/contaminants
The preceding discussion indicates the degrees to which bondline strain can be used to create a solid-state bond. However, implicit in that discussion and related modeling were two related facts: some level of contamination was always present in the joint area, and joint strengths could only asymptotically approach parent material strengths. The relationship here is straightforward.
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Mechanisms of solid-state bonding processes
15
1.11 Resistance butt weld on mild steel indicating a highly deformed zone down the bondline.
As long as contaminants exist in the joint, they reduce effective bond area, and act as initiation sites for subsequent mechanical failures.21 To achieve improved joint properties, particularly in industrial applications, some further reduction in the residual bondline contaminant content is advantageous. Fortunately, for many metallic systems, oxides are soluble in the matrix at elevated temperatures. The degree of solubility of a specific oxide in MxOy in a matrix of metal A, at equilibrium can be defined by the solubility product: Keq = Z(CM)x(Co)y
(5)
where Keq is the equilibrium solubility product, and CM and CO are the compositions of the metal M and oxygen in the matrix metal A, and Z is a proportionality constant relevant to activity coefficient. If the oxide is of the matrix metal, this expression reduces to: Keq = Z(Co)y
(6)
with Z a different proportionality constant. This suggests that the solubility product is similarly a power function of the maximum soluble oxygen content in the base material. For oxides of the base materials, Eq. 6 suggests that the solubility product can be estimated from the maximum solubility of oxygen in the matrix as taken from the appropriate phase diagram. In addition, the shape of the oxygen solvus provides some indication of the temperature
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Microjoining and nanojoining
dependence for the solubility product. In a similar manner, Eq. 5 suggests that the solubility product for non-base material metal oxides relative to the base metal (as well as the temperature dependence) can also be estimated from the phase diagram. Here, the solubility product is estimated from the maximum solubilities of the secondary metal and oxygen in the base material. If Raoultian behavior of oxygen in the base material is assumed, the proportionality constants in Eqs 5 and 6 become equal to 1. Then, knowing the stoichiometry of the oxide present, and using the appropriate phase diagram, some approximate solubility products for some different metals can be done directly. Table 1.2 lists some of these approximate solubility products for some standard engineering materials with their most common oxide. These solubility products are calculated for approximately bonding temperatures estimated for forge welding processes (0.9 Tm). Materials shown include Al, Fe, and Ti. These solubility products cover about 30 orders of magnitude, indicating, on one extreme, the difficulty dissolving Al oxide into an Al matrix, and the relative ease with which Ti can dissolve its own oxide. Similar calculations can be done for non-matrix metal oxides. The approximately solubility product for Al2O3 on Fe is calculated, and compared to the similar calculation for Fe’s own oxide (Fe2O3) in Table 1.3. These calculations were done at the approximate bonding temperature for forge welding Fe. In this case, the stoichiometry for the two oxides is similar (x = 2, y = 3), so the difference between the two solubility products is directly related to the solubility of the Al in the steel. This fact appears to account for the relatively low solubility of Al2O3 in Fe. Table 1.2 Estimated solubility products for oxides present on some common engineering materials assuming Raoultian behavior of oxygen in solution in the parent material (calculations are done for temperatures roughly representing bonding temperatures for forge welding processes (0.9 Tm) (internally generated).) Metal
Oxide
Keq
Al Fe Ti
Al2O3 Fe2O3 TiO2
3 × 10–29 1 × 10–15 1 × 10–01
Table 1.3 Comparison of solubility products for Al2O3 and Fe2O3 in an Fe matrix at the appropriate bonding temperature for forge welding Fe (0.9 Tm) (internally generated) Base material
Oxide
Keq
Al Fe
Al2O3 Fe2O3
1.8 × 10–19 1.0 × 10–15
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Mechanisms of solid-state bonding processes
17
Such solubility products, and diffusivities of oxygen in the matrix, have been used as a basis for modeling oxide dissolution during diffusion bonding.22 However, this analysis was largely based on continuous oxide films, and focused on the maximum thickness of these films for relatively long (diffusion bonding) heating cycles. Such an analysis does not take into account the breakup of the oxide film into discrete particles caused by the applied surface strain during forge welding processes, or the relatively short heating times. A better analysis can parallel that done by Ashby and Easterling for the dissolution of carbide particles during welding.23 That analysis examines the dissolution of discrete particles. The approach used attempts to estimate the roles of both the solubility of the carbide constituents into the matrix, and diffusion of these constituents away from the decomposing carbide. The approach is based on the assumption that distinct spherical particles can be dissociated completely into a volume of matrix with radius I. Further, that the particle will dissociate into this volume at a locus of times and temperatures defined by: 1
l = ( D*t *) 2
(7)
where D* and t* are the combinations of the diffusivities (D, a function of temperature, defined at T*) and times over which the particle can be completely dissolved into the volume matrix defined by I. Combining this approach for examining the role of diffusion can be combined with an expression for the temperature dependence of the solubility product of the particle, to examine particle dissolution behavior. The discussions on solubility products for oxides detailed above can be used to adapt these equations for oxide particle dissolution. The resulting governing equations include for base metal oxide particles: Ts =
B (O) y A – In f
(8)
and for non-base metal oxides: Ts =
B ( M ) x (O) y A – In f
(9)
where Ts is the dissolution temperature, A and B are the temperature coefficients for the appropriate solubility product, and f is the matrix volume fraction affected by the decaying oxide, defined by: f=
1 3
Q 2 1 + t * exp – 2 1 – 1 R T * Ts t
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(10)
18
Microjoining and nanojoining
In this expression, Q2 is the activation energy associated with the appropriate diffusivity (oxygen or metal + oxygen), and R is the ideal gas constant. With some estimate of t* and T*, Eq. 10 combined with either Eqs 8 or 9 (as appropriate) define an implicit relationship between the time–temperature profile for the process, and the degree of oxide dissolution. In these expressions, t* and T* are direct functions of the oxide particle size and distribution. Values for these can be presumably estimated from the original distribution of oxides on the bonding surfaces, and estimations of surface strain as described above. From these equations, some qualitative estimate can be made of the role of both the degree of forging, and time–temperature profile on bond quality. With increasing strain applied to the contacting surface, particle size and density will both inevitably fall. These factors reduce amounts of oxygen (and potentially second metal), which must be diffused, and increase the kinetics of oxide dissolution. Extended heating (welding) times are important, in that again diffusion is promoted. Increasing welding temperatures are advantageous both for increasing rates of diffusion, but for increasing the solubility product for the dissolution reaction.
1.8
Breakdown of the interfacial structure
A third mechanism of bonding results from the decomposition of any interfacial structure. As described above, straining of the bond surface, extrusion of material around residual oxide particles, and matching crystallographic structures across the bondline result in a highly dislocated bondline structure. This highly dislocated structure is both of relatively high energy, and planar. An example of such a highly dislocated bond is shown in Fig. 1.12. This structure obviously develops during straining the contact surface. However, the presence of various particulates from the contaminated bond surfaces may also stabilize this structure. Decomposition of this structure is largely a thermally assisted process. To develop this highly dislocated structure, considerable energy for deformation is required. Much of this energy is stored in the interfacial structure itself. With varying degrees of activation energy, this structure can quickly decompose to a lower energy variant. Parks15 has done considerable work understanding the breakdown of contacting interfaces. In his work, he suggests two regimes for breakdown of this interfacial structure. These parallel the concepts of recovery and recrystallization. Recovery of the interface implies a realignment of the dislocated structure in order to reduce the overall strain energy of the system. This is typically done at relatively low temperatures, permitting only local movement of the dislocations which make up the boundary structure.24–26 During recovery, these dislocations realign themselves into dislocation cells. An example of such cells is shown in Fig. 1.13.27 Parks
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19
1.12 Interfacial structure on resistance projection-welded mild steel.
1µ
1.13 Dislocation cells in a dynamically recovered Eron microstructure.24
found that very little effective bonding occurred if interfacial decomposition was limited to recovery. Rather, substantial bond strengths were found if higher annealing temperatures were used, resulting in bondline recrystallization. This is shown in Fig. 1.14. Recrystallization is essentially the nucleation and growth of new grains. Provided activation energies are high enough, this mechanism of interfacial decomposition shows the greatest reduction in
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Microjoining and nanojoining
50,000
Silver Copper Aluminium 70–30 brass SAE 1020 steel 75A titanium
Titanium
Shear strength, PSI
40,000 SAE 1020 steel
30,000 70–30 brass 20,000
Copper
Silver 10,000 Aluminium
200
400
600 800 Interface temperature, °F
1000
1.14 Weld strengths as a function of annealing temperature for a range of materials.15
bondline energy, and is suggested by Parks as essential for forming highintegrity bonds. During welding, residual stored energy (as local deformation) can play a role in the kinetics of recovery and/or recrystallization. Parks has demonstrated that actual bonding temperatures can be reduced depending on the degree of deformation in the material. Required recrystallization temperatures as a function of the degree of deformation for a range of materials are shown in Fig. 1.15. Obviously, the extent to which this interfacial structure can decompose is a function of both the amount of strain applied, and the temperature cycle experienced. Increasing amounts of strain (upset) obviously increase the amount of work in the material, and promote subsequent breakdown of the interfacial structure. Time at temperature, however, provides the activation energy to allow this aspect of the bonding process to proceed. It is interesting from this discussion that greater levels of upset might permit bonding at shorter times and lower temperatures. However, in practice extended times
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Mechanisms of solid-state bonding processes –2.0
–1.0
21
0
1.0
Platinum 1200
Nickel
Recrystallization temperature, °F
1000
Iron
70–30 brass
800
Aluminum 600
Magnesium 400 Copper Silver 200
Recrystallization deformation
10
20
30 40 50 60 Percent cold work
70
80
90
1.15 Recrystallization temperatures as a function of degree of deformation for a range of materials.15
and temperatures are almost always advantageous permitting maximum homogenization of the joint microstructure.
1.9
Summary
This chapter has taken a systematic look at mechanisms of bond formation during solid-state (forge) welding processes. Discussions have been limited to those processes which can be characterized as having two stages: heating and forging. Explicitly excluded were those processes which do not use
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heating (cold-pressure welding processes) or forging (diffusion bonding processes). For the forge welding processes, three distinct mechanisms of bonding have been discussed. These include contaminant displacement/ interatomic bonding, dissociation of retained oxides, and decomposition of the interfacial structure.
1.9.1
Contaminant displacement/interatomic bonding
This mechanism of bonding relates to displacement of contaminants by local strain at the contacting surface. Displacement of these contaminants allows exposure of clean surfaces for direct interatomic bonding. The basics for modeling this mechanism were largely taken from the cold-pressure literature. Although several models are available with increasing levels of complexity, all predict that bond strengths asymptotically approach that of the base metal with increasing surface strain. For the forge welding processes, the developed temperature distribution also plays a role, increasing metal plasticity, assisting in localizing strain at the bondline, and reducing required upset loads.
1.9.2
Thermal dissolution of oxides/contaminants
The applied surface strains described above permit considerable bonding, but leave a residue of oxide/contaminant particles dispersed over the bond surface. As a mechanism for further facilitating bonding, many of these particles can be thermally dissolved in the matrix. The relative solubility of specific types of particles can be assessed directly by examining solubility products between the constituent elements of the particle compared with solution in the base material. For base material oxides, this solubility product is only a function of the solubility limit of oxygen in the matrix. This analysis was used to indicate relative stability of a range of base metal oxides. This examination was extended, using previous work done for dissolution of carbide particles in steel, to examine the kinetics of dissolution. This analysis incorporates both solubility product and diffusivity factors. The results indicate the effects of residual oxide particle size, as well as the role of the severity of the thermal cycle for dissolving these oxide particles.
1.9.3
Breakdown of the interfacial structure
The side result of the first two mechanisms is a highly dislocated interfacial bond structure. This structure results largely from the application of bondline strain, but can be stabilized by the presence of discrete oxide particles. This structure is of relatively high energy, and can be a detriment to weld quality. Decomposition of this structure does improve bond quality. The mechanism of decomposition, however, depends on the thermal cycle employed. For
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23
relatively short or low-temperature cycles, the structure may only recover, resulting in a series of dislocation cells. At higher temperatures and longer times, recrystallization of the metal at the bondline can also occur. Some results suggest that recrystallization of the bondline structure results in better weld properties. Increasing stored energy at the bondline (caused by higher levels of strain) also appears to aid the kinetics of recrystallization, and improve weld quality.
1.10
References
1. Standard Welding Terms and Definitions, ANSI/AWS A3.0-94, American Welding Society, Miami, FL (1994). 2. D.R. Miller and G.W. Rowe, “Fundamentals of Solid Phase Welding,” Metallurgical Reviews, 28(7): 433–480 (1962). 3. R.F. Tylecote, “Investigations on Pressure Welding,” British Welding Journal, 1(3): 117–135 (1954). 4. R.F. Tylecote, D. Howd and J.E. Furmidge, “The Influence of Surface Films on the Pressure Welding of Metals,” British Welding Journal, 5(1): 21–38 (1958). 5. L.R. Vaidyanath, M.G. Nicholas and D.R. Milner, “Pressure Welding by Rolling,” British Welding Journal, 6: 13–38 (1959). 6. H.A. Mohamed and J. Washburn, “Mechanisms of Solid State Pressure Welding,” Welding Journal Research Supplement, 54(9): 302s–310s (1975). 7. P.K. Wright, D.A. Snow and C.K. Tay, “Interfacial Conditions and Bond Strength in Cold Pressure Welding by Rolling,” Metals Technology, (1): 24–31 (1978). 8. N. Bay, “Mechanisms Producing Metallic Bonds in Cold Welding,” Welding Journal Research Supplement, 62(5): 137s–142s (1983). 9. K. Inoue and Y. Takashi, “Recent Void Shrinkage Models and their Applicability to Diffusion Bonding,” Materials Science and Technology, 8(11): 953–964 (1992). 10. Y. Takashi, K. Inoue and K. Nishiguchi, “Identifications of Void Shrinkage Mechanisms,” Acta Metallurgica Materials, 41(11): 3077–3084 (1993). 11. T. Enjo, K. Ikeuchi and N. Akikawa, “Effect of Oxide Film on the Early Process of Diffusion Welding,” Transactions of the JWRI, 10(2): 45–53 (1981). 12. T. Enjo, K. Ikeuchi and N. Akikawa, “Effect of the Roughness of the Faying Surface on the Early Process of Diffusion Welding,” Transactions of the JWRI, 11(2): 49–56 (1981). 13. A. Nied, “private communication,” General Electric Company, Research and Development Center, Schenectady, NY (1991). 14. V.M. Zalkin, “Theoretical Problems of Cold Pressure Bonding of Metals,” Svar. Proiz., (11): 41–42 (1982). 15. J.M. Parks, “Recrystallization Welding,” Welding Journal Research Supplement, 32(5): 209s–222s (1953). 16. J.E. Gould and T.V. Stotler, “An Examination of Morphological Development during Flash Butt Welding,” EWI Cooperative Research Report (1995). 17. E.F. Nippes, W.F. Savage, J.J. McCarthy and S.S. Smith, “Temperature Distribution during the Flash Welding of Steel,” Welding Journal Research Supplement, 30(12): 585s–601s (1951).
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18. E.F. Nippes, W.F. Savage, S.S. Smith, J.J. McCarthy and G. Grotke, “Temperature Distribution during the Flash Welding of Steel – Part II, Welding Journal Research Supplement, 32(3): 113s–122s (1953). 19. E.F. Nippes, W.F. Savage, G. Grotke and S.M. Robelotto, “Studies of Upset Variables in the Flash Welding of Steels,” Welding Journal Research Supplement, 36(4): 192s– 216s (1957). 20. Resistance Welding Manual, fourth edition, Resistance Welding Manufacturers Association, Publs., Philadelphia, PA (1989). 21. W.F. Savage, “Flash Welding – Process Variables and Weld Properties,” Welding Journal Research Supplement, 41(3): 109s–119s (1962). 22. Z.A. Munir, “A Theoretical Analysis of the Stability of Surface Oxides during Diffusion Welding of Metals,” Welding Journal Research Supplement, 62(12): 333s–336s (1983). 23. M.F. Ashby and K.E. Easterling, “A First Report on Diagrams for Grain Growth in Welds,” Acta Metallurgica, 30: 1969–1978 (1982). 24. J.D. Embury, A.S. Keh and R.M. Fisher, “Substructural Strengthening of Materials Subject to Large Plastic Strains,” Transactions of the Metallurgical Society of AIME, 236(9): 1252–1260 (1966). 25. J.H. Cairns, J. Clough, M.A.P. Dewey and J. Nutting, “The Structure and Mechanical Properties of Heavily Deformed Copper,” Journal of the Institute of Metals, 99: 93– 97 (1971). 26. A.L. Wingrove, “Some Aspects of Relating Structure to Properties of Heavily Deformed Copper,” Journal of the Institute of Metals, 100: 313–314 (1972). 27. J.E. Pratt, “Dislocation Substructure in Strain Cycled Copper as Influenced by Temperature,” Acta Metallurgica, 15(2): 319–327 (1967).
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2 Mechanisms of soldering and brazing T T A K E M O T O, Osaka University, Japan
2.1
Introduction
Soldering and brazing are important methods of joining materials without melting the base metals. The common necessary characteristic of these processes is called wetting. Excellent wetting offers good joint quality such as mechanical properties and reliability. These two methods can be classified by several criteria, such as filler metals, joining methods, and atmosphere of joining. One major characteristic of these methods is that they can join many parts collectively. Examples are the soldering of printed circuit boards (Klein Wassink, 1984; Hwang, 1996; Rahn, 1993) with high density packaging, and aluminum brazing of heat exchangers for automobile use: in electronic packaging (Judd and Brindley, 1992; Hwang, 1995), the number of joints on one substrate board easily exceeds several thousand points and popular aluminum heat exchangers can contain more than 1,000 braze joints. It is extremely important to obtain high quality joints in fine pitch surface mount technology to secure the reliability of product. In conventional soldering and brazing the penetration characteristic of filler metal by capillary force is important to make sound joints in lap type joints. The wettability plays a great role in the penetration process of molten filler metal and is influenced by many factors. However, the main factors are material and compositional such as filler metal, base metal, and flux, and physical factors such as operation temperature, surface condition and atmosphere. In both methods, surface oxides on base metals and filler materials are obstacles against wetting, therefore the removal of these oxides is an important process. Usually these processes use fluxes to remove surface oxide films for obtaining good wetting. The removal process is an electrochemical process similar to corrosion, in which the flux residue has corrosive characteristics. The removal of flux is also important to secure corrosion resistance of joints. The products using microjoining are widely used in electronics such as mobile phones, personal computers, and popular electric products including 25 WPNL2204
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white goods and brown goods. In these products so called microsoldering is the key technology for assembly of high density fine pitch wiring boards. A variety of heat exchangers for automobile use are representative products using microjoining technology that brazes thin and fine pitch fins to thin tube or plate. Deviation from the correct brazing conditions introduces severe erosion damage at high temperature conditions and defects due to poor wetting and reduced fillet size at low temperature conditions offers low endurance pressure of heat exchangers. In order to reduce the weight the reduction in thickness of brazing sheet fins is ongoing in the automobile industry. Even though the fillet size is also being reduced to enhance the efficiency and compactness of heat exchangers, the joints should have high strength to withstand high pressurization. Using microjoining, microsoldering and microbrazing in the production of these products, care for the reaction between molten filler and base metal becomes more important to make a sound joint. Any excess reaction enables thick intermetallic compounds (IMCs) to form; undesirable deformation by erosion that reduces the quality of joint properties.
2.2
Definition of soldering and brazing
There are several important characteristics of soldering and brazing. The most important being that the process can be conducted under the melting point of base metals. Melting of base metal is essential in various welding processes; however, soldering and brazing offers relatively reliable joints without melting of the base metal. The definition of filler metals used in soldering and brazing are listed in Table 2.1. The boundary between soldering and brazing is divided by the liquidus temperature of filler materials. In brazing, the filler metals having liquidus temperature exceeding 450°C are called brazing filler metals or fillers (Schwartz, 2003; Humpston and Jacobson, 1993) and in soldering, the filler material having liquidus temperature less than 450°C is called solder. The practical brazing temperature varies from 600 °C for aluminum brazing to 1200 °C for nickel brazing (Humpston and Jacobson, 1993; 2004; Jacobson and Humpston, 2005). Many commercially used brazing processes are not included in the microjoining process; the exception is the production of heat exchangers with fine fin/plate structure. On the other hand, soldering is the major microjoining process in the assembly of printed circuit boards for advanced high-performance electronics. Furthermore, soldering has become a key technology in the assembly of high density surface mounted electronic boards with fine pitch. Figure 2.1 shows the lowest solidus and the highest liquidus temperatures in the same filler metal groups; the plots are based on JIS. Aluminum brazing in the manufacture of various heat exchangers for automobile use may be the
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Table 2.1 Definition and distinction of filler metals dividing solders from brazes Soldering
Brazing
Name and definition of filler material
Solder Liquidus less than 450 °C Sn-based solders usually form IMC Brittleness due to IMC
Filler, filler metal, filler alloy Liquidus more than 450 °C Usually forms solid solution Erosion due to dissolution of base metal into molten fillers High High High oxidation rate High reducing effect in reduction atmosphere such as H2 gas Seldom exceeds 0.5 Tm of fillers Heat exchangers, pipes, various equipment, jewelry
Reaction between filler and base metal Characteristic of interface Activity of flux Corrosiveness of flux residue Oxidation and reduction during operation Service temperature Major application
Low Low Low oxidation rate Low reducing action even in H2 atmosphere Often exceeds 0.5Tm of solders Electric and electronics
Tm: Melting temperature
Mechanisms of soldering and brazing
Definition and distinction
27
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Microjoining and nanojoining 1350 Maximum liquidus Minimum solidus
1250
1050
B-Ni
B-Cu-P
550
B-Ag
650
B-Pd
750
B-Au
VB-Au
850
B-Cu B-Cu-Sn B-Cu-Zn
950
B-Al
Temperature (°C)
1150
450
2.1 The lowest solidus and the highest liquidus temperatures in the same filler metal groups listed in JIS.
largest brazing market in yield of products, variety of brazing methods and types of heat exchangers. In this brazing process, the filler metals are supplied as clad material with core base material. The clad sheet is called the brazing sheet. Reduction in weight is important in this type of engineering, therefore, the thickness of fins and plates is constantly being reduced. The combination of filler metals for several types of heat exchanger are given in Table 2.2. The composition of filler metal in each heat exchanger is selected by considering the brazing process, and required reliability including mechanical and chemical properties. Among the various filler metal systems, the microbrazing is used in jewelry manufacture. Several carat gold alloys are brazed using similar colored brazes, they are mainly gold based Au-Ag(-Cu-Zn) with a melting temperature range around 415 ~ 960°C. Figure 2.2 shows similar plots as Fig. 2.1 for solder systems. The main solder systems can be classified according to composition; Zn, Cd, Pb and Sn based solders being most popular. Zn-Al and Cd-Zn are mainly used for the joining aluminum and Al alloys. The Sn-Pb eutectic system has been the conventional solder for electric and electronic assembly; however, recently Sn-based system solders have become more important for practical use, because the ban due to Restrictions on Certain Hazardous Substances (RoHS) related to Waste Electronic and Electronic Equipments (WEEE) has forced the use of lead-free solders (Ganesan and Pecht, 2006) in the assembly of electric and electronic equipment with only a few exceptions. Some important lead-free solders are listed in Table 2.3 together with the definition of each system. Sn-based lead-free solders can be classified according to the melting
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Table 2.2 Filler metals used for production of heat exchangers and main brazing methods Main filler metal system
Main base materials
Main brazing methods
Automobile (Radiator, evaporator, condenser, oil cooler, etc.)
Al-Si Al-Si-Mg
Furnace, non-corrosive flux, N2 gas Vacuum
Instantaneous water heater Inter cooler Chemical plants High pressure fluids, ammonium tiller
Cu-P, Cu-Ag-P Cu-Ni-Sn-P Ni-Cr-Si-P Ni-Cr-Si-P
Sea water, ships, medicine production line, ultrapure water production, chemical engineering, medical equipments
Ti-based, T-Cu-Ni
Aluminum alloys (A 3003, A 1100) Aluminum alloys (A 3003, A 1100) Copper Copper, brass Ni-based super alloys Stainless steels (Type 316, 304, 321) Pure titanium, titanium alloys
Fluxless Fluxless, atmosphere Vacuum Vacuum Vacuum
Mechanisms of soldering and brazing
Application fields of heat exchanger
29
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Microjoining and nanojoining 450 Maximum liquidus Minimum solidus
400
50
Cd-Zn Sn-Pb
Sn-Pb-Ag
Sn-Sb
Sn-Zn
Sn-Cu
Sn-Cu-Ag
Sn-Ag-Cu Sn-Ag
Sn-Bi
100
Sn-ln
150
Sn-Zn-Bi
200
Sn-ln-Ag-Bi Sn-Ag-Bi-Cu
250
Zn-Al
300
Sn-Pb-Bi
Temperature (°C)
350
0
2.2 The lowest solidus and the highest liquidus temperatures in the same solder groups listed in JIS.
Table 2.3 Representative lead-free solders adopted in standards Solder
Melting temperature range (°C)
System and definition
Composition, mass (%)
Solidus
Liquidus
High temperature solidus more than 217 °C liquidus more than 225 °C
Sn-5Sb
238
241
Sn-0.7Cu
227
227
Sn-0.7Cu-0.3Ag
217
226
Sn-3.5Ag
221
221
Sn-3.0Ag-0.5Cu
217
219
Sn-3.5Ag-0.7Cu Sn-3.8Ag-0.7Cu
217 217
217 217
Sn-2.5Ag-1Bi-0.5Cu Sn-3.5Ag-4In-0.5Bi
213 207
218 212
Sn-3.5Ag-8In-0.5Bi
196
206
Sn-9Zn
199
199
Sn-8Zn-3Bi
190
196
Sn-58Bi, Sn-57Bi-1Ag Sn-52In
138 119
138 119
Mid high temperature solidus more than 217 °C liquidus less than 225 °C Mid temperature solidus less than 217 °C and more than 150 °C liquidus more than 200 °C Mid low temperature solidus more than 150 °C liquidus less than 200 °C Low temperature solidus less than 150 °C
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31
temperature range. They are divided into the five systems shown in Table 2.3. The precise composition of fillers and operation conditions are explained in Ganesan and Pecht (2006). Figure 2.3 shows the important factors influencing the quality of joints made by brazing and soldering methods. The main three factors are materials, equipment and atmosphere, and process. It is well known that the quality of filler metals and fluxes have great influence on joint characteristics. The composition and microstructure of base metal also play an important role during the brazing and soldering processes by influencing the reaction characteristics including wetting, dissolution and reaction with molten filler metal.
2.3
Basic metallurgical reaction during soldering and brazing
2.3.1
Wetting
The important phenomena that occurs during brazing and soldering processes using flux is schematically demonstrated in Fig. 2.4 together with the temperature rise profile of the joint. The first process is the activation of flux corresponding to the rise in temperature during the preheat stage. The removal
Material Filler material
Flux
Composition Microstructure Surface oxide Level of impurity
Composition Activity Corrosiveness Ease of removal
Base metal
Soldering Solder land on printed wiring board, plating, OSP Electrode of components, plating Brazing Microstructure of base metal
Equipment and atmosphere Source of heating Atmosphere Length of heating zone
Process Temperature profile Preheating Surface treatment
2.3 Important factors influencing the quality of joints made by brazing and soldering.
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Microjoining and nanojoining
4 Initiation of wetting 5 Dissolution of base metal into molten filler 6 Formation of reaction layer, IMC
fil m
3 Melting of filler
lo fo xi de Re
tiv
2
Ac
Preheat
m
at io
ov a
n
of
flu
x
Temperature
Melting point of filler
1
Temperature profile
Time
2.4 Example of temperature profile and phenomena occurring during soldering and brazing processes. γlf
γsf
Filler metal
θ
Flux
γsl Base metal
2.5 Wetting of filler metal on base metal, contact angle θ is an indicator of wettability.
of oxide on filler and base metal subsequently proceeds. This process is important to enhance wetting just after the melting of filler metal. The oxide removal process is the key for promoting wetting. The metallurgical reaction of molten filler and solid base metal initiate after wetting; the first step is dissolution of base metal into molten filler metal. The final process is the formation of the reaction layer. The formation of intermetallic compounds depends on the combination of filler metal and base metal. The conventional explanation of wetting is indicated in Fig. 2.5. The balance of several indices is indicated by the following Young’s equation. γsf = γlf cos θ + γls
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Mechanisms of soldering and brazing
33
where γsf is the interfacial tension between flux and solid base metal, γlf is the interfacial tension between molten filler and flux, γls is the interfacial tension between molten filler and solid base metal, and θ is a contact angle. A small wetting angle is judged to be superior wetting. To reduce the wetting angle, the reduction of γlf is effective, accordingly, the use of active flux is useful to reduce γlf. The reduction of γsl also contributes to enhancing the wettability. To observe the wetting angle, the simplest method is the conventional spread test using a fixed amounts of solder and flux. The other important element in brazing and soldering is the formation of the fillet. The fillet size influences the joint quality especially the mechanical properties and endurance against the thermal, mechanical and chemical environment. Fillet size is restricted by the parameter surface tension, density of filler and gravity, 2 γ / ρg , where, γ is the surface tension, ρ is the density and g is gravity. To secure a large fillet, the reduction of density and increment of surface tension is effective if the quality of wetting is secured. The maximum fillet height can be calculated using the infinite length of vertical plate (Klein Wassink, 1984). The profile of the fillet is indicated in Fig. 2.6 for several filler metals (Miyazaki et al., 1997). The profile is shown as the shape of molten filler wetted on a vertically immersed plate of infinite length, with perfect wetting and a contact angle of 0 degrees. The values used for calculations are shown in Table 2.4. It is clear that the largest fillet can be obtained from aluminium filler; however, smaller fillet formation in soldering is evident under the same 10 Plate θi = 0 rad
8
Height (mm)
B-Al 6
B-Cu-Zn B-Ag
4
B-Au Solder
2
0 0
2
4
6 8 Distance (mm)
10
12
2.6 Meniscus profile of various filler metals and Sn-Pb solder formed on the vertical plate with infinite length under perfect wetting, wetting angle, q, is zero.
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Table 2.4 Physical properties of brazing filler metals and Sn-40Pb solder Filler
B-Al
B-Cu-Zn
B-Ag
B-Au
Solder
Composition Temperature (K) γ (N/m) ρ (kg/m3)
Al-10Si 873 0.89 2400 8.6
Cu-40Zn 1203 0.97 7200 5.3
Ag-28Cu 1123 1.0 9000 4.7
Au-20Cu 1273 1.2 14000 4.2
Sn-40Pb 503 0.38 8100 3.1
2 γ /ρg (mm)
B-Al; γ = 0.89 N/m, ρ = 2400 kg/m3 5 Rod 4 Radius; 1 mm Height (mm)
Surface of rod 3 θi = 0 rad 2
θi = π/6 rad θi = π/3 rad
1
0 0
1
2
3 4 Distance (mm)
5
6
2.7 Meniscus profile of aluminum filler metal formed on the vertical pipe with radius 1 under several wetting angles θ.
wetting condition. Of course the figure only shows the ideal state assuming perfect wetting. In fact the practical fillet size is strongly dependent on wettability; poor wetting reduces fillet size. Figure 2.7 (Miyazaki et al., 1997) shows the reduction in fillet size due to the decrease of wettability, that corresponds to the increase of the wetting angle. To enhance the wettability, the reduction of γsl is effective. The use of active flux can usually reduce γsl. In microjoining the size of works to be brazed is small and they usually have a convex plane. Fillet formation on a convex surface is difficult due to the reduction of the fillet size. Figure 2.8 (Miyazaki et al., 1997) shows the theoretical fillet height of the effect of radius of the vertical rod on the fillet height of three fillers. It is clear that the reduction of rod radius decreases the fillet height. Figure 2.8 illustrates that the reduction of work size with a
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Mechanisms of soldering and brazing
35
10 Rod
Height (mm)
8
B-Al
θi = 0 rad
6
B-Ag
4
2
0 0
Solder
2
4
6 8 Radius (mm)
10
∞
2.8 Theoretical fillet height of each filler showing the effect of radius of vertical rod on fillet height, wetting angle is assumed to be zero.
small convex surface creates difficulty in making large a fillet. This fact explains one of the difficulties in microjoining.
2.3.2
Dissolution
Kinetics In microjoining, the phenomena occurring after wetting greatly influences the quality of joints. After wetting, a metallurgical reaction occurs between the molten liquid and the solid base metal. This reaction process includes the dissolution process of the solid base metal into molten liquid and the formation process of a reaction layer between them. The reaction continues during the joining process until solidification of the solder. The dissolution phenomena becomes especially noticeable under the same alloy system of filler and base metals, for example, aluminum brazing using Al-Si and copper brazing using Cu-P. The dissolution rate of base metal in molten filler can be explained by the following equation (Moelywyn-Hughes, 1947; Shoji et al., 1980; 1982). dC/dt = K(A/V)(Cs – C)
(1)
where, C is the concentration of solute in liquid after reaction time t, K is a constant, A is the interface area between the solid and liquid, V is the volume of liquid, and Cs is the saturation concentration of solute in liquid. The parameters A and V are usually constants under the practical brazing process and the dissolution rate is proportional to (Cs – C). The concentration
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Microjoining and nanojoining
difference (Cs – C) is the driving force for dissolution. The saturation concentration of solute Cs increases with a rise of temperature, therefore, a rise of operation temperature also increases the dissolution of base metal. The integral form of Eq. (2.1) becomes: C = Cs[1 – exp(–KAt/V)]
(2)
Equation (2.2) indicates that the dissolution rate becomes smaller with operation time because the dissolved base metal increases C, which decreases the dissolution rate. In microsoldering the fillet size is reduced. In the small fillet, the composition of filler easily reaches saturation concentration. Severe dissolution and subsequent formation of the reaction layer is called erosion. Excess erosion deteriorates joint quality and in extreme cases the molten filler completely penetrates through thin base metal. In such cases the joint and fillet cannot hold the appropriate shape. The selection of filler metal and operation temperature is important to avoid severe erosion. Effect of some parameters The factors influencing erosion except for the parameters shown in the above equations are the composition of filler and base metal, and microstructure of base metal. Figure 2.9 (Okamoto et al., 1983) shows the effect of the content
300
Erosion depth (µ)
Ni 200
Fe 100 Zr Si
0
0
0.5 1.0 1.5 Content of additional elements (wt%)
2.0
2.9 Effect of additional elements to pure aluminum on erosion depth after spread test at 615 °C for 3 min.
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Mechanisms of soldering and brazing
37
of elements in aluminium base metals on erosion depth after a spread test at 615 °C for 3 min. The increase in the amount of additional elements enhances erosion. Figure 2.10 (Okamoto et al., 1983) shows the plots between erosion depth and grain size of base metal. It is clear that the reduction of grain size increases erosion depth in each binary alloy system. The effect of grain size is clearest in the Al-Ni system. In this system, where no intermetallic compound forms, large grain size and low erosion were observed. Figure 2.11 (Takemoto et al., 1989) shows the effect of P content of Cu15Sn-P filler on erosion of oxygen free high conductivity copper, tested at 700°C for 30s using a spread test specimen. The increase of P content increases erosion. Usually the erosion process proceeds linearly with the
Pure Al Al-Si Al-Zr Al-Fe Al-Ni
Grain size (µ)
1500
1000 200
100
0
100 200 Erosion depth (d(µ))
300
2.10 Plots between grain size of aluminum base metal and erosion depth after spread test at 615 °C for 3 min.
Erosion depth (mm)
0.4
Cu-Sn-5P 700°C, 30 s
0.3
0.2
0.1
0
10
15 Sn content (mass%)
20
2.11 Effect of phosphorus content in Cu-15Sn-P filler metals on erosion depth of pure Cu after spread test at 700 °C for 30s.
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Microjoining and nanojoining
square root of holding time. If the volume of molten filler is extremely large when compared to the contacting surface of base metal, erosion proceeds linearly with holding time. Figure 2.12 (Takemoto et al., 1989) is an example of the effect of holding time on erosion depth for a combination of Cu-15Sn7P filler and pure Cu base metal. In this case, using a spread test specimen, erosion depth increases linearly with the square root of holding time. Plots between the inverse of test temperature and the logarithm of erosion rate, Arrhenius plot, showed a linear relationship. The apparent activation energy obtained was 195 kJ/mol; this value is close to the activation energy of diffusion of Sn in solid Cu (188 kJ/mol). Intermetallic compound formation Brazed and soldered joints should have excellent mechanical properties because the joints are often used as strength members. To maintain high strength with ductility, it is better to avoid the combination of filler metals and base metals that form intermetallic compounds (IMCs) between them. This is especially the case with braze joints that are often used under a higher mechanical load than soldered joints; therefore, the formation of IMCs with a hard and brittle nature should be avoided. The practical selection of filler and base metals usually considers these criteria. In soldering, the use of tin-based lead-free solder has become a popular process, but the formation of Sn-containing IMCs is inevitable. The growth of IMCs deteriorates elements of the joint quality such as strength, thermal
Cu-15Sn–7P
740 °C
Erosion depth (mm)
0.6
720 °C 0.5 700 °C
0.4
0.3
0
5
10 Holding time (s1/2)
15
20
2.12 Relation between erosion depth of pure Cu and square root of holding time tested at 700–740°C using Cu-15Sn-7P filler metal.
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fatigue and impact strength. Another important issue in electronic packaging is the growth of IMCs during the practical use of electronics. Even room temperature is a relatively high homologous temperature for solders and the growth of IMCs in a soldered joint is undesirable. Ni-P plating on Cu conductors is effective in suppressing the IMC growth rate, but still the formation of SnNi system IMCs and the growth of IMCs proceeds. During the growth of SnNi IMCs, the formation of and growth of P-rich layer continues. This layer also decreases the solder joint quality; therefore, it is important to avoid the formation of a P-rich layer. In the soldering of a Cu base metal with pure Sn, the equilibrium phase after saturation of Sn with Cu is Cu6Sn5 according to the Sn-Cu binary phase diagram. In this case, the formation of IMCs occurs at the solder/base metal interface. The formation of such reaction products is frequently observed in soldering, because Sn makes a variety of IMCs with many elements (Manko, 1964). The thick reaction layer, especially the IMS layer certainly reduces the joint quality by reducing ductility.
2.4
Materials for soldering and brazing
2.4.1
Solders
Several solder systems are classified in Fig. 2.2 according to the compositional systems. Except for tin-based solders, the amount of practical use is extremely small. In microjoining, the most important solders are lead-free solders. Since July 1, 2006, electric and electronic systems must be assembled using toxic free substances within the restrictions of RoHS related to WEEE in the EU directive. Accordingly, the conventionally used Sn-Pb eutectic system solders cannot be used in electronic packaging for the EU. The main leadfree solders in ISO 9453 (2006) are given in Table 2.3. The listed style is based on JIS Z 3282 (2006). The solder groups are divided according to melting temperature range. Figure 2.2 also showed several Sn-Pb and SnPb-Ag solders. Only Pb-based solders for die bonding can be used in electronic packaging for the EU region. The most important lead-free solders are SnAg-Cu system solders such as Sn-3.0Ag-0.5Cu in Japan and Sn-3.8Ag-0.8Cu in USA and EU. The solidus temperature is 217°C (the ternary eutectic temperature of the Sn-Ag-CU system), which is 34°C higher than the Sn-Pb eutectic temperature. The main characteristics of lead-free solders compared to Sn-Pb eutectic solders are summarized as follows: higher melting temperature ranges, relatively poor wettability and higher dissolution rate of solid metals and alloys. Pb-based solders having a soft and ductile nature are used for die bonding. The replacement of Pb-based solder is preferable, however, no suitable alloys have been developed.
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Microjoining and nanojoining
Zinc-based and Cd-Zn solders are used for the joining of aluminium. Fluxless soldering is available for aluminum using ultrasonic application and zinc-based solders. Zinc-based solder shows severe erosion against aluminum, therefore furnace soldering of aluminium using Zn-based solders is not recommended. To avoid severe erosion, a short operation time is important in practical application. The application of ultrasonic power meets this criteria because it breaks down the surface oxide film that leads wetting.
2.4.2
Soldering fluxes
Table 2.5 shows the classification of soldering fluxes based on JIS Z 3284 (2005). A similar classification is also adopted in ISO 9454–1 (1990). The composition of flux cannot be determined from any of the standards. However, the basic constituents are explained in some books. The main materials are rosin-based carbonic acid, resin and activators (Hwang, 1989; Klein Wassink,1984). However, activators have corrosive properties, and therefore the removal of flux residue is essential to prevent corrosion problems. Recent improved soldering fluxes have less residue corrosiveness. They are called no clean flux and can be used without removal or cleaning of flux residue. Owing to the poor wettability of lead-free solders, relatively active fluxes tend to be used for lead-free soldering. Accordingly, the removal of flux residue after soldering by cleaning or deactivation is essential. After the Table 2.5 An example of classification of soldering fluxes according to their main ingredients, partially based on ISO 9454-1 (1990) and JIS Z 3284 (2005) Constituents Flux type ———————————————————————————— Flux basis Flux activation Fluoride Resin
Colophony (rosin) Colophony modified Non-colophony (resin)
Water soluble substance based Organic solvent based (non-water soluble) Inorganic Water soluble Non-water soluble
No activator added Halide activated Amine halide Organic acid, amine organic salt
Added No addition
Flux form*
Liquid Solid Paste
Organic
With ammonium halide With halogenated zinc With halogenated tin With phosphoric acid With halogenated hydrogen acid
* Flux cored solder wire and pre-flux coating are supply methods of flux.
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banning of the use of Freon in the cleaning process for flux residue, the use of no clean flux has become popular. A variety of deactivating fluxes can be used after the soldering operation. Precise information on fluxes can be found in books by Hwang (1989) and Klein Wassink (1984).
2.4.3
Filler metals
A variety of brazing filler metals have been used for joining various materials with complicated shapes. Wettability is also the most important characteristic in brazing. The combination of flux, filler metals, atmosphere and process conditions are important to achieve excellent wetting and high joints quality. Microjoining is conventionally used in the brazing process in the production of heat exchangers. The main use of several filler metals for heat exchangers have already been given in Table 2.2. Homogeneous heating is essential to obtain high pressure endurance of heat exchangers; poor heating results in brazing defects such as non-wetting. Excess heating offers inadequate microstructure at the filler metal/base metal interface based on the excess reaction between molten filler metal and solid base metal. Homogeneous heating is especially important in the production of large-scale heat exchangers. Avoiding deformation during the brazing stage including the heating and cooling processes is extremely important when brazing heat exchangers. The fixing technology and temperature profile during brazing is important to make sound products. Precise compositions and their uses can be obtained in books by Schwartz (2003), Jacobson and Humpson (2005) and in the standards related to filler metals specified by the American Welding Society. The filler metals used for heat exchangers and for ornaments are important in microbrazing. To avoid deformation, erosion, discoloration, etc., the brazing temperature should be kept as low as possible in order to make a sound fillet. The microstructure of base metal has great influence on the fillet size and erosion, especially in aluminium brazing using a brazing sheets. A brazing sheet is a clad material consisting of a base metal core and a filler metal clad. After melting the clad filler melting the liquid filler fully reacts with the core material. The liquid filler metal diffuses or penetrates rapidly into the fine-grained core material through the grain boundary; therefore in order to suppress the liquid/solid reaction, it is necessary for the core material to have large stable grains. Such core material offers a sound large fillet. The penetration or erosion of core material reduces the fillet size.
2.4.4
Brazing fluxes
Brazing flux plays a great role in removing the oxide film on both base metal and filler metal, and this reaction enhances wetting. Common brazing fluxes contain halogen compounds such as chlorides and fluorides. These chemicals WPNL2204
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have strong corrosion properties against metals and alloys, and therefore their removal by rinsing is essential to secure the reliability of joints. In aluminum brazing the use of non-corrosive AlF-KF system fluxes is popular. A thin flux layer is retained after brazing, therefore, the process requires no subsequent cleaning. The use of protective nitrogen gas is necessary to secure excellent joint formation. The popular fluxes for aluminum brazing are listed in Table 2.6. Currently, one of the problems is that there is no standard indicating the precise composition of these fluxes. A precise indication of flux constituents is important for awareness of several important matters related to the risk of corrosion, health and safety, and environmental issues. The introduction of appropriate flux standards containing a precise indication of content is an important future requirement.
2.5
Soldering and brazing process
2.5.1
Classification of method
The joining processes can be classified according to the following factors: heat source, equipment, process and atmosphere. Aluminum brazing methods are classified in Fig. 2.13. The first step is to divide them according to flux. Many processes are adopted in the production of a variety of heat exchangers for automobile use. The most recent major method is non-corrosive flux using AlF-KF system fluxes and there are several compositions of this type. Also CsF-AlF system fluxes are used for low-temperature brazing; however, the system fluxes are not practical due to their high cost. The other major method is vacuum brazing which produces drawn cup-type evaporators using Al-Si-Mg filler metal. The addition of Mg makes fluxless brazing possible through its better action in removing oxygen from the vacuum atmosphere during Mg evaporation while heating. The following equations illustrate the best reactions of Mg (Terrill et al., 1971). Mg + 1/2O2 → MgO
(3)
Mg + H2O → MgO + H2
(4)
Mg + 1/3Al2O3 → MgO + 2/3Al
(5)
Equations (2.3) and (2.4) are the best reactions to remove oxygen and moisture from the brazing atmosphere and enhancing brazability in the fluxless process. Eq. (2.5) is a direct reduction of surface oxide. The important issue is the evaporation of Mg. If the MgO reaction product remains and covers the surface of base metal, brazability is lowered, therefore, vacuum atmosphere is necessary for the self-removal of reaction products by evaporation.
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Table 2.6 Classification and typical constituents of aluminum brazing fluxes Recommended brazing method
Composition
Melting temperature Range (°C)
Solvent of flux
Concentration (%)
Activity level
Corrosive
Torch Furnace Dip Dry air
NaCl, KCl, LiCl, ZnCl2, LiF
490–600
Water or alcohol
50–75
Medium
480–560 550–570
– Water
100 5–15
High Low
Reaction flux
NaCl, KCl, LiCl, AlF3 NaCl, KCl, LiCl, SrCl2, Na3AlF6, KF ZnCl2
Around 380
Alcohol, ketone
Dry nitrogen gas (Nocolok)
KAlF4, K3AlF6, K2AlF5·H2O
562–580
Water or alcohol
3–10
High
Non corrosive
High
Mechanisms of soldering and brazing
Classification of flux
43
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Flux process
Torch brazing Dip brazing Furnace brazing Controlled atmosphere brazing
Alcoa 393 process Alcan Nocolok process (Non-corrosive flux process)
Particular brazing
Automated flame process Induction heating process Block heating process Molten filler dip process
Fluxless process
Vacuum brazing
High vacuum level process Low vacuum carrier gas process
Inert gas atmosphere brazing
VAW process Borg Warner process KD 206 process
2.13 Classification of aluminum brazing method.
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Brazing of aluminum
Mechanisms of soldering and brazing
2.5.2
45
Soldering for electronics assembly
Soldering is an important process in electronics packaging. Table 2.7 shows the classification of soldering processes, types of solder and corresponding heat sources and equipment. Soldering processes can be classified into three groups: flow, reflow and manual. Flow soldering using a molten solder bath, contains the following three methods: wave, drag and dip. However, wave is the latest major method of high density fine pitch soldering. Recently, some soldering processes, that previously used soldering irons have introduced soldering robots. Figure 2.14 shows a classification of soldering processes together with a simple illustration of the representative processes. In recent lead-free soldering, high quality reliable joint formation is important to secure the reliability of products (Shangguan, 2005). Wave soldering Wave soldering is the main flow-type soldering process, however, the process is better known as flow soldering in Japan. The process joins axial lead components and SOP on the under side of through-hole type boards by contact with a molten solder wave. To enhance wettability, several modifications of wave related to shape, distance between waves, height, etc., are characteristics of equipment designed for lead-free soldering.
Table 2.7 Classification of soldering process, type of solder and corresponding heat sources and equipment Soldering process
Solder material
Heat source, equipment
Flow
Wave Drag Dip
Bar Ingot Wire (for liquid height adjustment)
Solder material
Reflow
Infra-red Hot air Laser
Solder paste Preform Wire Flux cored Plating
Solder material
Manual (robot)
Soldering iron
Flux cored Solder paste Preform Wire Plating
Solder material
Note: the use of N2 gas is possible in all soldering processes.
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Microjoining and nanojoining Soldering process
Manual
Flow
Wave
Drag
Reflow
Dip
Infra-red
Hot air
Laser
(a) Soldering iron
Flux Component cored solder Board wire
Heater
Fan
Hot air
Solder paste
Molten solder Lead
Board
Manual Through-hole type assembly assembly/Repair
Solder wave Wave Through-hole assembly
Board
Component
Hot air reflow Surface mount assembly
(b)
2.14 Classification of main soldering process (a) and simple illustration of representative processes (b).
Reflow soldering This process is important for production of high-density fine-pitch packaging by using solder pastes. The process includes printing of solder paste, setting of components, preheating, soldering, inspection and cleaning of flux residue. The final process can be skipped by using a no clean flux. The most important material in this process is solder paste which contains fine solder powder, flux and chemicals to adjust various rheological properties such as printability, slump and tackiness. The appropriate supply of solder paste is the key to achieve defect-free solder joints. For lead-free soldering the most popular solder paste contains powders less than 25 µm in diameter to adapt fine pitch soldering. Manual soldering Manual soldering is the traditional process used to join component leads to solder lands using a soldering iron and flux-cored solder wire. To prevent solder defects such as glazing and delamination of boards skillful, rapid operation is important. To meet this criteria, a rapid temperature responding manual soldering iron is necessary for lead-free soldering. However, owing to the high erosion rate of lead-free solder, the iron plating on the soldering
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iron tip is easily damaged. In addition, some fluxes in flux-cored solder wire react with iron plating, and as a result offer poor wetting of the iron tip with solder; this phenomena causes trouble in the soldering operation. Of course the correct selection of appropriate input power and the shape of tip are also important.
2.5.3
Brazing
Flux brazing Brazing processes are often classified according to the equipment or heat sources as indicated in Fig. 2.15. The most commonly used commercial brazing equipment are torch (TB), furnace (FB) and induction (IB). The following processes are also used: resistance (RB), dip (DiB), infra-red (IB) and diffusion (DB). In the DiB process, the specimens are immersed in molten flux; however, the dip soldering process uses molten solder – the under layer of the printed circuit board is soldered by touching with a molten solder bath. The molten material is different in both processes. Fluxless brazing To suppress the growth of oxide film on filler metal and base metal, a low oxygen and low humidity atmosphere such as a vacuum and nitrogen gas is used. The major merit of this process is that it requires no flux residue removal process. The usual flux removal process requires a huge amount of water, therefore, the fluxless process may be recognized as an environmentallyfriendly process. Examples of the fluxless processes were previously indicated in Fig. 2.10. In fluxless processes it is important to ensure an excellent atmosphere with low oxygen concentration and a low dew point.
Heat source, equipment, process Laser Induction Ultrasonic Gas flame
Furnace Continuous Torch Batch
Atmosphere Vacuum Carrier gas Atmosphere (nitrogen gas) Decomposed ammonia gas
Flux Fluxless
Type of filler Filler complex • Brazing sheet (filler + core) • Brazing paste (filler + flux) Filler only • Wire • Sheet • Powder • Preform • Plating
2.15 Classification of brazing methods and type of filler metals.
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Table 2.8 Systematic solution of microjoining technology for future development Classification
Future subject
Present requirement
Future subject
Lead-free solder (LFS) with melting temperature range near 183 °C (Sn-Pb eutectic) and 300 °C (Pb-based) LFS having excellent wettability Plating having high long time reliability after aging
Completely lead-free solder for aluminum with excellent corrosion resistance
Reduction of erosion Environmentally conscious (Cd free)
Shift to soldering of Al after development of new solder for Al
Plating no IMC formation with Sn
Low erosion
Prevention of diffusion of Si into core material in Al brazing sheet
Environmentally conscious (toxic free)
Complete no clean process (complete flux removal free)
Fluxless soldering
Flux removal free process Reduction of brazing flux consumption
Target
Self-solidification process similar to diffusion brazing Homogeneous heating in reflow process
Adoption of new heat source Non-contact process instead of soldering iron
Homogeneous pressurizing jig
Brazing under low vacuum level Low oxygen and humidity atmosphere of low cost Low temperature process below recrystallization temperature of base metal
Homogeneous and rapid heating equipment Feed back system during soldering
Development of new heat source
Homogeneous and rapid heating furnace
Development of safety manual process
Decrease of reflection
Filler metal
Base metal Surface treatment Process
Microbrazing (Al brazing)
Present requirement
Equipment Heat source and furnace New heat source (Laser)
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Microjoining and nanojoining
Material
Microsoldering
Mechanisms of soldering and brazing
2.6
49
Summary and future trends
Soldering and brazing in microjoining technology is evolving day-by-day especially in the environmentally conscious microsoldering for electronics assembly. To meet the demands of miniaturization, the reduction of fillet size, thickness of base metal, and weight and size of products is advancing. When considering the trends of products, it is important to develop the microsoldering and microbrazing technology from every viewpoint. Table 2.8 summarizes the important points for consideration of the important key word ‘micro’. The points are classified into three groups: material, process and equipment. In each microjoining process, the requirements of present and future subjects are listed. Microsoldering will develop gradually; however, the systematic development of these items is important to secure joint reliability under severe conditions. The development of lead-free solder with a high performance, homogeneous heating system and IMC formation free plating are the important areas for development. In the microbrazing of aluminum, the reduction of operation temperature to prevent anneal softening of base metal is a key to reducing the weight and size of aluminum heat exchangers. In both microjoining processes, the environmentally conscious view point is quite important. The no clean process without the need for a flux removal process reduces environmental impact because it does not use cleaning water and organic solvents. The fluxless low-vacuum level operation is an environmentally-friendly future process. The further innovation of laser microjoining processes is a key to meeting the demands for dissimilar material joining.
2.7
References
Ganesan S and Pecht M (2006), Lead-free Electronics, John Wiley & Sons, New York. Humpston G and Jacobson D M (1993), Principles of Soldering and Brazing, ASM International, Ohio. Humpston G and Jacobson D M (2004), Principles of Soldering, ASM International, Ohio. Hwang J S (1989), Solder Paste in Electronics Packaging, Van Nostrand Reinhold, New York. Hwang J S (1995), Ball Grid Array and Fine Pitch Peripheral Interconnections, Electrochem. Publishers, Scotland. Hwang J S (1996), Modern Solder Technology for Competitive Electronics Manufacturing, McGraw-Hill, New York. ISO 9453 (2006), Soft solder alloys – chemical compositions and forms. ISO 9454–1 (1990), Soft soldering fluxes – classification and requirements. Part I: Classification, labelling and packaging. Jacobson D M and Humpston G (2005), Principles of Brazing, ASM International, Ohio. JIS Z 3282 (2006), Soft solders – chemical compositions and forms. JIS Z 3284 (2005) Solder paste.
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Judd M and Brindley K (1992), Soldering in Electronics Assembly, Newnes, Oxford. Klein Wassink R J (1984), Soldering in Electronics, Electrochem. Publishers, Scotland. Manko M M (1964), Solders and Soldering, McGraw-Hill, New York. Miyazaki M, Mizutani M, Takemoto T and Matsunawa A (1997), ‘Analysis of Meniscus Configuration Formed on Circular Rod’, Q. J. Japan Weld. Soc., 15(4), 674–680. Moelwyn-Hughes E A (1947), The Kinetics of Reaction in Solution, 2nd ed., Clarendon Press, Oxford. Okamoto I, Takemoto T and Uchikawa K (1983), ‘Brazability and Erosion by Molten Filler Alloy of Aluminum Base Plates with Intermetallic Compounds’, Trans. JWRI, 12(1), 57–64. Rahn A (1993), The Basics of Soldering, John Wiley & Sons, New York. Schwartz M (2003), Brazing, 2nd ed., ASM International, Ohio. Shangguan D (2005), Lead-free Solder Interconnect Reliability, ASM International, Ohio. Shoji Y, Uchida S and Ariga T (1980), ‘Effect of Specimen Size on Dissolution of Copper in Molten Tin under Forced Convection’, Trans. Japan Weld. Soc., 11(2), 148–155. Shoji Y, Uchida S and Ariga T (1982), ‘Dissolution of Solid Copper Cylinder in Molten Tin-Lead Alloys under Dynamic Conditions’, Met. Trans. B, 13B, 439–445. Takemoto T, Okamoto I and Matsumura J (1989), ‘Erosion of Pure Copper by Copper Phosphorus Brazing Filler Metals containing Tin and Silver’, Trans. JWRI, 18(2), 205–209. Terrill J R, Cochran C N, Stockes J J and Haupin W E (1971), ‘Understanding the Mechanism of Aluminum Brazing’, Weld. J., 50(12), 833–839.
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3 Fundamentals of fusion microwelding G A K N O R O V S K Y, Sandia National Laboratories, USA and V V S E M A K, Pennsylvania State University, USA
3.1
Introduction
What is fusion microwelding? A strict definition is probably impossible to agree upon, so a working definition is suggested, wherein fusion microwelding (or joining) is the joining by fusion of parts which have at least one dimension <100 µm. Thus, according to this definition, the joining of two <100 µm thick sheets (which might be much larger in-plane) is microjoining, but the joining of one <100 µm sheet to a much thicker part is not. With this definition, much of what is presently called microjoining (even in other chapters of this book) is more properly classified as ‘milli’ or perhaps, ‘meso’ joining. In this chapter the basic phenomena that interact in the creation of a fusion microjoint are introduced. These include (a) thermal aspects (heat input, heat input distribution, heat flow within the fusion zone and surroundings, the effect of boundary conditions, and heat sources), (b) mechanical aspects (the evaporative force generated by high energy density heat input, constrained thermal expansion, pool dynamics, surface tension, body forces, and externally applied forces), and (c) geometrical aspects (joint designs, gaps, shape stability). The main process variations used to accomplish fusion microjoining, as determined by their heat sources, are noted, and how these heat sources’ characteristics influence their use in microjoining. A few examples highlighting the major processes are given; however, as process examples will be given in other chapters of this book, an extensive collection is not provided here. Similarly, production concerns (variability, process monitoring, inspection, part manipulation and fixtures, testing, properties) while mentioned, will be treated in subsequent chapters and hence not addressed here. It is assumed that the reader is already familiar with macro welding; thus while this chapter is an introduction to fusion microwelding, it is not intended to be an introduction to welding per se. 51 WPNL2204
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3.2
Fundamental aspects
3.2.1
Thermal
The equations of heat and mass transfer applied in the macroscopic welding world are readily extrapolated to the micro world. However, the characteristic times and distances are distorted to shorter and smaller values than are characteristic of the macro world. Also, when dealing with transient heat flow in micro parts, near-adiabatic conditions are more likely to occur, as are near-steady-state conditions, when an effective heat sink is nearby. Nearadiabatic conditions can result in extreme heating rates when low masses and high energy density heat sources like laser or electron beams are involved! Finally, though not treated here, the basic equations of heat transfer (Fourier’s parabolic heat conduction equation) and fluid flow (the Navier–Stokes equations) eventually break down, and need to be replaced by more complicated relations taking into account the non-continuuum nature of matter and energy transport at micro and smaller scales.1 Using a commonly-applied one dimensional (1-D) diffusion approximation,2 a localized input of heat (such as obtained from a laser pulse) would be expected to diffuse according to the equation: x ~ (4αt)1/2
(1)
where x is distance, t is time, and α is the material’s thermal diffusivity. Taking steel as an example (α ~8 mm2/s)3 a thermal pulse would require ~3 × 10–4 s to arrive at a boundary ~100 µm distant, which is shorter than many welding lasers’ shortest pulse length. While one can argue that the 1D approximation is a poor choice for many real geometries, the key issue here is that nearby boundaries (nearby on the micro scale, at least) can have a dramatic effect on the thermal profile, causing it to either approximate a lumped parameter solution (if adiabatic) or reach a stationary state in relatively brief periods of time (if isothermal). For the case of a micro electro-mechanical system (MEMS) device, where feature sizes on the order of a few 10’s of µm are common, the absence or presence of a connection to the substrate (which is a massive heat sink) can make the difference between very rapid heating (i.e. adiabatic conditions) and near-stationary state conditions, where the maximum temperature reached is a balance between the power, material thermal conductivity (typically polycrystalline Si for MEMS devices, an excellent thermal conductor) and the distance to the nearest heat sink. A 1-D analytical solution (thermal conduction only) is given by Carslaw and Jaeger4 to the problem: ‘The slab with prescribed flux at its surface,’ which approximates the thermal heating of the end of a shaft to which a gear might be joined (see Chapter 15). The boundary conditions for two cases are described as follows. Case 1: the region 0 < x < l at zero initial temperature, constant flux F0 into the solid at x = l, no flow of heat at x = 0. Case 2: the
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Fundamentals of fusion microwelding
53
same, except instead of a zero flux condition at x = 0, the temperature is maintained at zero. Case 2 is of interest to this discussion. The solution is given by: T=
2 F0 ( α t )1/2 ∞ (2 n +1) l – x (2 n +1) l + x Σ (–1) n ierfc – ierfc 1/2 K n =0 2 (α t ) 2 ( α t )1/2 (2)
where K is the material’s thermal conductivity. As the solution involves the summation of exponential terms (the ierfc functions), it was evaluated numerically using MatLab.5 Resulting temperature distributions for a material with properties equivalent to alloy steel are plotted in Fig. 3.1 as a function of time. Steel pin Tm ~ 1400°C
K = 51.9 W · m–1 · K–1 κ = 1.36 × 10–5 m2 · s–1
10 s
1000 1s 500 0.1 s 0 0
0.5
1
1.5
2
2.5 3 Pin length, mm
3.5
4
4.5
(a) 30 10 s 25 Temperature, °C
Temperature, °C
1500
1 ms 20 15 0.1 ms
10 5 0 0
0.01 ms 0.001ms 0.2
0.4 0.6 Pin length, mm
0.8
1 × 10–4
(b)
3.1 Temperature distribution for isothermal end condition for (a) 5 mm and (b) 100 µm long pins vs. length and time.
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For a 5 mm long shaft, an appreciable period of time (several seconds) is required to attain melting at the heated end for the heat flux used (which was 15 × 106 W/m2, appropriate to a micro E-beam welder made by converting an SEM). With this heat flux, melting can occur only if a long enough pin is present. This is necessary to reduce the isothermal end’s heat sink effect. Notably, much of the pin is heated to appreciable temperature before melting occurs at the end. This is a two-edged sword. In some cases minimal total heat input is desired to avoid damage to nearby heat sensitive areas; on the other hand, for brittle materials such as glass and Si, sharp thermal gradients should be avoided to prevent fracture. If the pin is shortened to 100 µm, the time needed to achieve essentially steady state drops to about 1ms. However, because the heat sink is much nearer the source, the maximum temperature (for the same heat flux) drops in proportion to the new vs. old shaft lengths. Thus a much higher value of F0 would be needed to achieve melting at the end of the 100 µm shaft, but it would not need to be applied for very long. The requirement for brief periods of intense heat application is met by the ‘SHADOW’ laser welding technique,6 in which the laser beam is rapidly scanned along the joint, which will be discussed in detail elsewhere (see Section 3.6.2, and Chapter 14). Contrariwise, if the part to be melted is effectively adiabatic, such as a ‘floating’ MEMS-device feature, much less heat input is required, as the only heat loss is via convection or radiation. Particularly for high conductivity materials (such as Si), a simple lumped parameter calculation embodying conservation of energy will give adequate results for determination of part temperature vs. time. As an order-of-magnitude example, it takes 7.2 J/mm3 (= 7.2 nJ/µm3) to raise Si from ambient to a molten state.7 For an ‘adiabatic’ MEMS feature on the order of 100 µm × 10 µm × 1 µm, this would require only 7.2 mJ, which would be applied in ~10 ms at the relatively low full beam power of the converted SEM in the previous example (~750 mW). Since the part’s temperature will effectively be raised isothermally, it would be difficult to prevent complete melting, rather than controlled localized melting desired for a joining process. Fusion microwelding thus requires a careful balance between heat input and heat sinking, particularly for parts that are of micro scale extent and are not effectively heat sinked.
3.3
Forces acting on the pool
The forces acting on a molten pool can be split into two categories: materialrelated and process-related. In Table 3.1 we compile the major phenomena in each category. Further, we note which phenomena are important for the three types of heat sources commonly employed for fusion welding. Since this is an introductory chapter, the discussions presented are highly simplified, and
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Table 3.1 Forces acting on molten pools Category
Force phenomenon
Important for
Thermal expansion/contraction Evaporation Surface tension Momentum/inertial Viscosity Gravity
RW Beam Arc, beam Arc, beam, RW Arc, beam Arc
Aerodynamic Electromagnetic Extrinsic applied mechanical force
Arc Arc, RW RW
Material-related
Process-related
RW = resistance welding, Arc = various arc welding processes, Beam = laser and electron beam welding.
the reader is directed to later chapters and the literature for a fuller exposition of the phenomenon in question. Relative to the macro-to-micro size transition, again, certain factors are affected greatly, while others are totally insensitive. In the following, quantitative estimates on the relevance of each phenomenon and how it is affected by size are given.
3.3.1
Thermal expansion/contraction vs extrinsic applied mechanical force
Thermal expansion against restraint during welding is the source of residual stress and distortion in all welding processes, macro and micro. However, it is a large topic and will not be treated here. Instead we will look at how it directly affects resistance welding (RW), where an actual external force is applied by the current conducting electrodes, and the thermal expansion of the material being heated reacts against it. In particular, a standard opposed electrode geometry is assumed, composed of two collinear round electrodes on either side of two metallic sheets being welded at a spot. The thermal expansion-induced hydrostatic stresses and strains created by melting are calculated in two steps (the constants used below are for steel). First the volume expansion under no stress from ambient to the molten state is calculated from: εm = (3α ∆Tm + δm)
(3)
then the pressure needed to recompress back to the original volume is calculated from: σm = K εm
(4)
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where K is the isothermal bulk modulus for the melt: ~80 GPa,8 α is the average linear coefficient of thermal expansion: ~10–5, ∆T is the temperature difference from ambient to melting: ~1500°C and δm is the volume expansion strain upon melting:9 ~0.036. Substituting into equations 3 and 4, εm ~0.081 and σm ~6.5 GPa – truly enormous pressure! Further, none of the above factors are dependent upon the size scale. In reality, resistance welding electrodes do not try to prevent this expansion; they only try to maintain a relatively constant applied force. An estimate of the expansion that the electrodes must accommodate is obtained from: δexp ~ C (εm/3) × thickness
(6)
The product of the uniaxial strain and thickness is reduced by a factor C which takes into account that the entire ‘column’ of metal between the electrodes does not reach the melt temperature and that some plastic deformation in the metal being welded also takes place. This is neglected here for simplicity, as is the fact that the molten zone may be superheated above the melt temperature, which would increase εm. The average electrode separation velocity Ves is then approximated by: Ves ~ δexp/current duration
(7)
and the average electrode acceleration ael is obtained from: δexp = 1/2 aes (current duration)2
(8)
In a macroscopic steel weld of overall thickness 10 mm and current duration ~1/6 s, equations 6–8 give δexp ~0.27 mm, Ves ~1.6 mm/s and aes ~19 mm/ s2. In a micro resistance weld, of thickness 0.1 mm and pulse length 1 ms, δexp ~2.7 µm, Ves ~2.7 mm/s and aes ~5.4 m/s2 are similarly obtained. The effective mass of an electrode plus its holder for an industrial sheet metal welder can vary enormously depending upon the sophistication of its design, but probably ranges from 1–100 kg, while for a small bench-top weld head intended for electrical interconnections a range of values from 10–100 g is appropriate. Combining these masses with the above calculated accelerations, the average inertial forces acting on the electrodes during melting obtained from Newton’s 2nd law: f = ma
(9)
are estimated to range 0.02–2 N and 0.05–0.5 N, respectively. While the former range is negligible compared with the electrode force of 2500 N which would be typical of a sheet metal welder, the latter range is not when compared with electrode forces on the order of 1 N used for small electrical interconnection welds, implying that there may be some variation in electrode force during the weld heating. Again employing equation 9, one can estimate the acceleration that the electrode/holder combination is capable of from the
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nominal electrode force and effective mass, which is found to range from 25–2500 m/s2 for the large industrial welder and 10–100 m/s2 for the micro welder. Once the fusion zone has formed and the electrical current has shut off, a rapid solidification and cool down, accompanied by a similarly rapid contraction of the weld occurs. At least for the period of time that the molten zone is present, the electrodes should apply pressure to prevent the formation of fusion zone defects. Thus they need to ‘follow’ the surface under the influence of the electrode force. In equation 6 the molten zone fraction term C was neglected in the earlier on-melting scenario, but it is essential here (and set to 0.5), as contraction due to the solidification phase change only is relevant and not that due to cooling back to ambient. Furthermore, only the final term in the parentheses of equation 3 is applicable, thus equation 6 becomes: δcontr ~ 0.5 (δm/3) × thickness
(10)
For the industrial sheet metal weld, δcontr is estimated as 0.06 mm, and the solidification time is estimated as 0.03 s.10 Substituting δcontr for δexp, and the solidification time for the current duration in equation 8 gives the required acceleration for the electrodes to maintain contact: 133 mm/s2. While this is much larger than required during melting, it is still well within the available acceleration rate for the slowest electrode plus holder combination (25 m/ s2), so the dynamics of the industrial sheet metal welder appears to be adequate to its task. Cooling rates for micro resistance welds are at least as fast as those for laser welds of similar size, and can be on the order of 104–105 °K/s, so a solidification time range of 0.1–1 ms is suggested. At an overall thickness of 0.1 mm, δcontr is estimated as 6 × 10–4 mm. If the fusion zone takes 0.1 ms to solidify, the required acceleration to maintain pressure on the solidifying fusion zone is 120 m/s2. If the molten zone takes 1 ms to solidify, the required acceleration is only 1.2 m/s2. Thus for the micro resistance welder estimates of the available (10–100 m/s2) and required values of acceleration overlap. Clearly, changes in weld schedule, such as the electrode force, which varies both the acceleration rate and the heat extraction rate (and hence the solidification time) can have a significant effect on the ability of the welding system to maintain a desired force level on the solidifying weld pool.
3.3.2
Evaporation
A high heat flux incident on a surface may cause melting and simultaneous intense vaporization. The evaporating vapor causes a recoil force which depresses the molten surface, in severe conditions forming a deep, narrow ‘keyhole’. This phenomenon is most often encountered during high energy density beam welds (laser and/or E-beam). The depression aids energy transport
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to regions below the surface of the material being irradiated, and results in much deeper melting (penetration) than would occur by convection or conduction heat transport. In extreme cases the recoil force may cause molten material to be ejected from the weld pool, causing ‘spatter’ and drilling. It was shown by S.I. Anisimov,11,12 that the recoil pressure produced due to the surface evaporation at a given temperature is proportional to the pressure of saturated vapor at this temperature: prec(T) = C psat(T)
(11)
For the cases of evaporation into a vacuum or strong evaporation into atmospheric pressure, (where the evaporation is much greater than the equilibrium vapor pressure) C = 0.54. The saturated vapor pressure as a function of temperature can be found from the Clausius–Clapeyron equation. Assuming the vapor behaves ideally, this is given by: dp/dT = (SV – SL)/(VV – VL) = (L/T)/(VV – VL) ~ L p/RT2
(12)
where L is the molar latent heat of evaporation, p is the saturated vapor pressure, T is the temperature, R is the universal gas constant, and SV, SL and VV, VL are molar quantities of entropy and volume (where the subscript indicates vapor: V or liquid: L), respectively. The approximation in equation 12 occurs by neglecting the molar volume of the liquid relative to the vapor and then employing the ideal gas law to calculate the vapor’s molar volume. Most materials’ vapor phases behave non-ideally and consequently, the dependence of the saturated vapor pressure on temperature differs from the theoretical formula deduced above. Experimentally-measured and analyticallyexpressed dependencies of saturated vapor pressure on the temperature for pure metals are tabulated.13 For metal alloys and non-metallic materials, the relatively little data available is not conveniently accessible.14–18 Also, the available experimental data typically provide the dependence of saturated vapor pressure only for temperatures up to normal (ambient pressure) boiling temperature. Because of non-equilibrium caused by high heat flux, in electron and laser beam welding the surface temperature often exceeds the normal boiling point. Without providing details given elsewhere,19–21 the reader is asked to accept that the beam/material interaction is too rapid for bubble formation. Thus, boiling, i.e. volumetric evaporation, as opposed to surface evaporation, is improbable in beam welding. Hence, values for the saturated vapor pressure for temperatures exceeding the boiling temperature must be determined by extrapolation of the existing data. Experience accumulated in laser welding modeling and validation experiments22,23 demonstrates that acceptable accuracy of numerical predictions can be obtained for a wide variety of materials such as steels, aluminum, copper, and titanium aerospace alloys using the following formula for evaporation recoil pressure:
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pr(Ts) = 0.54 B0 (Ts)–1/2 exp(–L/RTs)
59
(13)
where B0 is an evaporation constant (determined for a given material by equating the saturated vapor pressure at the boiling point to 1 atmosphere), and Ts is the surface temperature. Recoil pressure values of 1.5 (meas.) – 7 kPa (calc.) for keyhole beam welding conditions have been found.24,25 Some investigators have reported much higher measured values,26 but did not correctly take into account the area that the keyhole pressure acts upon (essentially the whole top surface, not just the keyhole only, since the pressure wave expands at the sonic velocity) in their experiments. None of the factors in this equation are size dependent.
3.3.3
Surface tension
Surface tension (σ), also referred to as surface energy, can act as either a driving force or as a restraining force. It acts as a driving force when differences in surface tension between connected interfaces may be present, i.e. the classic case of wetting and spreading of a droplet on a surface where the vector sum of the surface tension forces of the liquid/vapor and liquid/solid interfaces is less than that of the solid/vapor (so critical to brazing and soldering processes), and also in the case where a thermal gradient in a liquid drives a flow because of the temperature dependence of the liquid/ vapor surface tension (Marangoni flow). Surface tension acts as a restraining force which tries to minimize the surface area of a fluid volume, e.g. where relieving supersaturation of gas in a liquid expands a pore, or where the surface tension tries to collapse a deep, narrow keyhole being generated by evaporative recoil pressure in a high energy density weld. Some magnitudes of these forces are calculated in the following. As a restraining force minimizing the surface area of a liquid droplet of spherical shape of radius r (in the following, the applicable surface tension is at a liquid/vapor interface) the surface tension generates a hydrostatic pressure p (also called capillary pressure) given by: p = 2σ/r
(14)
Values of σ ~ 0.5–2 N/m (or N-m/m = J/m ) are typical for liquid metals.27 The capillary pressure is inversely related to the radius of curvature of the liquid surface. Table 3.2, showing the variation of capillary pressure vs. radius of curvature demonstrates how important surface tension becomes when sub-mm dimensions are involved. The temperature dependence of surface tension (Marangoni effect) also results in a driving force according to the following: 2
dσ/dr = (dσ/dT)(dT/dr)
2
(15)
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where dσ/dT ~ –0.3 × 10–3 N/m-K for a pure metal or 0.36 × 10–3 N/m-K when a surface active element is adsorbed on the surface.28 (It may be easier to understand that equation 15 represents a force by identifying it as the gradient of an energy.) If dT/dr ~1200°/0.005 m (dT is approximated as the difference in boiling and melting temperatures), then dσ/dr ~ +/–70 Pa. Note that this is a shearing stress, rather than a pressure as in equation 14. Also, in a manner similar to the capillary pressure, the shearing stress, though negligible when the temperature gradient extends over centimeters, increases by 4 orders of magnitude, and becomes quite appreciable when it extends over microns. Finally, since the surface tension temperature coefficient can be either positive or negative, the flow can be up or down gradient, causing significant changes in metal flow. Note that the flow is driven from the area of low surface tension to that of high surface tension. A final effect caused by surface tension is its capability to smooth the surface, such as the ability of oil on the surface of water to damp wave action.29 In mm-size or larger welds, ripples are extremely common, especially in pulsed welds. When the weld itself becomes a few tens of microns in extent, the forces required to perturb the surface on a finer scale become appreciable (see Table 3.2), and other than solidification-induced irregularities which are driven by intersections of grain boundaries with the surface, the surfaces of such small welds tend to be quite smooth.
3.3.4
Pool stability
Drilling vs welding In the previous section the balance between the cavity vapor pressure and the capillary surface tension is noted as controlling the stability of a deeply penetrating keyhole weld. If the evaporative forces are too high, metal is actually expelled from the molten pool (‘spatter’), resulting in drilling or cutting, rather than welding. Even in the micro world, where workpieces to be welded are thin, there is a need for keyhole welding, as its high depth/ width ratio minimizes the volume of melt, and thus reduces heat input and resulting thermal effects such as distortion. However, this must be balanced by increased concern to avoid drilling instead of welding.
Table 3.2 The internal pressure vs. radius of curvature relationship @R
1 cm
1 mm
1 µm
p (Pa)
2 × 102
2 × 103
2 × 106
Note: 2 MPa corresponds to ~20 atmospheres!
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In laser welding, the heat input comes from a surface source whose extent is given by the focused spot size. If a molten zone of radius r is desired, the heat input required Qr is thus proportional to r3. The laser heat input is given by the time integral of the laser power P(t), or Pt for simplicity (where t is the laser pulse duration and where 100% energy transfer is assumed). If we wish the laser weld to be efficient, we must restrict thermal diffusion, thus the laser pulse duration will be given by equation 1 (with the distance x replaced by r). Equating laser heat input Pt to required heat input Qr and rearranging t to the opposite side gives the following equation (where C and C′ are constants involving geometrical and material thermal constants): P = Qr/t = C r3/t = C r3/(r2/4α) = C ′ r
(16)
The intensity of the laser irradiation required is proportional to P/r2, which gives an overall laser intensity vs. fusion zone size dependency of C ′/r. Thus, as the fusion zone size decreases, the required laser intensity increases inversely. Eventually, the intensity increases to a point where it is impossible to avoid keyholing and a transition to a drilling mode. Fortunately, the creation of a raised ‘crown’ of metal around the vapor cavity caused by the central keyhole depression (shown in Fig. 3.2) will produce an extremely high radius of curvature, which provides a substantial retardation force against complete expulsion. A simplified analysis of this situation has been given,30 which Laser beam
Wl
2wm
h
rm
vm (a)
Wl
2wm
rm h
vm (b)
3.2 Schematic diagram of vapor recoil vs. capillarity balance.
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while simplifying the geometry and averaging properties to allow for computational convenience, nevertheless includes all the factors acting: heat transfer, mass transfer, recoil pressure and capillarity. Basically it was found that capillarity forces prevent expulsion during fusion micro melting. The results for some typical laser pulse conditions are shown in Fig. 3.3. It should be emphasized that this is an extremely dynamic condition, and that the balance between the vapor recoil force and capillarity does not necessarily exist during the entire pulse. Thus, at early stages, before the crown is formed, some spatter may occur, which is then stabilized as an appreciable melt volume builds up. Similarly, if too much energy is provided, the crown increases in size, and decreases its radius of curvature, again departing from balance, which may then allow spatter late in the pulse. De-coalescence upon solidification In making an actual weld, there are two or more pieces to be joined across a gap or gaps. Often the parts are heated simultaneously, and form initially separate molten zones. When these grow sufficiently, and are brought into contact, they coalesce and form a single pool. If the pools are brought together quiescently, i.e. they are brought together by thermal expansion, or by the rounding of square corners, etc., and there is not a lot of oscillatory motion, the merged pool will generally be stable. However, if an appreciable gap
Absorbed laser pulse energy, mJ
10000
1000
100
Edr, 0.1 ms Edisp, 0.1 ms Em, 0.1 ms Edr , 1 ms Edisp, 1 ms Em, 1 ms Edr , 5 ms Edisp, 5 ms Em, 5 ms
10
1
0.1 0
100
200 300 400 500 Melt pool (laser beam) radius, µm
600
700
3.3 Calculated threshold absorbed pulse energies for surface melting (Em), melt displacement (Edisp), and melt ejection (Edr) as a function of the melt pool radius for pulse durations of 0.1 ms, 1 ms and 5 ms for iron, assuming 2wm = 0.1wl.
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must be bridged, which may occur because of pool motion rather than growth, there is a real possibility that the lowest surface energy configuration will not be a common pool, but instead two separate ones. In this case, it is possible that de-coalescence will occur before solidification takes place. In the macroscopic world, where it is common to add filler material to make up the presence of gaps, this is rarely a problem. However, in the micro world, where the addition of filler material is a practical impossibility, de-coalescence is a real concern. Two examples are given in Figs 3.4 and 3.5.31 In the former, a round column of liquid metal is made through a sheet (which models a spot weld through two tightly clamped sheets), and it is assumed that some fraction of metal is lost due to spatter. The process diagram indicates that below a certain amount of metal loss, which depends on the aspect ratio, a hole will not be formed, but instead a thinned membrane will result. In the latter case, a joint is being made between two butted bars across a gap. Again, the process diagram indicates the transition between a stable bridge across the two versus separation into two separate zones.
3.3.5
Momentum/inertial forces
Many instances of the effect of momentum (P) and/or inertia in welding can be found. An instance was brought up in the discussion of external forces applied during resistance welding in an earlier section. In this section several more important instances are treated. 0.8 0.7
Drilling
Ejected fraction
0.6 0.5 0.4 0.3 0.2 0.1 0 0
Self-healing 0.5
1 1.5 2 2.5 Weld pool aspect ratio, h/r
3
3.5
3.4 Process diagram showing critical ejected fraction as a function of pool aspect ratio. Below the critical fraction, the weld pool is selfhealing.
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Filler volume, µm3
3
2.5 Stable ligament
2 De-coalescence
1.5
1
0.5 1
1.5
2
2.5 3 3.5 Beam diameter, µm
4
4.5
5
3.5 Process diagram showing filler volume needed as a function of beam spot diameter. Above the critical volume, the weld ligament is stable and spans the gap; below it, the ligament decoalesces. Dots with error bars are data points; the line is a parabolic curve fit to guide the eye.
Newton’s 2nd law in elementary form has already been encountered as equation 9; its more comprehensive statement relating force to momentum change is: f = m a = dP/dt
(17)
The effect of size enters directly via the masses involved. As the mass depends upon the cube of the dimension, these forces will vary with the cubic power of size, as compared with the inverse dependence affecting capillary forces. Similar to surface tension, the momentum force may be a driving force or a resisting force. Consider a stream of droplets impinging upon a molten pool’s surface. This is commonly encountered with the GMAW process at moderate currents, but could also be represented by a stream of laser-melted powder in a deposition process. The following example uses data for steel filler being welded at 150 amperes.32 Droplets of 1.2 mm radius (r) created from a consumable electrode (the density ρ of liquid iron at the melting point is ~7 × 103 kg/m3), have a mass given by: m = ρ (4π/3) r3 = 5 × 10–5 kg
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Combined with a droplet velocity ~ 0.76 m/s, this mass yields a droplet momentum P = mv ~ 3.8 × 10–5 kg m/s
(19)
Assuming the droplet’s entire momentum P is transferred in a time period ∆t approximated by the droplet diameter divided by its velocity: ∆t ~ 2r/| v | ~ 0.0032 s
(20)
the approximate average force will be given by: | f | = d | P |/dt ~ | P |/∆t ~ 12 mN
(21)
Applying this force over an area equal to the droplet’s projected area (~4.5 mm2), gives an average pressure of ~2.7 kPa (~0.03 atm). If one assumes instead that the particles are 0.12 mm in radius (i.e. micro-sized), but maintain the same velocity, then P decreases by 1000, ∆t decreases by 10, and f decreases 100-fold to 0.12 mN. The projected area also decreases 100-fold, so the pressure remains approximately unaffected. In laser-based powder deposition processing,33 while the 0.12 mm particle radius is approximately correct (though at the high end of the range of particle sizes typically used), the typical particle velocities are in the range of 2–5 m/s. Hence, P ~ 1–2.5 × 10–7 kg-m/s, ∆t ~ 24–60 µs, f ~ 1.7–10 mN, and the pressure ranges from 3.8–20 kPa. Another instance of momentum affecting the fusion zone is a consequence of rotational motion caused by a forced vortex in high current GTAW welding. At sufficiently high current (>250 amperes) the vortex forces the molten zone to assume a whirpool shape with a depressed center.34 This central depression acts similarly to the vapor-formed keyhole seen in laser or Ebeam beam welding in promoting enhanced penetration, though the mechanism is quite different. The body force created by centrifugal motion is given by: fv = ρ ω × (ω × r)
(22)
where ω is the pool rotational speed in radians/s and r is the radius of gyration. (Note: this is a consequence of the electromagnetic Lorentz force, given by j × B, the cross product of the current density and the magnetic field, respectively, and which will be treated in a later section.) For liquid steel in a macroscopic weld pool,34 with ω = 188 radians/s at r = 1.5 mm, fv = 3.7 × 10–4 N/mm3. This exceeds the body force due to gravity on the equivalent volume of steel: 6.9 × 10–5 N/mm3, and consequently allows the periphery of the pool to be raised above the initial surface until surface tension and gravity combine to oppose it. In a micro scale pool r would be at least an order of magnitude smaller; thus to achieve an equivalent force, ω2 would need to increase by 10 (an increase of ~3.16-fold for ω). While the forced vortex appears to be limited to high current GTAW, and thus a micro
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fusion comparison is moot, stirring of pools due to the Lorentz force causing a streaming, inverted fountain-like pool motion35 is commonplace in lower current arc welding (and in resistance welding as well). It has been found that the mixing time of a macroscopic fusion zone stirred electromagnetically (which is inversely related to the pool’s average flow velocity) is approximately proportional to the pool volume at constant current, and inversely proportional to current squared at constant volume.35 It has also been found experimentally that the volume of a GTAW weld pool is proportional to the weld current to the power 2.3, all other factors (voltage, travel speed, thickness) held constant.36 Combining these relations, if the pool size decreases 10-fold (from 1 cm to 1mm), the volume will decrease by 1000-fold, requiring a current decrease of 20-fold (since the minimum weld size in the above study was ~3 mm wide, this is a slight extrapolation, however the author’s experience is that ~mm wide GTA welds can indeed be made at ~25 amperes predicted from the above study). This change in current implies that the average pool flow velocity will decrease by 400-fold; however, since the volume of the pool has decreased by 1000, the inverse volume dependence implies the velocity will increase by the same factor. Thus the net change in velocity will be on the order of 1000/400, ~ 2.5-fold. Given the approximations and extrapolation used in arriving at this estimate, especially since it does not take into account the effect of viscosity which should increase with the much increased surface/volume ratio for the smaller pool, it is reasonable to assume that the average pool velocity (and momentum/ volume) will increase by less than an order of magnitude upon decreasing from the macro to the micro world. Another example is the sloshing of a molten pool,37 approximated by: f = dP/dt ~ 2 (pool mass × pool velocity)/(slosh period/2)
(23)
where the factor of 2 comes from the reversal in direction of the pool motion, and where: pool velocity ~surface displacement (peak-to-valley) /(slosh period/2)
(24)
Given a pool mass of 0.3 g, a surface displacement of ~0.2 mm, and a period of 0.005s, a value for f of ~20 mN is calculated via equations 23 and 24. This is of the same order of magnitude reported for arc forces.34 In order to predict how the pool inertial force will vary with size, we must look at the terms in equations 23 and 24. It is clear that the pool mass varies as the pool volume, or as a characteristic dimension (r) cubed. For varying pool size, the natural period has been found to correlate with the square root of the pool mass.37 This implies that the period will correlate with the square root of the pool volume, or a characteristic dimension of the pool to the 3/2 power. One would expect the surface displacement to be proportional to the
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size of the pool in the direction of motion (basically, this assumes geometrical similarity); however, for micro welds, in addition to the force of gravity, we know that surface tension will also play a greater role as a restraining force at small sizes, and it has an inverse size dependence (equation 14). Since these combine additively within a term, rather than multiplicatively, we can only suggest that the surface displacement will correlate with the pool characteristic dimension to a power less than one, especially at smaller size scales. Collecting terms for equation 23 (with equation 24 substituted for the pool velocity), in the numerator we have pool volume (from the mass) and pool dimension<1 (from the surface displacement) combining to give a factor of r<4, while in the denominator the square of the period appears, which correlates with ((r3)1/2)2 or r3. Thus the overall dependence of momentum on pool size is weak, at less than the first power. Experimental measurements of pool flow velocities suggest that in mm-size laser welds flow velocities in the range of 0.25–2.5 m/s are present, depending upon whether they are surface (faster)38 or bulk (slower).39 Similar measurements in GTA welds40 of cm-size suggest a bulk value of 0.85 m/s, which is of the same order of magnitude, in agreement with the weak dependence suggested above. A final instance where momentum/inertia and Newton’s 2nd law of motion may be invoked is in the movement of parts. Assume a CNC stage is making a part ‘turn a corner’ of 1 mm radius at 25 mm/s travel speed, under an arc or beam. Equation 22 again applies (however, recall from elementary physics, that equation 22 is equivalent to mv2/r). The associated centrifugal force forcing the pool to slosh outward is calculated to be 4.4 mN for a 1 cm3 weld pool. If the corner radius is 0.1 mm, the force increases 10-fold to 44 mN. If the travel speed is increased, the force increases with the square of the velocity. In terms of acceleration, the previous cases correspond to 0.06 and 0.6 g’s, respectively. Since the mass of the pool is proportional to its volume, the force decreases proportionately for micro welds. For a 1000 µm3 weld pool, the forces acting become 4.4 pN and 44 pN, for the 1 mm and 0.1 mm turn radii, respectively. However, at the micro scale, it is likely that the radius of curvature of the corner will be proportional to the size of the weld. At a 10 µm radius the force increases to 0.44 nN, and at 1 µm, 4.4 nN (corresponding to 6 and 60 g’s) If instead we consider the forces acting at the level of the part and fixturing, rather than at the weld pool, the overall force trying to upset the part is calculated by replacing total mass for that of the weld pool in equation 22. Since macro parts plus fixtures can weigh many tens of kg, the forces required can be many tens of N. While masses, and hence forces, are greatly reduced for individual micro parts and fixtures, in order to increase production throughput, multiple parts are usually grouped on the same part holder, somewhat cancelling this benefit. Thus approaches which avoid moving parts and fixtures, like beam scanning of laser or electron beams (wherein
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the momentum of lightweight mirrors or electrons is controlling, and which is not a variable depending upon the fixture) may be attractive.
3.3.6
Viscosity
In the liquid state viscosity resists fluid shearing, which of course affects velocity distributions in the molten pool caused by other forces. The controlling equation is: τyx = –µ dux/dy
(25)
where τyx is the shear stress due to viscosity, ux is the velocity in the x direction, (the direction in which the shearing motion is occurring), and y is the direction normal to the shearing direction. Typical values of µ (the absolute viscosity) for molten metals are ~5–7 × 10–3 Ns/m2 for molten Fe and ~1 × 10–3 Ns/m2 for molten Al.41 As an example, for fluid flow in the unmixed layer next to the solid HAZ of a steel macro scale weld, dux ~0.1 m/s, dy ~0.1 mm, resulting in a value of τyx ~6 Pa. In the micro world, the entire fusion zone is smaller than the stagnation layer thickness noted above. This implies that any lamellar flow layer in the pool is close enough to a boundary which may affect it, implying that viscosity has an increased importance. Indeed, for micro scale fusion zones, calculation of the Reynolds number, (ρuxD/µ) which is the ratio of inertial to viscosity terms, implies that it will be quite small, since it is proportional to the characteristic size D. For the above macro weld example, using data for steel, the Reynolds number is ~1000; by simply changing D to 0.1 mm, it drops to 10. With a decrease in velocity, it will decrease even more. Among other implications, low Reynolds number flows tend to exhibit little turbulence, and do not mix very well, thus dissimilar material micro welds may be inhomogeneous.
3.3.7
Aerodynamic forces
Similar to the previous section, this effect is due to the flow of a gaseous fluid (instead of a liquid) across the molten surface of the pool (instead of across the liquid/solid interface), which creates a drag force which promotes flow in the same direction of motion, and may even lead to ‘waves’.29 The needed relation is again equation 25, except that the appropriate constant and values of gas motion and gas boundary layer thickness must be used. Even though the gaseous viscosity is typically 10 times less than that of a liquid, because the stagnation layer thickness is thinner, and the gas velocity is faster, the shear force is ~100 times larger. Some calculated peak values of aerodynamic shear stress in arc welds include ~70 Pa for a 3 mm, 150 ampere arc,42 and ~400 Pa for a 1 mm, 200 ampere arc.43 No obvious changes are seen when transitioning to the micro scale.
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Perpendicular impingement of a gas jet of density ρg and velocity v on the surface of a molten pool creates a peak stagnation pressure p (which results in a surface depression) where the center line of the jet intersects the pool surface.44 This is given by: p = 1/2 ρg v2
(26)
In a welding arc, an arc jet is caused by the diverging electric field, and various investigators have modeled and measured the stagnation pressures resulting, calculating ~400 Pa for a 3 mm long, 150 ampere arc,42 and measuring ~1–6 kPa for 8 mm long, 300–600 ampere arcs.43 The arc pressure increases as the current squared.34 The surface depression expected due to a given stagnation pressure in the micro scale will be reduced somewhat because capillary forces due to high surface curvature will become involved as well as fluid flow and gravity forces. While arc welding is prevalent in the macro world, arc heating processes are rarely encountered in the micro world due to problems of instability and control at currents below 1 ampere. This is mitigated by pulsing the arc. In other words, a relatively high current is used, but it is only sustained briefly. Two processes that employ a brief, transient arc are pulsed (or pulse) arc welding, and percussive arc welding. The difference is that in the latter the arc is drawn between two parts which are impacted together shortly after the arc is initiated. Both of these processes are used to accomplish spot welds, rather than seam welds.
3.3.8
Electromagnetic force
The Lorentz force F caused by the interaction between an electrical current and a magnetic field (which includes its self-induced field) is given by the well-known cross product of the current density j and magnetic field B: F=j×B
(27)
Several possible scenarios in arc and resistance welding which are affected by this physical effect include: (a) the stirring of a molten pool, with the possible creation of a whirlpool if a free surface is present, as has been already mentioned, (b) pinch-off of molten droplets from a consumable electrode from which an arc is drawn to the workpiece (Haidar and Lowke calculated 6 kPa on GMA droplets45), (c) creation of a plasma jet in a diverging (constricted at the electrode, less constricted at the work piece) arc column, and (d) ‘kick’ of conductor cables or electrode arms due to current passage in resistance welding. The magnetic field as a function of radial distance r around a straight conductor carrying a current of i amperes, and the attractive force F (per
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length L) between two such conductors separated by a distance d are given by:46 B = µ0 i/(2πr)
(28)
F/L = µ0 i2/(2πd)
(29)
(If the current in the two conductors is in opposite directions, the force is repulsive.) As noted earlier in the rotational inertial body force section, the factor of current squared appears in the force equation. In the case of arc welding, the volume melted was found proportional to the current to the power 2.3. In resistance welding, the heating value of electrical current/unit volume is given by j 2ρ where ρ is the resistivity. Since ρ is not size dependent (at least, not until the thickness of the material is <<1 µm), this implies a current squared dependence of the molten volume. Since we reduce the volume by about a million-fold in moving from cm to sub-mm sizes, the current level will decrease a thousand-fold, and the force will also be decreased a million-fold, similar to the molten volume. For a macroscopic resistance weld in steel of ~1 cm size, ~10,000 amperes are recommended for 1/6 of a second.47 From equation 29 the resulting force/length acting to separate the electrode arms (since the currents will be in opposite directions), which are typically spaced by ~10 cm, is 200 N/m. Since the force holding the electrodes together is typically 2kN, this is a relatively unimportant effect. For an analogous micro weld, perhaps 100 amperes is needed (typically the pulse will be only a few ms long), and the electrode arms will only be ~2.5 cm apart, so the force will only be 80 mN/m, and small compared with the electrode force, which would typically be about 1 N. A number of conductor spacing/current level/resultant force combinations representative of both macro and micro resistance welding scenarios are collected in Table 3.3.
Table 3.3 Force/length between current-carrying conductors spaced by distance d d I 10 A 100 1000 10000
5 cm
5 mm
0.5 mm
0.05 mm
0.0004 N/m 0.04 4 400
0.004 N/m 0.4 40 4000
0.04 N/m 4 400 NA
0.4 N/m 40 NA NA
NA: not applicable. If current flow in both conductors is in the same direction, the force is repulsive, if in opposite directions, it’s attractive.
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71
Gravity
Finally, the last force acting is due to gravity, and is sometimes called, in analogy with hydraulics, the metallostatic head. As in hydraulics, it is given by: ∆p = ρLg ∆z
(30)
where the liquid density ρL = 2000–8000 kg/m , and the gravitational constant g = 9.80 m/s2. Typical values of the level difference ∆z = 0.0001–0.01 m, result in ∆p of 2 × 100–8 × 102 Pa. This is an important factor in out-ofposition arc welding of large structures such as pipelines or industrial structures, where the molten zone is on the order of cm in extent, and is only balanced by surface tension (with relatively large radii of curvature) and arc pressure. Gravity also plays a role in generating a buoyancy force in the molten pool, caused by differences in density (Archimedes’ principle). This is usually due to either a temperature gradient: 3
∆ρ = ρ (3α ∆T )
(31)
or composition gradient (replace ∆T with the change in composition: ∆C, and 3α with ∆ρ/∆C). The linear thermal expansion coefficients of metals are on the order of 10–5 °K–1, while the temperature differences in a weld pool are usually less than 103 °K implying that relative density differences due to temperature are ~0.03, reducing the effect substantially compared to the metallostatic head at an equal (vertical) distance. Further, in a quiescent pool, the hottest metal is usually found on the pool surface where the heat source impinges, and hot metal is less dense, so the buoyancy force tends not to have a convective effect.35 However, if dissimilar density metals are being welding together, appreciable density differences (up to a factor of 5 or more), can occur, and unstable density distributions (i.e. higher density material above lower density material) can occur, which will create convection currents. This also occurs in autogenous keyhole welds, where the hottest material may be deep in the keyhole, and may be readily swept into the bottom of the pool.
3.3.10 Summary of forces Table 3.4 summarizes the calculations and measurements of the previous sections and allows comparison of which effects are significant in the macro and micro worlds. While both a pressure/stress and force are usually provided for the macro and micro cases, in some instances it makes no sense to do so. It may readily be seen that the thermal expansion force is always important, no matter the size scale. Other important forces acting on the macro scale pool include the evaporative recoil force, several types of inertial force, and the aerodynamic stagnation pressure. On the micro scale, as noted already surface tension-related forces are important, as are inertial-origin forces.
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Table 3.4 Macro vs. micro comparison of various origin stress/force magnitudes acting on the weld pool Force origin
Macro example magnitude pressure/stress force
Micro example magnitude pressure/stress force
Thermal expansion (M)
6.5 GPa (ambient-to-molten) 6.5 kN
6.5 GPa (ambient-to-molten) 0.65 N
RW electrode inertial force, melting (P)
0.02–2 N
0.05–0.5 N
RW electrode inertial force, solidifying (P)
0.13–13.3 N
0.1–10 N
Laser or E-beam driven evaporation (P)
1.5–7 kPa 1.5–7 mN
1.5–7 mPa 0.15–0.7 µN
Surface tension pressure (M)
2 kPa 2 mN
200 kPa 20 µN
Marangoni shear stress (M)
70 Pa 70 µN
0.7 MPa 70 µN
Impinging stream of droplets (P)
2.7 kPa 12 mN
67 kPa 3 mN
Lorentz force vortex centrifugal (P)
3.7 × 10–4 N/mm3 37 mN
NA
Pool sloshing (M) Stage motion, centrifugal (P)
200 mN 44 mN
20 µN 44 nN
Liquid viscosity shear stress (M)
6 Pa 6 µN
6 Pa 0.6 pN
Aerodynamic shear stress (P)
70–400 Pa 70–400 µN
70–400 Pa 7–40 pN
Aerodynamic stagnation pressure (P)
0.4–600 kPa 0.4–600 mN
0.4–600 kPa 0.04–60 µN
Lorentz GMAW pinch (P)
6 kPa 6 mN
NA
Current carrying conductor ‘kick’ (P)
200 N/m 200N (assumes 1 m cable)
80 mN/m 8 mN (assumes 10 cm cable)
Gravity (M) metallostatic head
0.8 kPa 0.8 mN
20 Pa 0.002 µN
Gravity (M) buoyancy force (∆ρg): Autogenous* Dissimilar†
2.1 kN/m3, 21 mN 52 kN/m3, 52 mN
2.1 kN/m3, 2.1 µN 52 kN/m3, 5.2 µN
M: material-related force; P: process-related force; NA: not applicable. Equivalency between stress/pressure and force assumes 1 mm square or 1 mm cube for macro forces and 10 µm square or 10 µm cube for micro forces. *Due to temperature gradient only in steel, †due to compositional gradient only, Al vs. Cu dissimilar density combination.
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Practical aspects
3.4.1
General discussion of major process variations
73
Beam welding Some heat sources are directional: lasers, electron beams, constricted arcs (as in plasma arc welding); while others are not: resistance heating, unconstricted arcs (as in gas tungsten arc welding), ultrasonic energy. If a directed energy source is present, the energy must usually pass through the ‘top’ part to the ‘bottom’ part, if the parts are not coplanar. This lapped geometry helps mitigate the very difficult edge-to-edge alignment problems encountered with thin and flimsy microparts. However, unless very effective thermal contact between the two is available, the top part may perforate (i.e., a physical separation may develop at the junction of the fusion and unmelted zones). Another way of achieving this joint is to accept the fact that the top sheet will be perforated, and mitigate this. In such a scenario, one would supply enough energy to create a molten pool in the bottom part that is ‘splashed’ outward by the recoil force, which then ‘reconnects’ to the perforated top part. (This scenario assumes that the bottom part is somewhat more massive than the top one, otherwise it, too, will perforate, and there will be no splash.) Using a broadly similar mitigation scheme, a procedure for welding thin parts together which is robustly tolerant of interfacial gaps has been patented by Philips, Eindhoven, Netherlands. It is called laser ‘spike’ welding.48 In this procedure, a specifically tailored laser weld pulse is used. Initially, a low power density is used to melt the top layer of a multi-layer stack; next a higher power density pulse is added, which via the recoil force mechanism described earlier, pushes the molten top layer into contact with the layer below it. This state must be maintained until wetting and coalescence of the top and bottom layers occurs. By balancing the power densities and amounts of power in each portion of the pulse, gaps which represent a large portion of the thickness of the material being welded can be joined. One could also conceive of a similar process employing an electron beam. Lasers and electron beams are of course ‘directional’ sources of energy. Electrons used in E-beam welding (typically accelerated at 10–40 or 50–150 keV for micro and macro applications) have appreciable penetration depth and generally couple well with industrial materials (except perhaps for nonconductors, where a charge may build up), whereas laser light has extremely little penetration capability and may be reflected away depending upon the wavelength and the target material’s reflectivity. Thus, in the micro world, the two processes may not be blindly treated as equivalent. The laser is predominantly a surface source of energy, while the E-beam is a volume source of energy, unless the accelerating voltage is atypically low (less than a few keV). One should always check the relative depth of electron penetration vs. the part thickness to determine which regime is active.
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Arc welding The ability of an arc to be a directional heat source depends upon several factors. These include its effective size at the workpiece and the current density, which control the velocity of the arc plasma and its force at the workpiece. In turn, the arc ‘size’ is controlled by many factors including the arc gap, whether or not the electrode melts (consumable filler metal electrode vs. non-consumable tungsten electrode), details of the electrode’s end preparation (if non-consumable) or diameter (if consumable), and in the case of a consumable electrode, the form of metal transfer involved (short-circuit, globular, or spray). Also important is whether any arc-constricting orifices are used to ‘channel’ the arc, as in plasma arc welding (PAW). The smallest orifice diameter typically available for the PAW process is ~1/2 mm, which would be an approximate upper limit on the arc’s size. This is still too large for true microwelding. Without arc constriction, the standard (i.e. no orifice) gas tungsten arc process has an even larger arc ‘size’, and hence is even less applicable to microjoining. However, in addition to constricting the size of the arc column by an orifice, it is possible to limit its size by the geometry of the parts being welded. An obvious geometry employing this concept is that of joining small diameter wires. Another manner of constricting an arc is to arrange for its generation between a sharp feature and a nearby featureless target. This approach is employed in the percussive arc process, wherein a sharp nib is placed on one of the parts to be joined. At least two methods of initiating the arc are employed. In one, short circuit contact occurs between the two parts to be joined. The second variation of this process employs a high frequency a.c. breakdown potential. In both process variants, one of the parts is held with a slip fit by an actuator which is attached to a charged capacitor. In the first variation described, upon initial contact a rush of high current from the capacitor basically explodes the nib, creating a localized metal vapor-filled small gap, across which an arc immediately establishes itself. In the other, breakdown occurs across the gap before contact occurs. Continued motion of the actuator then causes contact, which quenches the arc and forces coalescence of the parts, which at this point have molten surfaces from the action of the arc. Given that the nib is small (a fraction of a mm) and the actuator velocity is appreciable (~1 m/s), the arc lifetime is generally on the order of a few hundred microseconds at most. In addition to constricting the arc by the mechanical nib’s diameter and its short length and brief lifetime, a dielectric medium (most often a droplet of water) may be used to provide an additional form of arc constriction. The liquid may also act to contain the quite explosive ejection of liquid metal droplets and vapor, leaving the part surfaces surrounding the joint considerably cleaner than when no droplet is employed. The challenge to developing a robust procedure involves balancing melting enough material to provide an adequate joint while providing a good contact event that promotes coalescence without
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splashing most of the molten zone away, leaving excessive porosity in the joint. While commercial equipment is available from a number of manufacturers for both of these variations, one of the authors’ first contacts with microjoining was in welding 50 µm diameter intrinsic thermocouple wires to small single crystal specimens by holding the wire in a pair of tweezers attached to a capacitor bank and ‘jabbing’ them against the specimen. Needless to say, the tweezers were electrically insulated! With practice, and conducted under a binocular microscope, this was a surprisingly effective means of precision thermocouple attachment. Resistance welding The problems of melting across gaps would seem to be eliminated by employing resistance welding, which uses an external pair of electrodes to press the parts being joined together. However, since actual contact is made, concerns about damaging the workpieces (or the electrodes, which are also fragile) must be addressed. Further, since the contact footprints between the electrodes and the workpieces control the current density, the electrode tips must be shaped reproducibly and dressed to this size and shape as wear accumulates. Finally, at the sub-mm scale, it is very difficult to use the ‘classic’ opposed electrode configuration, where the bottom electrode is normally stationary and the parts being welded are placed against it as a ‘steady rest’ while the top electrode is brought into contact. This procedure is difficult to implement when the parts are very small, requiring precise horizontal and vertical motion capabilities of the part manipulator, or motion control of both electrodes, which is rare. It is far easier to use a one-sided approach, where only horizontal manipulator motion is needed, since all vertical motion is accomplished by the welding head itself. Also, with the one-sided approach, both electrodes are completely visible, unlike the opposed electrode case, where one is partially or completely obscured by the parts being welded. Heating during resistance welding involves a balance of Joule heating, interfacial resistance heating (particularly for materials which may form low conductivity surface films such as the oxide films which form on aluminum alloys) and the Peltier effect.47 The latter two effects, since they only occur at interfaces, become increasingly important as the parts being welded become thin. This is because while Joule heating is proportional to the thickness of the parts being welded (at a given current density), interfacial and Peltier heating are not. Interfacial resistance is difficult to model because it is dependent upon surface roughness, the presence of surface contaminants or films, the applied force, the material’s plastic deformation behavior with temperature, and the interface temperature. These factors all change during the initial stages of the weld process, so it is a transient effect. The Peltier effect, which is proportional to the difference in conduction electron energy levels on
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opposite sides of the interface, is especially important in d.c. resistance welding at dissimilar interfaces. It has the fascinating and significant effect of reversing sign with current polarity, e.g. with one polarity it is a heating effect, while for the reverse polarity it is a cooling effect (this also means that it tends to average out for multiple pulse alternating current). This effect can be manipulated to increase the heat at the weld interface surface and decrease it at the electrode interface, decreasing electrode ‘sticking’ to the part. Sticking is not desirable in macro resistance welds, as it decreases electrode life, but in micro resistance welds it is devastating, almost certainly damaging both the electrode and part. Conversely, if the Peltier effect is misapplied it has the opposite effect. Material combinations that have large Peltier effects are Cu vs. Ni (especially notable, as many electrodes are made from Cu) Ni vs. Mo and Au vs. Pd. Nevertheless, at least at the larger limits of fusion microwelding (and brazing and soldering), and particularly for joining wires or ribbon, resistance welding is a useful technique. In addition to the examples shown later of welding ~25 µm foil and wires to much thicker parts (and incidentally, violating our rule of thumb for defining fusion microwelding), a perhaps more common geometry is the cross wire weld.49
3.5
Some problems in applying fusion microjoining
3.5.1
Handling and fixturing
Fusion microjoining inherently has all the problems of macrojoining processes with the possible sole exception that one does not have to worry about large masses falling on your toes and crushing them. However, it can be equally frustrating to deal with extremely small masses. For one thing, when parts get so small that they are hard to see with the naked eye, they also tend to be hard to put down, because of residual magnetism, static electricity and Van der Waals forces. Further, if a small part is inadvertently dropped, while one’s feet may not be at risk, one’s back and head will be, from rooting about under tables and benches looking for that dropped part! Even if retrieval is successful, and the part is not damaged, it will need to be re-cleaned before use. Most fusion microwelding operations will need to be conducted in a clean bench or clean room environment, with augmented vision (microscopes, video imaging), and with scrupulously clean piece parts. In high volume production, manipulation of parts will be automated. As such, there is no possibility of looking for dropped parts, but issues of damage- and contamination-free part manipulation will become important for a successful process. Since small parts tend to be fragile, a trade-off between secure
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handling and part damage may require specific handling features to be designed into the parts. Even with augmented vision, alignment of micro parts is difficult. Self-aligning ‘keyed’ features are recommended to aid part alignment for joining. Note that even after joining, the small size of the joints may leave them quite fragile. Any fusion microjoining process contemplated for production will require a significant investment of time and thought relative to fixturing issues, which in addition to allowing efficient loading and unloading of parts, and holding them in correct alignment for joining, must also have reproducible and effective thermal contact.
3.5.2
Process characterization for repeatability
In order to produce reproducible joints, heat input vs. heat extraction must be finely balanced. In the previous section, the issue of fixturing to control thermal boundary conditions was mentioned. Additionally, the amount of heat applied and its distribution relative to the joint geometry must also be controlled. If the heating process is open loop, the process parameters which control heat input to the part must be measured and controlled to within a small tolerance. Doing this is quite a challenge, because the normal tools for process control are simply not suited to the regime of µm, µJ, and mN. Even if the tools are present (such as microammeters), extra care to eliminate extraneous noise must be made, so additional training and calibration requirements may be imposed. If feedback controls are to be used with the energy sources, such control loops must have exceptionally short time constants, because the processes take place over very short periods of time. Thus a dedicated digital signal processor (DSP) device will probably be needed.
3.5.3
Part fit-up
Finally, the part geometry reproducibility is equally important. If the parts are not carefully dimensionally controlled, even with good fixturing, the fitup at the joint will be variable. It must be remembered that the relative tolerance value is applicable, not an absolute value. So an extremely tight absolute tolerance at the macroscale (+/–0.0025 mm) may be barely acceptable or even unacceptable at the micro scale. Further, if the part dimensions vary, even if joint fit-up is not compromised, a change in the process heat input may be necessary.
3.5.4
Metrology/inspection
The same considerations apply to deciding if the process has been successful as apply to the process characterization. The quality assurance tools for
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measuring on the micro scale are just not adequate. Imagine trying to apply X-ray or dye penetrant inspection to a fusion microjoint? About all that is available is microscopic visual examination, perhaps some small-scale mechanical testing along the lines of wire pull or ball bond shear testing as used in the wirebonding industry, and of course, final device functional testing. Given that the latter will probably be chosen, adequate production quantities will need to be started so that the final yield will suffice for contracted delivery quantities.
3.5.5
Problems with modeling at small size scales
Even computer modeling of the processes can cause problems. For example, if a finite element model employs the SI unit system, dimensions of element volumes expressed in cubic meters can get so small (a 1 µm cube would have a volume of 10–18 m3) that errors may result unless extra precision calculational methods are employed.
3.5.6
Materials problems
At the micro scale, the conventional materials used in fusion macrojoining are not encountered. The materials of micro devices tend to be made by photolithographic procedures. Thus one expects to encounter Si-based materials, polymers, and electroplated materials (i.e. metallic parts made by LIGA processing*). Even the most ‘conventional’ materials encountered, e.g. gold and aluminum wires, controlled thermal expansion metal alloys (Kovar, Alloy 52, etc.) are specialty materials. While these materials may not be especially problematic for joining, the tools needed to study problems must become more sophisticated. For example, metallography of a few µm extent part is not something done in a typical metallography laboratory. Instead, electron microscopy is needed, and probably transmission electron microscopy, not ‘just’ scanning electron microscopy. Conventional electron microprobe analysis becomes less useful, as the volumes measured by their probe beams become similar in size to the fusion zone. Again, more sophisticated means of analysis are necessary. Regarding metallurgical problems, LIGA materials (which are electroplated) may have brighteners or residual stress-reducing agents added. A common agent is saccharine. When welds are made in such material, it is not unusual *LIGA: LIthographie, Galvanoformung, Abformung, acronym for high aspect ratio (depth to width), micron-level resolution, metallic part production process using synchrotron radiation to expose a masked photoresist material, followed by dissolution of the exposed photoresist, electroplating the developed cavities, dissolution of the mold and then chemically-etched release from the substrate.
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to observe porosity and cracking. An example of a Ni–Fe magnetic alloy exhibiting several simultaneous defects is shown in Fig. 3.6. As noted already, microwelds typically have a large surface-to-volume ratio. If sufficient heat sinking is available, this can lead to exceptionally rapid cooling rates. In some materials (such as 304 stainless steel), the crystal structure of the material which solidifies can be affected by the cooling rate.50 This in turn can produce solidification cracking due to microsegregation of impurities if the structure is unable to relieve stresses caused by solidification contraction. Fusion zones solidifying from welds invariably have epitaxial relations with the grain structure of their constituent parts. Unless the microstructure of these parts is exceptionally fine, it is likely that only a very small number of grains will be present in the solidified fusion zone. Since any impurities will tend to segregate to the grain boundaries, this will make them even more likely to promote cracking. The solidification interface morphology of a macroscopic weld typically starts as a planar interface and then may break down to cellular or dendritic forms. This breakdown phenomenon is controlled by the likelihood of achieving constitutional supercooling,51 which occurs when the actual temperature of the melt in front of the solidification interface is less than the local, compositiondependent liquidus temperature in the same location. The local composition is perturbed both by solute partitioning across the solid/liquid interface as solidification occurs and by diffusion and convection into the main weld pool. A steady-state liquid composition perturbation profile develops after
100 µm
3.6 Defects (hot cracks, cold cracks and gross porosity) in Ni-Fe LIGAbase material. Metallurgical section in plan view of laser spot weld.
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an initial transient. This profile is affected by the partition ratio between solid and liquid: k, given by the phase diagram, the solute diffusion rate (and convection if present, though in front of the interface there is usually a stagnant layer where convection is minimal) and the interface growth velocity. Mathematically, this can be stated as the following criterion for planar interface instability in a dilute binary alloy:52 G/v < ∆To/D
(32)
where G is the temperature gradient in front of the interface, v is the solidification velocity, ∆To is the difference between the liquidus and solidus temperatures at the nominal composition of the solidifying material, and D is the diffusion rate of the alloying element in the liquid. As noted above, an initial transient takes place which is of approximate length δ given by the following equation.53 δ ~ 4D/vk
(33)
During this transient, because the composition gradient in front of the interface is opposite to that established during steady state, planar interface stability is essentially guaranteed until steady state is approached. In equation 32, the right hand side (rhs) is fixed by the material system, so the ratio of G/v is the controlling factor as to whether non-planar front solidification and the greater segregation which accompanies it will tend to occur more or less in micro welds vs. macro welds. Calculations54 indicate that while both G and v separately increase as laser weld size decreases, the ratio actually decreases for both spot and seam welds. For 304 stainless steel, ∆To ~ 30 °K, while D for liquid metals is typically given as 10–5 cm2/ s, implying that the rhs ratio is ~30 K-ms/µm2. G/v was found to range from 0.1–1, 54 indicating the instability criterion is satisfied. The transient distance for the same conditions would be <0.2 µm (using k = 0.9), much smaller than the fusion zone of even a micro weld. It has also been found that capillarity can stabilize the interface.55 Below a given wavelength λi, perturbations of the interface will tend to decay, rather than grow. The critical wavelength is given by: λi = 2π[DΓ/(v∆To)]1/2
(34)
where Γ is the Gibbs–Thomson coefficient (given by σ/∆sf, the surface tension divided by the entropy of fusion per unit volume). Calculating λi for the above case of 304 SS gives a value of 0.025 µm, implying that this is not a stabilizing influence at the scale of interest, which is on the order of 0.1– 100 µm. Thus, it appears that planar solidification is no more likely in micro welds than in macro welds, though any cells or dendrites formed will be of small diameter, given the rapid solidification rates.
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81
Examples: laser, arc and resistance welding
In this section, a few examples of several different process fusion microwelds from the authors’ experience are shown (or as close to micro, according to our definition, that we could find!). As many more examples will be given in other chapters, this collection is not intended to be extensive, but only to illustrate significant points.
3.6.1
Laser spot weld
An example of an assembly weld for a developmental flexure device is shown in Fig. 3.7. In this scenario, LIGA-produced pure Ni cantilever springs are attached to shafts made of gauge pin material (alloy steel). The entire device (not pictured), including top and bottom plates, shafts and springs will all fit into an approximately 2 mm cube. In this case, each ‘fork’ of the spring is intended to snap to the appropriate shaft, in order to simultaneously facilitate good thermal contact between the two piece parts, and to minimize the need for fixturing. Fixturing was minimized, but pre-weld assembly was quite difficult. Additionally, the LIGA process was not quite accurate enough (at the time when these parts were being evaluated) to provide good reproducible fitup. Each short ‘seam’ weld is composed of three individually placed spot welds made with a pulsed Nd:YAG laser using an intra-cavity aperture and short focal length lens to reduce the spot size to about 30 µm diameter.
Weld 2
Weld 1
3.7 Laser welds of LIGA-based Ni cantilever springs to alloy steel shaft. Springs are 250 µm wide (in direction parallel to shaft).
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Figure 3.8 shows a single spot weld in a similar shaft-to-part geometry, except in this case the size is slightly reduced, particularly for the lower picture. The same laser as noted above was used. The materials combination is LIGA-produced Ni disk to wrought Ni shaft.
3.8 LIGA Ni disk spot welded to ~250 µm diameter wrought Ni shaft. Lower picture shows ~150 µm diameter shaft welded through hole at 3 o’clock of upper picture.
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3.6.2
83
Laser seam weld
Figure 3.9 illustrates a ‘seam weld’. In this case, the part (which consisted of two stacked disks, mechanically pinned together) was spun rapidly under a stationary beam. The entire weld was accomplished during a single ~10 ms duration laser pulse. The materials were similar to those used in Fig. 3.7. The cw-like seam weld was successful in staking the ‘pins’ in place, whereas individual spot welds were not, due to hot cracking. The elimination of hot cracking is attributed to the different dilution, as well as the less dramatic thermal cycle caused by the seam weld, compared to a spot weld. In practice, this method is called ‘SHADOW’ (Stepless High-Speed Accurate and Discrete One-Pulse Welding6,56), and is more usually accomplished by guiding the beam via galvo-controlled mirrors, while the part is stationary. Examples of fusion microelectron beam welds are shown in Chapter 15.
Pins
Part rotation
Stationary laser spot
Weld
3.9 SHADOW-like method 5 micro seam weld (part spun under stationary laser beam of ~10 ms pulse length), locking alloy steel pins holding two LIGA Ni disks together in place. This part was intentionally welded off the centerline of the pins to show their remnants.
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3.6.3
Microjoining and nanojoining
Arc seam weld
A schematic drawing of a fabricated portion of a metallic bellows is shown in Fig. 3.10. In this case, each convolution of the bellows is made by either a GTA or PAW process. This edge geometry weld serves to localize the arc heating in the width direction, as the arc localizes itself to the entire double thickness width. One advantage of fabricating the bellows by welding, rather than by forming, is that by varying the number of convolutions to weld together, custom lengths of bellows may be fabricated without investing in a unique set of tooling for each length desired.
3.6.4
Arc spot weld (percussive arc)
Besides the manufacture of thermocouple beads, (another instance where the part localizes the arc, rather than the welding process itself), another application where a spot arc weld is used is in attaching electrical ground pins or small threaded studs. In the percussive arc welding process, an arc is drawn between two parts, while one part is being accelerated at the other. Thus immediately upon creating the arc, it is extinguished by the shorting of the two parts together. In addition to stopping the arc, much of any molten zone is splashed away.
3.10 Schematic of edge geometry arc (GTAW or PAW) weld joining individual convolution of a bellows made from material a few thousandths of an inch thick, typically stainless steel or Ni alloy.
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The example shown in Fig. 3.11 is between a Ni-48Fe material (aka Alloy 52, or Niron) and 304L stainless steel. In this case the two materials are compatible. In other cases, for example between Alloy 52 and Aluminum, the molten zone is highly susceptible to solidification cracking. By splashing most of the material away, very thin fusion zones are left, which tend to be less susceptible to cracking. Thus this process is useful for incompatible material couples. Several approaches to initiating the arc are available industrially. In the part pictured, a small nib (seen in the lower photo) is formed on one of the parts. The contact of the nib to the other part completes a circuit containing a charged capacitor. The current flow resulting literally explodes the nib, which provides a continuous metal vapor between the two
3.11 Metallographic section of percussive arc welded Alloy 52 (Ni48Fe) ground pin to stainless steel substrate. Lower picture shows arc-initiating nib before welding. Dimensions of pin: 1.0 mm shaft diameter, 1.6 mm flange diameter, 0.38 mm flange thickness, 0.2 mm × 0.35 mm diameter die-formed nib.
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parts which then supports an arc. (While the pin being joined is actually 1 mm in diameter, we include this as an example of a fusion microweld because it is the nib which is being melted, and this is much smaller.) Other variants of this process use a high frequency discharge to initiate the arc. Additionally, by controlling the capacitance and charging voltage appropriately, the process may actually dispense with the arc and become a micro projection resistance weld, the subject of the next section.
3.6.5
Resistance welding
The examples shown in Figs 3.12 and 3.13 involve welding a small diameter wire onto a much larger substrate. This is yet another example where the part itself is used to concentrate the heat source. Metallography of such welds
3.12 Resistance microweld of Tophet C (Ni-16Cr-24Fe) wire to Alloy C276 pin.
3.13 Resistance microwelds of Ni wire to Alloy C276 pin, low and high energy. Note better wetting of pure Ni than Ni alloy of Fig. 3.12.
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may or may not show a fusion zone in the substrate, but as may be seen clearly in Fig. 3.13, the wire itself is usually melted. In this case, the contour of the welding electrode in contact with the wire is an extremely critical parameter. In addition to providing the electrical contact, it effectively forms a mold for the fusion zone. To ensure a low stress concentration factor, the curvature of the electrode face must be smooth and reproducible. Further, to ensure a long lifetime for the electrode, it should be made of a very durable material (in this case W) and polished. In some instances, an inert shielding gas may be applied to assist in wetting to the substrate (especially if it does not melt) and to avoid oxidation of the electrode. Figure 3.14 shows an example of a very thin sheet being welded to a substrate. In this case, the electrode diameter controls the current density. Metallography shows that melting has occurred mostly below the weld interface, except at the weld periphery. Thus the weld is mostly solid state, probably due to the overwhelming influence of the thermal mass of the upper electrode (in this case again made from pure W). Also seen is a small, innocuous pore, as is often found in macroscopic resistance spot welds.
3.7
Summary and conclusions
In this chapter the point has been made that much conventional knowledge, while useful as a guide, will need to be reconsidered for micro scale fusion joining. While there are no insurmountable barriers to prevent joining from being successful at the micro fusion level, all the usual procedures and tools need to be attended to with new insight. Some phenomena, while being important at the macro scale become much more critical (i.e. surface tension, part fixturing) at the micro scale. Hopefully a few guidelines to achieving this transition were provided.
3.14 Resistance spot welds of 25 µm thick PH 17-7 ribbon to 304L case. Left: top view of two welds, right: metallographic cross section.
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3.8
Acknowledgements
Sandia is a multiprogram laboratory operated by Sandia Corporation, a Lockheed Martin Company, for the United States Department of Energy’s National Nuclear Security Administration under contract DE-AC0494AL85000. We would also like to thank Dr R.C. Dykhuizen and Mr D.O. MacCallum, both of Sandia National Laboratories for reviewing the text.
3.9
References
1. Khadraw, A.F., Othman, A., Al-Nimr, M.A., Transient free convection in a vertical microchannel as described by the hyperbolic heat conduction model, International Journal of Thermophysics, 26, 2005, pp. 905–918. 2. Shewmon, P.G., Diffusion in Solids, McGraw-Hill, NY, 1963, p. 8. 3. Metals Handbook, Volume 1, 9th Edition, ASM, Metals Park, OH, pp. 145–151 (approximate average value for 1008 steel over temperature range ambient – 1200 °C). 4. Carslaw, H.S., Jaeger, J.C., Conduction of Heat in Solids, 2nd Edition, Oxford Science Publications, Oxford, 1959, p. 112. 5. MatLab V7.01, the Mathworks, Natick, MA, 2005. 6. Olowinsky, A., Kramer, T., Dumont, N., Hanebuth, H., New applications of laser beam micro welding, Proceedings of ICALEO, 2001. 7. Kelley, K.K., Contributions to the Data on Theoretical Metallurgy, XIII, Bulletin 582, U.S. Bureau of Mines, USGPO, 1960. 8. Stacey, F.D., Physical properties of the Earth’s core, Surveys in Geophysics, 1, 1972, pp. 99–119. 9. Kurz, W., Fisher, D.J., Fundamentals of Solidification, Trans Tech Publications Ltd, Switzerland, 1986, p. 240. 10. Chien, C.-S., Kannaty-Asibu Jr., E., Displacement measurement using a fiber optic sensor in resistance spot welding, Proceedings of 5th International Conference on Trends in Welding Research, ASM, 1999, p. 622. 11. Anisimov S.I., Vaporization of metal absorbing laser radiation, Sov. Phys. – JETP, 27, 1968, p. 182. 12. Anisimov S.I., Khokhlov V.A., Instabilities in Laser-Matter Interactions, Boca Raton, FL, CRC Press, 1995. 13. Iida, T., Guthrie, R.I.L., The Physical Properties of Liquid Metals, Clarendon, Oxford, 1988. 14. Lynch, C.T., CRC Handbook of Materials Science, Vol. 3, CRC Press, Boca Raton, 1974. 15. Samsonov, G.V., Oxide Handbook, IFI/Plenum Press, New York, 1973. 16. Kohl, W.H., Handbook of Materials and Techniques for Vacuum Devices, Reinhold Publishing, New York, 1967, p. 100. 17. Zhao, H., DebRoy, T., Weld metal composition change during conduction mode laser welding of Aluminum alloy 5182, Metallurgical and Materials Transactions B, 32B, 2001, pp. 163–172. 18. Sahoo, P., Collur, M.M., DebRoy, T., Effects of oxygen and sulfur on alloying element vaporization rates during laser welding, Metallurgical Transactions B, 19B, 1988, pp. 967–972.
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19. Afanas’ev, Yu. V., Krokhin, O.N., Vaporization of matter exposed to laser emission, Soviet Physics-JETP, 25, 1967, pp. 639–645. 20. Batanov, V.A., Bunkin, F.V., Prokhorov, A.M., Fedorov, V.B., Evaporation of metallic targets caused by intense optical radiation, Soviet Physics-JETP, 36, 1973, pp. 311– 322. 21. Kelly, R., Miotello, A., Does normal boiling exist due to laser-pulse or ion bombardment, Journal of Applied Physics, 87, 2000, pp. 3177–3179. 22. Semak, V.V., Damkroger, B., Kempka, S., Temporal evolution of the temperature field in the beam interaction zone during laser material processing, Journal of Physics D: Applied Physics, 32, 1999, pp. 1819–1825. 23. Semak, V.V., Knorovsky, G.A., MacCallum, D.O., On the possibility of microwelding with laser beams, Journal of Physics D: Applied Physics, 36, 2003, pp. 2170–2174. 24. Knorovsky, G.A., MacCallum, D.O., Recoil force measurements during pulsed Nd:YAG laser spot welds, Proceedings of ICALEO03, Laser Institute of America, 2003. 25. Kanouff, M.P., internal memoranda, Sandia National Laboratories, (a) April 16, 1998, (b) March 13, 2003. (GOMA is a full-Newton FEM program for free and moving boundary problems with coupled fluid/solid momentum, energy, mass and chemical species transport described in ‘User’s Guide’, Sandia National Labs report SAND97-2404, Sandia National Labs, Albuquerque, NM, September 1998). 26. Fabbro, R., Chouf, K., Sabatier, L., Coste, F., Dynamical interpretation of deep penetration CW laser welding, Proc. of ICALEO’98, November 16–19, 1998, Orlando, FL, pp. 179–186. 27. Lancaster, J.F., The Physics of Welding 2nd Edition, Pergamon Press, Oxford 1986. 28. Sahoo, P., Debroy, T., McNallan, M.J., Surface tension of binary metal-surface active solute systems under conditions relevant to welding metallurgy, Metallurgical Transactions B, 19B, 1988, pp. 483–491. 29. Gottifredi, J.C., Jameson, G.J., The suppression of wind-generated waves by a surface film, Journal of Fluid Mechanics, 32(3), 1968, pp. 609. 30. Semak, V.V., Knorovsky, G.A., MacCallum, D.O., Effect of surface tension on melt pool dynamics during laser pulse interaction, Journal of Physics D: Applied Physics, 39, 2006, pp. 590–595. 31. Personal communication, 2005, Holm, E.A., Sandia National Laboratories. These calculations were performed using The Surface Evolver software, provided by Brakke, K.E., available from www.susqu.edu/facstaff/b/brakke/evolver/ 32. Essers, W.G., Walter, R., Some aspects of the penetration mechanisms in MIG welding, in Arc Physics and Weld Pool Behaviour, the Welding Institute, Cambridge, UK, 1979. 33. Grujicic, M., Hu, Y., Fadel, G.M., Keicher, D.M., Optimization of the LENS rapid fabrication process for in-flight melting of feed powder, Journal of Materials Synthesis and Processing, 9, 2001, pp. 223–233. 34. Burleigh, T.D., Eagar, T.W., Measurement of the force exerted by a welding arc, Metallurgical Transactions A., 14A, 1983, pp. 1223–1224. 35. Woods, R.A., Milner, D.R., Motion in the weld pool in arc welding, Welding Journal, 50, 1971, pp. 163s–173s. 36. Savage, W.F., Nippes, E.F., Zanner, F.J., Determination of GTA weld-puddle configurations by impulse decanting, Welding Journal, 57, 1978, pp. 201s–210s. 37. Renwick, R.J., Richardson, R.W., Experimental investigation of GTA weld pool oscillations, Welding Journal, 62, 1983, pp. 29s–35s. 38. Ki, H., Mohanty, P.S., Mazumdar, J., Modeling of laser keyhole welding: Part II.
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39. 40.
41. 42. 43.
44. 45. 46. 47.
48.
49.
50. 51. 52. 53. 54. 55. 56.
Microjoining and nanojoining Simulation of keyhole evolution, velocity, temperature profile, and experimental verification, Metallurgical and Materials Transactions A, 33A, 2002, pp. 1831– 1842. Matsunawa, A., Keyhole dynamics in laser welding, Technical Report, Lecture Note from a Course given at ICALEO 99, San Diego, CA, 1999 (cited in reference 36). Walsh, D.W., Savage, W.F., The mechanism of minor element interaction in autogenous weld pools, Advances in Welding Science and Technology Conference Proceedings, TWR ’86, ASM International, 1986, p. 59. Szekely, J., Fluid Flow Phenomena in Metals Processing, Academic Press, New York, 1979, pp. 16, 19. Lee, S.Y., Na, S.J., A numerical analysis of a stationary gas tungsten welding arc considering various electrode angles, Welding Journal, 75, 1996, pp. 269s–279s. Matsunawa, A, Yokoya, S, Fluid flow and its effect on penetration shape in stationary arc welds, Recent Trends in Welding Science and Technology, ASM International, 1990, pp. 31–35. Szekely, J., op. cit. pp. 401. Haidar, J., Lowke, J.J., Predictions of metal droplet formation in arc welding, Journal of Physics D, Applied Physics, 29, 1996, pp. 2951–2960. Halliday, D., Resnick, R., Physics, Part II, Wiley & Sons, New York, 1962, pp. 848, 854. Knorovsky, G.A., Overlooked fundamentals of resistance welding, The Metal Science of Joining, The Minerals, Metals and Materials Society, Warrendale, PA, 1992, pp. 249–255. Kijken, D.K., Hoving, W., De Hosson, J.Th.M., Laser penetration spike welding: A microlaser welding technique enabling novel product designs and constructions, Journal of Laser Applications, 15, 2003, pp. 11–18. Fukumoto, S., Zhou, Y., Mechanism of resistance microwelding of crossed fine nickel wires, Metallurgical and Materials Transactions A, 35A, 2004, pp. 3165– 3176. Elmer, J.W., The Influence of Cooling Rate on the Microstructure of Stainless Steel Alloys, Sc.D. thesis, M.I.T., September 1988. Kurz, W., Fisher, D.J., op. cit., p. 54. ibid., p. 57. ibid., p. 51. He, X., Elmer, J.W., DebRoy, T., Heat transfer and fluid flow in microwelding, Journal of Applied Physics, 97, 2005, 084909. Kurz, W., Fisher, D.J., op. cit., p. 61. Olowinsky, A., Klages, K., Gedicke, J., SHADOW® a new welding technique – basics and applications, Proceedings of the SPIE, 5662, 2004, pp. 291–299.
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4 Modeling of solid state bonding Y T A K A H A S H I, Osaka University, Japan
4.1
Introduction
Solid state microjoining methods are used to produce the interconnections between integrated circuit (IC) chips and lead frames (or electric circuit on substrate) in electronic packages. These methods include tape automated bonding (TAB), wire bonding and flip-chip bonding (Takahashi, 1999). TAB is used in bumpless inner lead joint formation (Takahashi et al., 1999a,b). Wire bonding has become a popular process because of its flexibility, which can accommodate changes in microelectronic circuit design. Flip-chip bonding was developed for high density chip scale packaging (CSP). Solid state bonding processes require a certain amount of plastic deformation. Oxide surface film (or organic contaminant film) on the bonding surface breaks down and is dispersed across the bond interface by the plastic deformation which occurs during bonding. The breaking down process is related to the extension of the original bonding surface (Bay, 1981). Bay gave a formula for the relationship between interface extension and bond strength when the bond interface uniformly deforms and extends. However, solid state microjoining methods such as wire bonding and bump bonding do not exhibit uniform deformation. It is also rather difficult to model the breaking down behavior of thin oxide film during bonding. Therefore it is necessary to simulate the plastic deformation process of wire, lead or bump to understand the interfacial extension behavior during bonding, by modeling the deformation process around the bond interface. This chapter discusses modeling and simulation of the plastic deformation process during solid state microjoining.
4.2
Viscoplastic deformation model and interfacial deformation
Special, rate-sensitive materials such as gold, copper and aluminum exhibit viscoplastic behavior in which the flow stress (effective yield stress) depends on the strain rate. Viscoplastic behavior is observed when the materials 91 WPNL2204
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are heated. The steady state equivalent strain rate ε˙ s is generally expressed by D bG σ n Q exp – ε˙ s = Ao o , kT G RT
(1)
where Ao is the non-dimensional constant, Do the frequency factor, n the stress exponent, b Burger’s vector, k Boltzmann’s constant, G the shear modulus, T the absolute temperature, Q the activation energy of viscoplastic deformation, R the gas constant and σ the equivalent stress. If the transient behavior of work hardening has to be taken into account for the short process of micro-bonding, the equivalent strain rate ε˙ is given by D bG σ n Q ε˙ = Ao [1 + B exp (– C ε )] o exp – , kT G RT
(2)
where B and C are non-dimensional constants and ε is the equivalent strain due to plastic deformation. When ε is large enough, equation (2) is nearly equal to equation (1). The material constant for gold and aluminum is given in Table 4.1 (Takahashi and Inoue, 2002). From equations (1) and (2), it is clear that the time for the viscoplastic deformation must be taken into consideration. Takahashi et al. (1993) proposed modeling viscoplastic adhering processes using a finite element method (FEM). Non-linear, eight-node isoparametric elements with four Lagrange multipliers were adopted for analyzing large deformation processes. The finite element method was applied to modeling the contact process between two joining surfaces with surface waviness. Void shrinkage due to viscoplastic deformation was simulated. It was proposed that the contact process between the two joining surfaces was controlled by time-dependent viscoplastic deformation, Table 4.1 Material constants of pure gold and aluminum Name
Symbol
Au
Al
Al-Si-Cu
Unit
Frequency factor Constant Constant Constant Burger’s vector Activation energy Stress exponent Melting temperature Shear modulus
Do Ao B C b Q n Tm
0.1 ×10–4 8.38 ×107 1.50 ×102 1.00 ×102 2.88 ×10–10 132.9 6.57 1336
0.1 ×10–4 2.28 ×108 5.60 ×102 1.00 ×102 2.86 ×10–10 150.7 4.40 933
0.1 ×10–4 1.00 ×107 5.60 ×102 1.00 ×102 2.86 ×10–10 150.7 4.80 ~933
m2 s –1
where
aG bG cG
G(T ) = aG(T/Tm)2+bG(T/Tm) + cG –9.9 ×109 –6.2 ×109 3.1 ×1010
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m kJ mol–1 K Nm–2
–9.4 ×109 –5.3 ×109 2.9 ×1010
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93
regardless of uniform or non-uniform deformation or bulk constraint conditions. The contact process is naturally affected by boundary conditions, i.e., the bump shape or the size ratio between wire diameter and pad thickness. The interfacial deformation is significantly related to the contact (bonded area) growth process. Time-dependent plastic deformation is important for modeling interfacial deformation during microjoining.
4.3
Thermocompression bonding
Large plastic deformation occurs during solid state microjoining, such as wire bonding, regardless of whether or not ultrasonic vibration is applied. Ultrasonic vibrations enhance plastic deformation and frictional micro-slip (Schwizer et al. 1999), but thermocompression bonding without ultrasonic vibration is very useful in helping our understanding of interfacial extension behavior during bonding, the effect of the size ratio (pad thickness effect) and so on.
4.3.1
Inner lead bonding in TAB
Model and calculation procedure TAB technology is a useful joining method for IC packaging (Takahashi et al., 1999a). TAB is used for inner lead bonding, between the fine leads on the tape and the pads on the IC. When solid state TAB (gang bonding) is carried out, fine leads plated with gold are bonded to gold-plated bumps or Al electric pads. In modeling, the lead bonding is simplified as illustrated in Fig. 4.1(a) and (b), where (a) is the simple model and (b) the mesh pattern of finite elements at section Q in Fig. 4.1(a). Both lead and pad are made of gold. In this model, the tool interface FE and the bond interface GD are assumed to be fixed between lead and pad, because of thermocompression bonding. It is also assumed that the substrate OB and the bonding tool FE are too hard to deform, and the temperature of lead and pad is uniformly kept at 573 K. This is the case when M = 1. The parameter M is defined by M=
Aop exp(– Q p / RTp ) , AoL exp(– Q L / RTL )
(3)
where Aop and AoL are non-dimensional constant Ao for pad and lead, respectively. Qp and QL are the activation energy of viscoplastic deformation for pad and lead, respectively. Tp and TL are the mean temperature of pad and lead, respectively. The parameter M is useful in discussing the effect of mechanical properties depending on temperature differences between pad and lead (Takahashi et al., 1999b). The size ratio δo /Ho is also defined, where Ho is the initial height of the lead and δo is the initial thickness of the
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Bonding tool
Load
Q Lead
Pad
Substrate
(a)
Constant load
y Bonding tool
Lead
E
Ho
F
Ho /2 D
C
O
A
B x
δo
Pad
Ho /2
G
(b)
4.1 Simple model of inner lead bonding (a) illustration of bonding tool, inner lead and pad on substrate, and (b) the initial mesh devision of the cross section of Q in Fig. 4.1(a).
pad as shown in Fig. 4.1(b). The lead reduction (compression ratio) ∆H/Ho is also introduced, where ∆H = Ho – H and H is the lead height after bonding. Figure 4.2 schematically illustrates finite elements and nodal points at the interface between lead and pad in (a) and also defines the segment extension εseg in (b), which is expressed by εseg = (l – lo)/lo,
(4)
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Segment number, Ns Lead 1
2
3
4
5
6
7
8
9
10
G
D
Pad
(a) εseg =
Initial contact interface
lo
l – lo lo
l
Lead Pad
Before deformation
After deformation (b)
4.2 Schematic illustration of bond interface between lead and pad (a) definition of segment number at bond interface, and (b) definition of segment extension (strain).
where lo is the initial segment length between the neighboring nodes and l is the segment length after deformation. The segment extension in FEM means the local interfacial extension. Local interfacial extension is a factor in controlling the bond strength of pressure welding such as solid state microjoining in the absence of ultrasonic vibration. Interfacial extension is related to surface exposure ψ, representing the degree of solid state pressure welding between metal surfaces and given by ψ = εseg(1 – Gd)/(1 + εseg),
(5)
where Gd (= 0 ~ 1) is the ductility of surface contamination or oxide film (Zhang and Bay, 1997). Therefore, segment extension is an important parameter for bond formation. Figure 4.3 is the flow chart for computer programming. At first, proper displacement rates { u˙ } are assigned to all nodal points as initial values { u˙ }old. The bonding load Wb (per unit length of lead) applied by the bonding tool in the y direction represents the boundary condition for summation ∑ Fnodal of nodal force in the y direction at the tool interface FE in Fig. 4.1(b). A certain displacement rate ( u˙ x , u˙ y ) = (u, v), which has a proper value of v, is initially given in the y direction in the tool interface FE in Fig. 4.1(b), and the value of v is fixed in each calculation step (loop). When the tool interface does not slide, u = 0. The displacement rate (value of v) of
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t=0 . Assumption of a proper displacement rate vector, {u} Calculation of element stiffness equations Formation of overall stiffness matrix . Calculation of perturbation, ∆{u} . {u}new = . . {u}old + ω1∆{u}
. ∆{u} < 5.0 × 10–4 . {u}
No
Yes Summation of nodal force at bonding tool, Σ Fnodal
0.999 ≤
Σ Fnodal Load
≤ 1.001
No
{u}new = ω2 {u}old
Yes Calculation of stress and strain rate Time increment ∆ t
t = t + ∆t Renewal of coordinates of nodal points No
t > tend Yes Output of result End
4.3 Flow chart for computer programming.
nodal points as the boundary condition is fixed in each calculation step, and the displacement rate at the fixed interface OB on the substrate in Fig. 4.1(b) is always given zero as the boundary condition. The proper displacement rates { u˙ }old given at other nodal points should be converged to the true rate vector { u˙ }new in each step by the perturbation method (Takahashi et al., 1993). The summation ∑ Fnodal of the nodal force in the y direction at the tool
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interface FE is usually not equal to the load Wb at the initial stage in each calculation step. The summation ∑ Fnodal has to agree with the bonding load Wb. In other words, the value of u˙ y = v at the tool interface is changed so that ∑ Fnodal is equal to Wb. The numerical solution is iteratively converged until the perturbation ∆{ u˙ } is much less than { u˙ }. In each loop, ∑ Fnodal is finally equal to the bonding load Wb in each step of time t. The time increment ∆t is properly changed from 10–10 to 10–2 s. The coefficient ω2 is greater than unity if ∑ Fnodal is less than the load Wb. When the two relaxation coefficients ω1 and ω2 are changed, the coefficient ω2 is less than unity if ∑ Fnodal is greater than the load Wb. The coefficient ω1 is in the range of 0.1 ~ 0.3 in the present simulation, although usually greater than unity. The bonding pressure P is defined by P = Wb /w, where w is the initial width of the lead. Numerical simulations Figure 4.4 shows the calculated effect of pad thickness on the distribution of interfacial extension. As FEM is used, the interfacial extension is discussed in terms of the distribution of the segment extension εseg and segment number Ns, shown in Fig. 4.2(b). As seen in Fig. 4.4, when the size ratio δo /Ho is unity, interfacial extension occurs easily at the central area of the bonding interface, but becomes smaller as δo /Ho decreases. The maximum value of εseg exists at Ns = 8 and the minimum value exists at the edge (Ns = 9 and 10) of the bond interface, i.e., the edge tends to shrink when the pad is thin enough. This is due to the constraining effect of the pad and the hard substrate. The peak of the segment extension appears inside the edge and εseg reaches 5% at Ns = 8 for ∆H/Ho = 10%, even if δo/Ho = 0.1. This implies that adhesion is more easily produced at the inside of the lead edge.
Segment extension, εseg (%)
10 8
δo /Ho = 1.0
6 4
δo /Ho = 0.3
2
δo /Ho = 0.1
0 –2 –4
M = 1.0 ∆H /Ho ≈ 10 (%) 1
2
3
P = 392 (MPa) T = 573 (K)
4 5 6 7 8 Segment number, Ns
9
10
4.4 Effect of size effect (δo /Ho) on interfacial (segment) extension.
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Viscoplastic deformation is controlled by the equivalent stress as indicated in Equation (2) (Takahashi et al., 1999a). Figure 4.5 shows the equivalent stress distributions on cross-section Q of Fig. 4.1(a). As seen in Fig. 4.5, the equivalent stress distribution is affected by pad thickness (the size ratio of δo/Ho). The equivalent stress is largest at the center of the lead and the perimeter of the interface of lead and pad. When the size ratio δo /Ho is large enough, like that of Fig. 4.5(a), the stress at the perimeter is dispersed into the pad and does not place a concentrated load on the substrate. On the other
P = 392 MPa T = 573 K ∆H/Ho = 10%
350 ≤ σ 300 ≤ σ ≤ 350 250 ≤ σ ≤ 300 200 ≤ σ ≤ 250 150 ≤ σ ≤ 200 100 ≤ σ ≤ 150 50 ≤ σ ≤ 100 0 ≤ σ ≤ 50 (MPa)
(a)
(b)
4.5 Stress distribution in cross section of lead and pad (a) δo /Ho = 0.5 and (b) δo /Ho = 0.1.
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hand, as the size ratio decreases (the pad becomes thinner), the stress is concentrated under the edge of the lead and can damage the substrate. Figure 4.6 shows the effect of M on interfacial (segment) extension. If the lead temperature TL is 573 K, the pad temperature Tp is 491 K at M = 10–2, Tp = 573 K at M = 1, Tp = 678 K at M = 102, and Tp = 740 K at M = 104. As can be seen in Fig. 4.6, interfacial extension is not achieved at M = 10–2 and this implies that direct bumpless lead bonding is very difficult for hard metals (thin and cool pad on hard substrate). However, as M increases, i.e., the pad heats up, the lead/pad interface begins to extend easily. When M ≥ 102, the segment extension is greater than 10% at Ns = 7~8 in spite of ∆H/ Ho = 10%. The extension of Ns = 9~10 decreases and the interface under the lead edge (i.e. the corner) is compressed. Even if M increases, the bonding interface at the lead edge tends to shrink due to the swelling of the pad. This swelling becomes striking under heating because the pad becomes soft. It should be noted that the interface extension is non-uniform and shrinks at the lead edge. Figure 4.7 shows the calculated distributions of the equivalent stresses in the cross-section of lead and pad. When the pad temperature is low (M = 10–2), the equivalent stress is concentrated in the shape of a cross pattern (Fig. 4.7(a)) within the lead and then flows into the pad under both sides of the lead edges, at the corner junctions. On the other hand, in Fig. 4.7(b) for M = 104, the black area, where σ > 300 MPa, is widely distributed within the lead. Because the pad is heated and soft, the flow stress within the pad is much lower than that of the lead. So, the pad deforms more easily than the lead and swells on either side of the lead edge. The forces are balanced between lead and pad (in the y direction) even though the equivalent stress is lower in the pad than in the lead. The numerical calculations predict
Segment extension, εseg (%)
25 20
P = 392 MPa T = 573 K
M = 104
15
M = 102
10
M = 1.0 5 0 –5 –10
∆H /Ho ≈ 10% δo /Ho = 0.1 1
2
3
M = 10–2
4 5 6 7 8 Segment number, Ns
9
4.6 Effect of M value on interfacial extension.
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P = 392 MPa T = 573 K δo /Ho = 0.1 ∆H/Ho ≈ 10%
(a)
350 ≤ σ 300 ≤ σ ≤ 350 250 ≤ σ ≤ 300 200 ≤ σ ≤ 250 150 ≤ σ ≤ 200 100 ≤ σ ≤ 150 50 ≤ σ ≤ 100 0 ≤ σ ≤ 50 (MPa)
(b)
4.7 Stress distribution in cross section of lead and pad (a) M = 10–2 and (b) M = 104.
accurately the effects of pad thickness and mechanical properties (Takahashi et al., 1999b).
4.3.2
Wire bonding
Wire bonding is an important microjoining technique, even though the bonding mechanism has not been perfectly solved yet. It is very difficult to clarify the slip and folding mechanism during wire bonding. The actual thermocompression bonds or thermosonic bonds often exhibit perimeter bond formation as the bonding load increases (or when wire deformation occurs rapidly at high temperatures), and a doughnut configuration with bonding
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near the periphery and no bond in the central zone is sometimes observed under high pressure conditions (Winchell and Berg, 1978). This is due to the frictional slip behavior, which depends on the bonding load and ultrasonic vibration (Lum et al., 2005). If the ultrasonic power is small in relation to the bonding load, frictional slip does not occur. Therefore, the bond interface may hardly slide after the wire touches (folds) to the pad surface during thermocompression bonding without ultrasonic vibration. Mayer et al. (2000) reported that metallic compounds begin to be produced from the central area and this extends to the periphery. Frictional slip plays a very important role in adhesion in Au ball bonding to Al pads. Lum et al. (2005) also suggested that the metallic compound formation at the periphery area is due to microslip annulus, which changes to gross slip when ultrasonic power increases, resulting in overall adhesion. Because frictional slip interacts with metallic compound formation, it is rather difficult to model the actual frictional slip behavior precisely. Liu et al. (2004) performed a numerical simulation of thermosonic ball bonding and discussed the impact of ultrasonic vibration on the wire bonding process, indicating that ultrasonic power enhances strain rate by raising the effective stress and facilitating wire deformation. On the other hand, in wedge bonding, ultrasonic vibration is usually applied in the longitudinal direction of the wire, creating great complexity for two dimensional modeling. The modeling and simulation of perimeter (periphery) bond formation in thermocompression wire bonding without longitudinal vibration is easier to understand, and this section, therefore reviews modeling and simulation of wire bonding without ultrasonic vibration. Pad thickness influences the wire thermocompression bonding process (Ishizaka et al., 1977). As the pad thickness increases, thermocompression bonding becomes possible in the central zone of gold ball bonds to Al pads. The effect of pad thickness on the bond formation should be taken into account in modeling thermocompression bonding. Models and computation procedure Saeki et al. (1991) reported a deformation (FEM) analysis of an Au wire ball and Al pad, using a rigid plastic FEM (axi-symmetric approximation). They simulated numerically a large deformation process in the Au ball due to the pressure of the capillary (bonding tool), but did not give any clear reasons against the perimeter bond from the viewpoint of interfacial deformation. Takahashi and Inoue (2002) proposed a simple model of wire bonding to describe the pad thickness effect on wire bonding and the interface extension effect on perimeter bonding behavior. The three points under review here are (i) the size ratio δo /Ho for flat tools or δ o / H o′ for grooved tools as illustrated in Fig. 4.8, (ii) the M value in Equation (3), which expresses differences in
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y Flat tool E
Ho
Wire
Ho /2
F
D
Ho
B
δo
C Pad O
A
x
(a) V groove tool Constant load
y E
G
F
D
Ho′
Ho
Wire
Ho /2
Ho δo
C
B
Pad O
A
x
(b)
4.8 Mesh pattern in cross section of wire and pad (a) for flat tool and (b) V groove tool.
the mechanical properties between wire and pad, and (iii) the tool shape effect, where δo is the pad thickness and Ho is the initial height (diameter) of wire, H o′ is the height as shown in Fig. 4.8 for grooved tools. Figure 4.8 illustrates the mesh pattern of the cross-section of wire and pad. The initial diameter of wire is 10 µm. Because of the plane symmetry (plane strain approximation), only the right-hand side is meshed. It is assumed that the interfaces between pad and wire or between wire and tool are fixed after the wire surface touches to the pad or the tool surface. The segment extension is
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referred to in Equation (4). Figure 4.9 illustrates the segment numbers Ns1, Ns2, and Ns3. In the wire bonding model, segments Ns1 = 1 and Ns3 = 1~18 are the interface from the beginning, and segments Ns1 ≥ 2 and Ns2 ≥ 2 are surfaces under initial conditions. A new interface is produced by the folding of wire surface (Ns1 = 2~11) to the pad surface Ns2 = 2~11. Figure 4.10 illustrates the method for obtaining the new node. When folding occurs, it is necessary to remove or add nodes in the finite element mesh. In Fig. 4.10, node 1 is the junction of the interface between wire and pad. Node 2 is a point on the wire surface and node 3 and node 4 are points on the pad surface. The initial mesh pattern in Fig. 4.9 is determined so that nodes 2′ and 3′ could meet at nearly identical coordinates as seen in Fig. 4.10. The new node is given as the midpoint between node 2′ and node 3′ unless the positions of nodes 2′ and 3′ are equal in the overall coordinate. This method avoids the complication of remeshing overall nodes and transferring data with respect to the displacement
Wire
C
6 7
5
4 1 2 3
8
9
11 10
Segment number, Ns1
(a) 1
2
3
4
5
6
7 8 9 10
Segment number, Ns2
Pad
1
2
3
4
5
6
11
7 8 9 10
Substrate
11
12
13 14 15 16 17 18
Segment number, Ns3 (b)
4.9 Definition of segment number (a) for wire surface and (b) for pad surface and interface on substrate.
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Wire
Pad
Folding
Wire
2
3
Node
4′
2′ 3′
1 1′
4
1′ Wire
4′
Pad
Pad
New node
4.10 Folding process of wire surface to pad surface. Black circles are the nodal points. Nodal point 1 is the junction between wire and pad. Nodal point 2 is on wire surface. Nodal points 3 and 4 are on pad surface. White points 1′, 2′, 3′ and 4′ are the position of the points 1, 2, 3, and 4 after the time increment ∆t given properly, respectively. The folding of 2′ and 3′ is judged by reading the coordinates. A new node is added as a new junction between wire and pad, if 2′ folds to 3′.
rate vector and the position of nodes, making the FEM solutions easier to obtain by increasing bonding time. The calculation is repeated until segment number Ns1 =11 of the wire surface folds to segment number Ns2 = 11 of the pad surface. The calculating method for wire bonding has been detailed elsewhere (Takahashi et al., 1996; Takahashi and Inoue, 2002) and is the same as that introduced in the section on lead bonding, except for the folding treatment. The constant (bonding) load Wb per unit length of wire is defined by P × d, where P is the bonding pressure and d is the initial diameter of wire. In the initial stage, a constant load is applied to the segment at point E in Fig. 4.8(a) or at point G in Fig. 4.8(b). The nodal displacement rates {u˙} are then converged to true rate vectors by the perturbation method in each calculation step (time t) so that the summation ∑ Fnodal of nodal force at the tool interface can be equal to Wb. Folding occurs at two places, i.e., there are two kinds of folding. One is concerned with the contact between wire and pad as illustrated in Fig. 4.10, and the other is produced between tool and wire. For wire bonding using a V-grooved tool, as in Fig. 4.8(b), folding occurs at both sides of point G. For tool folding where the wire touches the bonding tool, boundary conditions should be changed to incorporate nodal points fixed to the bonding tool. After the true displacement rate vector is solved at each calculation step at time t, the correct displacement rate vector ( u˙ x , u˙ y )
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is obtained for each nodal point and the coordinate of the nodal point is (x, y) at t = t. The coordinate (x, y) at t = t + ∆t is updated by (x, y) at t = t + ∆t = (x, y) at t = t + ( u˙ x , u˙ y ) ∆t. The two kinds of folding are judged and remeshing and the new boundary conditions updated. The numerical simulation for the next step of t = t + ∆t is carried out. The ending time tend is given by the time when the wire reduction of ∆H/Ho or ∆H′/Ho is greater than 50%, or the nodes at segment Ns1 = 11 on the wire surface are folded to the pad surface. The wire reduction (compression ratio) is defined in Fig. 4.11. The wire reduction ∆ H ′ / H o′ is used for wire bonding by V-grooved tools. In the thermocompression bonding model, bonding pressure P and bonding temperature T in numerical simulations were set higher than those of actual thermosonic wire bonding, because ultrasonic vibration was not applied in the simulation, i.e., P = 980 MPa for flat tool and P = 1568 MPa for Vgrooved tools, and T = 573 K was given for easily deforming wire. Numerical simulations As the size ratio δo /Ho increases, the pad deforms easily and is indented by the wire, because the constraint effect of the substrate on the pad decreases with increased size ratio. If the pad is thin (δo /Ho is small enough under the condition of M = 1), the equivalent stress is concentrated at the periphery inside both edges of the wire after folding, but is insufficient for interfacial deformation in the central zone. On the other hand, if the size ratio δo /Ho is large enough, the stress becomes high at both the central zone and at the ∆H = Ho – H
∆H
H
Ho
Flat tool
Deformation
(a) Reduction: ∆H / Ho ∆H ′ = Ho′ – H ′
∆H ′
H
Ho′
Ho
V groove tool
Deformation
(b) Reduction: ∆H ′/ Ho′
4.11 Schematic illustration for definition of wire reduction (a) for flat tool and (b) for groove tool.
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periphery (Takahashi and Inoue, 2002). It follows that interfacial extension becomes larger in the central zone of the bond area as pad thickness increases. The size ratio also affects the interfacial extension in wire bonding. As δo /Ho increases, the interfacial (segment) extension increases in the central zone. The pad surface exhibits a negative extension (compression) at the periphery if the pad is thin, even though the wire surface is extended. The effect of pad hardness (mechanical properties) is also very important for understanding interfacial deformation. As the pad becomes hard or M decreases, the wire reduction rate decreases and higher flow stress is needed for wire deformation, because the hard pad constrains the wire. The value of M also influences stress distribution on the interface between pad and substrate, as suggested by stress distribution changing with the segment number Ns3. If the pad is very thin and M is lower than unity, the stress is concentrated under the edges of the wire periphery and this makes it easy to form the periphery bond. As the wire diameter increases from 10 µm to 500 µm, the pad thickness δo increases 1 µm to 50 µm under the same conditions of size ratio δo /Ho = 1. In actual Al wire bonding, the pad thickness is a few microns even if the wire diameter is 300 µm. It follows then that stress concentration occurs in the substrate and damages the substrate or IC chip. When the pad becomes harder with precipitates, damage is often observed (Onuki and Koizumi, 1996). Load control (ramp loading) with ultrasonic vibration is necessary to facilitate bond formation in the central area for thin and hard pads, i.e., the slip and folding mechanism is very important. Figures 4.12 and 4.13 show numerical simulations of wire and pad deformation processes. The right-hand sides of the figures express the deformation patterns of the finite element mesh, and the left-hand sides give the displacement rate vectors for each nodal point. Figure 4.12 represents flat tools and Fig. 4.13 represents V-grooved tools. The initial wire diameter is 10 µm and M = 1, i.e., the materials of wire and pad are gold. Numerical values indicated at the top of each figure in Fig. 4.12 express the y components of the displacement rate vectors (length of arrow). Numerical values shown on the left-hand side of each figure in Fig. 4.13 express the value of the displacement rate at the first contact point G with the grooved tool. The center of the wire is compressed greatly in the flat tool. In the grooved tool, the compression is not so striking and the wire height not so changed from Ho because the upper side of wire is squeezed into the gap between the wire and the grooved tool. The constraint of the grooved tool significantly changes the overall wire deformation pattern and reduces the deformation rate. Figure 4.14 shows the changes in segment extension by segment number. Figure 4.14 suggests that the grooved tool facilitates interfacial extension if ∆ H ′ / H o′ is kept to 36% between flat and grooved tools. The grooved tool produces center bond formation more easily than the flat tool, because the
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107
By flat tool P = 980 MPa T = 573 K
t = 2.00 × 10–5 s (a) 6.39 × 10–3 (m/s)
t = 9.35 × 10–5 s (b) 4.86 × 10–5 (m/s)
t = 7.85 × 10–3 s
(c)
4.12 Numerical results of wire bonding by flat tool, with respect to the deformed patterns changing with the wire reduction (P = 980 MPa, T = 573 K) (a) ∆H/Ho = 0.0%, (b) ∆H/Ho = 21.3% and (c) ∆H/Ho = 39.8%.
segment extension becomes larger as it approaches the central zone. Negative extension is observed at Ns2 = 9~11 at the side of the pad as seen in Fig. 4.14(b). This implies that folding of the lateral surface is involved in the extension of wire surface and the shrinkage of pad surface in the periphery area, and thus the folding behavior is very complex.
4.4
Thermosonic bonding
Ultrasonic vibration can facilitate solid state microbonding. Interface slip and deformation can be enhanced by ultrasonic vibration and the bonding
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2.56 (m/s)
t = 1.00 × 10–20 s (a)
1.15 × 10–4 (m/s)
t = 1.86 × 10–3 s
(b)
3.27 × 10–7 (m/s)
t = 3.23 × 10–1 s
(c)
4.13 Numerical results of wire bonding by grooved tool, with respect to the deformed patterns changing with the wire reduction (P = 1568 MPa, T = 573 K) (a) ∆H ′ /H o′ = 0.0%, (b) ∆H ′ /H o′ = 20.4% and (c) ∆H ′ /H o′ = 35.9%.
interface heated up so that fretting or adhering is produced at the bonding interface. It is difficult to simulate frictional slip and interfacial deformation behaviors when ultrasonic vibration is applied. In particular, ultrasonic vibration applied in the longitudinal direction of the wire requires three dimensional analysis, which becomes highly complex. The two dimensional model for
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Segment extension, εs (%)
80 70 60 50
P = 980 (MPa) T = 573 (K) ∆H ′/Ho′ ≈ 36 (%) M = 1.0 δo /Ho = 0.1
Flat tool
40 30 V groove tool
20 10 0 0
1
2
3
4 5 6 7 8 Segment number, Ns1 (a)
9
10
11
12
9
10
11
12
30 Segment extension, εs (%)
V groove tool 20 10
Flat tool
0 –10 –20
P = 980 (MPa) T = 573 (K) ∆H ′/Ho ≈ 36 (%)
–30 0
1
2
3
M = 1.0 δo /Ho = 0.1
4 5 6 7 8 Segment number, Ns2 (b)
4.14 Effect of shape on interfacial extension (a) for wire surface and (b) for pad surface.
gold bump bonding is, therefore, introduced in this section. The model is simple but it is useful for understand the effect of ultrasonic vibration on plastic deformation and frictional slip.
4.4.1
Model of gold bump bonding
Gold bumps on IC chips are gang-bonded in flip-chip bonding (FCB). The gold bumps are bonded to the IC chip in advance, before FCB. Figure 4.15 illustrates the gold bumps at start time t = 0 s, pressed to the pad on the substrate by the bonding tool. It is assumed that the faying interface between the gold bump and the gold pad on the substrate is bonded. The cross-section
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Load Tool
Ultrasonic vibration
IC chip
Pad(Al-Si-Cu) Bump Pad (Au)
Substrate FCB Enlargement
Load + ultrasonic vibration
Tool
IC chip
Al pad Bump
Au pad Substrate Bonding interface
4.15 Illustration of bump bonding (flip chip bonding).
of the central area of the bump and the pads on the chip and the substrate is shaded in Fig. 4.15. Figure 4.16 is the FEM of the shaded area. The model is two dimensional and approximated to the plane strain problem. The pad (Area HIFG) on IC and the pad (Area LABC) on the substrate are, respectively, meshed with two and three layers of finite elements. The interfaces, LB, HG and JE are assumed to be fixed. It is assumed that the bonding interface KD can slip but not be separated.
4.4.2
Description of interfacial micro friction slip
It is assumed that the ultrasonic vibration amplitude Ao is small enough so that each nodal point cannot slip significantly over time increment ∆t, i.e., the slip model treats only micro- or nano-slips, the direction of which rapidly changes following ultrasonic vibration. Slip can occur when |τxy| > µ|σy|, where µ is the friction coefficient, τxy is the shear stress, and σy is the stress in the y direction in Fig. 4.16. It is hard to solve the slip rate self-consistently
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Load + Ultrasonic vibration
y Tool IC chip
G
H I
J
E
F
Pad on IC chip Bump Pad on substrate L K A
D O
x
C B
4.16 Illustration of mesh patterns of finite elements in bump bonding model.
at each step. It is assumed that the value of the micro frictional slip rate is given by x˙ s = αs(| τxy | – µ | σy |), where αs is a constant. The stress distribution at the next step is accommodated by the micro frictional slip, but the frictional slip model cannot be applied to large slip behavior because the stress relaxation at each step is ignored. The friction coefficient is generally expressed as a function consisting of three terms, µo, the summation of slip value and work hardening, i.e., µ = µ o + K1
∫
t
x˙ s dt + K 2 ε ,
(6)
t =0
where µo has two terms depending on local normal stress Pl and temperature T, and can be approximated to µo = ks(asPl + bs)(AsT + Bs), where ks is expressed by ks = ko1exp(–ko2| x˙ s |) + ko3, and ko1, ko2, ko3, as, bs, As, and Bs are constants listed in Table 4.2. For example, the value of µ is 3.47 × 10–1 for Pl = 392 MPa and T = 473 K at t = 0 s. The slip direction is in that of the shear stress. When ultrasonic vibration is applied parallel to the interface between IC chip and bump, the stress σy becomes positive (tensile) at the bonding interface edge opposite to the direction of ultrasonic vibration. In this case, the nodes should be separated, but the node separation procedure due to ultrasonic vibration is extremely difficult to programme. It is, therefore, assumed when σy > 0 at the bonding interface that the node is not separated and the same slip rate is given by x˙ s = αs | τxy |, which is the slip rate at the innermost node showing σy = 0 on the local bonding interface. This is a problem that remains in modeling ultrasonic bump bonding. Ultrasonic
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Value
Unit
αs k01 k02 k03 as bs As Bs K1 K2
2.0 × 10–12 1.0 1.0 × 104 0.005 1.967 × 10–4 1.03 × 10–1 5.37 × 10–3 –6.11x10–1 1.0 × 106 1.0 × 101
m3s–1N
(MPa)–1 K–1 K–1
Note: The constants are for Au/Au bonding interface.
vibration applied parallel to the IC chip surface HG of Fig. 4.16(b) is given by uν = Ao sin ωt. The rate is u˙ v = Ao cos ωt, where Ao is the amplitude and ω is equal to 2πfv, and fv is the frequency of ultrasonic vibration. The time increment ∆t was set as ∆t = Tv /8, where Tv is the period (Tv fv = 1).
4.4.3
Numerical simulations of FCB
Figure 4.17 shows the effect of ultrasonic vibration on the equivalent stress distribution. Figure 4.17 shows thermocompression bonding (a) without ultrasonic vibration (TC), and (b) with ultrasonic vibration (US) of Ao = 10 nm and fv = 60 KHz. The bonding condition is P = 392 MPa and T = 473 K and the size ratio Ho /Xo is unity, where Ho is the initial bump height and Xo is the initial bump width KD in Fig. 4.16. The thickness of the substrate pad is enlarged by 30 times in the drawing, it is actually 1 µm. In this example, the IC chip pad is ignored (0 µm). If ultrasonic vibration is not applied, the equivalent stress distribution is plane symmetrical as shown in Fig. 4.17(a) and the stress value is no more than P = 392 MPa. When ultrasonic vibration is applied normal to the y direction, the stress distribution is not symmetrical and, if the slip is small enough, the stress level is much higher than P = 392 MPa, as seen in Fig. 4.17(b). Referring to Equation (2), this means the increase in strain rate, i.e., the bump deformation is enhanced by applying ultrasonic vibration. In Fig. 4.17, the bump compression rate of US is about 90 times larger than that of TC. Such an enhancement is also simulated for thermosonic ball bonding by Liu et al. (2004). Figure 4.18 shows the slip rate distribution when ultrasonic vibration is applied to the right. Nodal points Ns = 1~25 are on the bonding interface. The frictional slip in TC is much smaller than that in US.
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(a) 900 ≤ σ 800 ≤ σ ≤ 900 700 ≤ σ ≤ 800 600 ≤ σ ≤ 700 500 ≤ σ ≤ 600 400 ≤ σ ≤ 500 300 ≤ σ ≤ 400 200 ≤ σ ≤ 300 σ ≤ 200 (MPa) (b)
4.17 Effect of ultrasonic vibration on the equivalent stress distribution for Barrel bump of Ho/Xo = 1 at t = Tv under the bonding condition of T = 473 K, P = 392 MPa and the ultrasonic vibration conditions are Ao = 10 nm, fv = 60 kHz.
Ultrasonic vibration not only facilitates bump deformation, but also interfacial slip. If ultrasonic vibration is applied to the right, the pressure is concentrated at the right-hand edge and frictional slip does not occur. On the other hand, because of the pressure reaction, tensile stress occurring at the left-hand edge and shear stress therefore occurs in the opposite direction at the left-hand side if the bump is high in relation to its width. So, the slip rate becomes negative on the left-hand side in this simple model, which assumes that separation does not occur. Even if separation does occur, frictional slip at the contact area does not become larger because the pressure increases at the contact interface. It follows that very high ultrasonic power only produces unstable interfacial deformation. The instability in the bump deformation causes unstable, complicated frictional behaviors. Unstable deformation can be restricted by bonding pressure, which decreases frictional slip. Therefore, setting suitable bonding conditions is necessary in successful FCB. Figure 4.18 suggests that if the slip is small enough, it occurs preferentially in the central area of bump bonding. If the vibration is applied to the right,
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Frictional slip rate, xs (10–4 m/s)
With ultrasonic vibration 0.5
0 Without ultrasonic vibration (TC) –0.5
–1.0 1
3
5
7
9 11 13 15 17 19 Number of nodal point
21
23
25
4.18 Interfacial (frictional) slip rate depending on the interface position. Number of nodal pont Ns is set along the interface KD in Fig. 4.16. Point K is correspond to Ns = 1 and point D is to Ns = 25.
the right-hand edge of the bonding interface scarcely slips because of the high pressure concentration. The left-hand interface can slip, being separated. This local separation and slip of the opposite side against the ultrasonic vibration may lead to the periphery sliding when the ultrasonic power is not so high. This is sometimes observed in thermosonic ball bonding because of the capillary shape (Lum et al., 2005). However, the slip usually begins from the central area even in thermosonic ball bonding (Mayer et al., 2000) if the adhesion occurs from the central area (the formation of metallic compounds begins from the central area). Figure 4.18 also suggests that the central area slips more easily. The right edge (Ns = 23~25) is in front of the central area and cannot slide. The central area slips towards the front area. Also, the absolute value of negative slip at the left edge (Ns = 1~3) is less than that of the central zone slip at Ns = 7~19 as seen in Fig. 4.18. This means that the bonding interface shrinks under interfacial slip behavior. As the initial bump height decreases, negative slip at the opposite side does not occur, because the direction of shear stress becomes the same as that of the ultrasonic vibration. It follows that frictional slip may shrink the bonding interface but does not extend it, i.e., the contact area (the interface area between bump and pad) is not increased by frictional slip. The increase in contact area is mainly produced by folding mechanisms due to bump deformation. Frictional slip, however, disperses the surface
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oxide or contamination film and facilitates adhesion (Harman and Albers, 1977). Frictional slip can also heat up the bonding interface, and is necessary to the formation of the central area bonding (Mayer et al., 2000). As ultrasonic power increases in wire ball bonding, gross slip begins to occur (Lum et al., 2005). The transition point from local slip to gross slip depends on the bonding load, the ultrasonic vibration value, bonding tool shape and so on. However, because the chip is flat in FCB, the normal stress distribution changes from that of ball bonding with capillary tools. Frictional slip occurs easily in the central area as seen in Fig. 4.18, if the ultrasonic vibration is not so large. Because the vibration amplitude of the IC chip is much less than that of wire bonding, it is not considered that large gross slip occurs in the FCB process to make adhesions simultaneously between many bumps and pads. Bump deformation accompanying large gross slip is, therefore, not discussed in this chapter, but it is suggested from numerical simulation results that a large gross slip causes a large stress relaxation and that bump deformation is not enhanced by ultrasonic vibration when large gross slip occurs, as Liu et al. (2004) pointed out for thermosonic ball bonding. In other words, the enhancing effect of ultrasonic vibration on bump deformation (see in Fig. 4.17) can be lost under large gross slip. Frictional bonding is a complicated, multiple process that depends on load, temperature, geometry of bonding tool, size ratio of bump and pad thickness and so on. This section discusses bump bonding where only a small slip occurs, but numerical modeling and simulation of more complex slip behavior accompanying interface separation will be necessary for a greater understanding of the thermosonic or ultrasonic bonding mechanism.
4.5
Numerical simulation of lap resistance welding
To model the resistance welding process, a coupled electrical–thermal– mechanical finite element analysis is required. Electric current distribution affects temperature distribution and consequently the elastoplastic deformation process. The deformation changes the electric current flow and hence joules heating, changing the temperature distribution. These electrical thermal and mechanical phenomena should therefore be taken into consideration simultaneously, self-consistently for unsteady state conditions and repeatedly in the time increment ∆t for transient behaviors. As the temperature increases, the deformation of the resistance welding zone becomes larger and electrical thermal conditions should be modified. A new field should be constructed according to the deformed geometry. It is necessary to consider the continuity of materials, and the mesh coordinates of the deformed field must be transferred to the new field. The FEM modeling and simulation results with respect to lap resistance welding are introduced in this section, based on results reported by Bao (2000).
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4.5.1
Microjoining and nanojoining
Modeling
Figure 4.19 illustrates a cross-sectional model of lap resistance welding. The electrode is assumed to be a rigid body with resistance Re negligibly small and thermal conductivity much higher than the specimen. The sample is a stainless steel (SUS 304) plate 5 × 5 mm2 with thickness δ = 0.5 mm. The overlap ratio is unity (Lv = δ), where Lv is the initial contact width and δ is the thickness of the plate specimens. The bonding force is set as F = 800 N and the voltage between electrodes is fixed as U = 0.7 V. The electric current is direct current (d.c.). For Fig. 4.19(b), the two dimensional finite element model (four nodal points in each element) was adopted. Mesh division is finer at bonding area HD, as seen in Fig. 4.19(b). No slip was assumed at the bonding interface, but free slide was assumed to occur at the interface between the electrode and the specimen because no bonding occurred between them. Thermal conductance at the electrode interface and contact electrical conductance were given by taking into account the contact ratio due to surface roughness, therefore the contact electrical conductivity and thermal conductivity depend on bonding pressure. Temperature dependences with respect to the electric conductivity and thermal conductivity of the bulk materials were taken into consideration. Plastic deformation was assumed to Force Electric current Electrode (rigid body) A δ
Specimen B
Lv
Rigid surface G
D
H
C Bonding interface
Specimen E
F
Resistant Re = 0
(a)
A
G
B
H D
F
C
E (b)
4.19 Schematic illustration of lap resistance welding (a) the cross section and (b) mesh pattern of finite elements.
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occur along the stress-log strain curve, and seven stress strain curves were given in the range of 294~1253 K. The yield stress at temperatures with no stress strain curve was interpolated at the same strain level, and the yield stress at T > 1253 K was the same as that at T = 1253 K. The time increment ∆t was fixed at 0.01 ms so that the temperature difference and the electrode displacement changed by 1~2% between two steps t = t and t = t +∆t.
4.4.2
Numerical simulation
In general, electrical conductivity decreases with temperature, but contact conductivity increases with temperature because the contact ratio increases with temperature. At the beginning (0~1 ms) of the bonding process, contact resistance is very high. After that, contact resistance rapidly decreases with increasing contact ratio as temperature increases. In the second stage (1~5 ms), bulk resistance increases, the specimens are heated up and plastic deformation begins. In the third stage (5~7.8 ms), the reduction (collapse) starts from about 20% to about 100%, as seen in Fig. 4.20. The increase in contact width and plate deformation (edge collapse) decreases bulk resistance, because bulk resistance is determined by temperature and geometrical size. Figure 4.20 shows the numerical simulation results with respect to the collapse and temperature distribution. The collapse process finishes at t = 7.5~7.8 ms under the bonding conditions (U = 0.7 V, F = 800 N), although Fig. 4.20(b) shows the deformed pattern and temperature distribution at t = 7.5 ms. The maximum temperature is about (0.73~0.83)Tm at t = 7.8 ms, where Tm is the melting temperature (1723 K) of the specimen (SUS304). It follows that the collapse process is produced under solid state without nuggets. The top and bottom edges deform heavily to touch the electrode, but they are cooled by the electrode and do not melt. The collapse process finishes at 7.8 ms, but the temperature of the bonded central area continues to increase even at t = 7.8 ms. So, if lap welding is conducted for t > 7.8 ms, keeping U = 0.7 V, the nugget will be produced at the central area. The melting process is not taken into account in the lap resistance welding model. The melting process in the bonded area is required to predict expulsion. In the model, surface potentials of U = 0.7 V and 0 V are given at the upper and lower electrodes, respectively. In reality, the surface potential will be changed, the electrode will not always be flat and will exhibit elastic deformation. It is, therefore, very difficult to model the lap resistance welding.
4.6
Concluding remarks
In modeling the joining and welding process, simplifications and approximations are necessary and the models only hold true under certain assumptions and prescriptions. The models of solid state microjoining
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0.5 mm (a) 293 K 415 K 538 K 660 K 783 K 905 K 1028 K 1150 K 1273 K 1443 K
(b)
4.20 Numerical simulation results of lap resistance welding (a) t = 5.3 ms and Ds = 30.36%, where Ds is the compression ratio and is obtained by Ds = ds /δ, where ds is the electrode displacement and δ is the plate thickness, and (b) t = 7.51 ms and Ds = 95.50%.
introduced in this chapter require a lot of assumptions and are two dimensional, but they make it possible to understand interfacial deformation and the effect of several factors, including size ratio, mechanical properties, temperature, pressure, and vibration. The contact mechanism can be discussed quantitatively, although modified and improved treatments may be needed to further advance understanding. However, as the model becomes more complex in order to consider various additional factors, it often becomes harder to understand the main process mechanism. Solid state microjoining is naturally complex, and complete understanding requires complex models. But simple models like those introduced in this chapter help us to understand the predominant mechanisms.
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References and further reading
Bao Y. (2000), ‘Fundamental study on resistance projection welding control system by using numerical simulation,’ Doctor thesis of graduate school. Osaka University, pp. 12–38. Bao Y. et al. (2000), ‘Numerical simulation method on resistance projection welding control process,’ ISIJ Inter. Vol. 40, pp. S1–S4. Bay N. (1981), ‘Cold pressure welding – theoretical model for the bond strength,’ Proc. of the Joining of Metal: Practice and Performance. Vol. 2, Coventry, UK, 10–12 April, pp. 47–62. Harman G. G. and Albers J. (1977), ‘The Ultrasonic welding mechanism as applied to alminum- and gold-wire bonding in microelectronics,’ IEEE Trans., PHP-13, No. 4, pp. 406–412. Ishizaka A., Iwata S. and Tamamoto H. (1977), ‘Formation of clean Al surface by interface deformation in Au–Al thermocompression bonding,’ J. Jap. Met. Soc., Vol. 41, No. 11, pp. 1154–1160. Liu Y., Irving S. and Luk T. (2004) ‘Thermosonic wire bonding process simulation and bond pad over active stress analysis,’ The 54 Electronic Component and Technology Conference(ECTC), IEEE, June 1-4, Las Vegas, Nevada, USA, pp. 383–391. Lum I., Jung J. P. and Zhou Y. (2005) ‘Bonding mechanism in ultrasonic gold ball bonds on copper substrate,’ Metall. Mater. Trans. Vol. 36A, pp. 1279–1286. Mayer M., Paul O., Bolliger D. and Baltes H. (2000) ‘Integrated temperature microsensors for characterization and optimization of thermosonic ball bonding process,’ IEEE Trans. Compo. Pack. Technol. Vol. 23, No. 2, pp. 393–398. Onuki J. and Koizumi M. (1996), ‘Effect of Si precimitation on damage in ultrasonic bonding of thick Al wire,’ Materials Trans. JIM, Vol. 37, pp. 1319–1323. Saeki H., Nishitake H., Uemura T. and Yotsumoto T. (1991), ‘Deformation analysis of Au wire bonding,’ Preprint of Autumn Meeting of Plastisity and Material Processing Soc., 25–27 Sept., Sapporo, Japan, pp. 687–690. Schwizer J., Mayer M., Bolliger D., Paul O. and Baltes H. (1999), ‘Thermosonic ball bonding: Friction model based on integrated microsensor measurements,’ Proc. of Inter. Electron. Manufact. Technol. (IEMT) Symposium, 18–19 Oct., Austin, TX, pp. 108–114. Takahashi Y., Koguchi T. and Nishiguchi K. (1993), ‘Modeling of viscoplastic adhering process by a finite element technique,’ Trans. of the ASME, J. of Eng. Materials and Technology, Vol. 115, (1993) pp. 151–154. Takahashi Y., Shibamoto S., and Inoue K. (1996), ‘Numerical analysis of the interfacial contact process in wire thermocompression bonding,’ IEEE Trans. on Components Packaging and Manufacturing Technology, Part A, Vol. 19, No. 2, pp. 213–223. Takahashi Y. (1999), ‘Numerical analysis of microjoining process in electronic packaging,’ Proceedings of the Materials Solutions Conferences’99 on Joining of Advanced and Specialty Materials, 1–4 Nov. Cincinnati, Ohio, The Materials Information Sco., pp. 40–49. Takahashi Y., Inoue M. and Inoue K. (1999a), ‘Numerical analysis of fine lead bonding –effect of pad thickness on interfacial deformation,’ IEEE Trans. on Components and Packaging Technol. Vol. 22, No. 2, pp. 291–298. Takahashi Y., Inoue M., and Inoue K. (1999b), ‘Numerical analysis of fine lead bonding –effect of mechanical properties on interfacial deformation,’ IEEE Trans. on Components and Packaging Technol., Vol. 22, No. 4, pp. 558–566.
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Takahashi Y. and Inoue M. (2002), ‘Numerical study of wire bonding – analysis of interfacial deformation between wire and pad,’ ASME J. of Electronic Packaging,’ Vol. 124, pp. 27–36. Winchell IIV. H. and Berg H. M. (1978), ‘Enhancing ultrasonic bond development,’ IEEE Trans., on Components Hybrids and Manufacturing Technol., Vol. CHMT-1, No. 3, pp. 211–219. Zhang W. and Bay N. (1997), ‘Cold welding – theoretical modeling of weld formation,’ Welding J., Vol. 76, No. 10, pp. 417s–420s.
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5 Modeling of fusion microwelding B H C H A N G, Tsinghua University, P. R. China
5.1
Introduction
Fusion welding refers to those processes in which parts are united by localized melting under certain types of heat when pressure is not applied. For large scale structures (thickness of sheet or diameter of tubes greater than 0.5 mm), filler materials are sometimes employed during fusion welding, while for small scale structures (thickness of sheet or diameter of tubes less than 0.5 mm), only fusion welding processes that do not employ filler materials are used. The fusion welding processes used for small scale structures, also named fusion microwelding processes, include various arc microwelding (e.g., gas tungsten arc microwelding and plasma arc microwelding), laser microwelding and electron beam microwelding [1]. Resistance microwelding is another type of microjoining process, in which a weld is formed between two workpieces through the localized melting and coalescence of a small volume of the material(s) due to resistance heating caused by the passage of an electric current [2, 3]. Pressure is usually exerted on workpieces by electrodes in resistance microwelding. As a typical resistance microwelding process, resistance spot microwelding, sometimes called small scale resistance spot welding (SSRSW), has been increasingly employed to join thin sheet metals with a thickness less than 0.2–0.5 mm. This kind of application of resistance spot welding has many differences compared with conventional resistance spot welding (RSW), also referred to as ‘large-scale’ resistance spot welding (LSRSW), and is mainly used for sheets larger than 0.6–0.8 mm in thickness [3]. Both fusion and resistance microwelding have been increasingly applied in the fabrication of electronic components and medical devices. Nevertheless, due to the relatively small amount of energy delivered to very small components during the microwelding processes, the microwelds behave very differently from the conventional large scale welds, and the conclusions drawn in conventional welding are not applicable for microwelding. It has proven to be an unacceptable practice to ‘scale down’ the welding procedures suggested 121 WPNL2204
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for large scale workpieces in order to select parameters for microwelding processes [4–7]. Therefore, to obtain good weld quality, the welding procedure development process for microwelding can be both time consuming and costly. As we know, the thermal process is the center for both fusion and resistance microwelding processes, because it is related to all of the physical and chemical reactions during welding. Actually the thermal process determines the metallurgical reactions, melting, solidification, and other phase transformations during welding, and the geometry, composition, microstructure and properties of the resulting welds. Therefore, accurate measurement or calculation of the thermal process is actually a prerequisite to control the welding process and final quality [8]. However, the limited size of the welding zone, insufficient time for measurement, high heating and cooling rate, in addition to the complexity from coupling effects make it very difficult to measure the thermal cycle during the welding process through even the most ambitious experiments alone. On the other hand, numerical modeling of the welding process is not influenced so much by the restrictions mentioned above, and is a powerful tool for better understanding of the thermal and other physical processes associated with welding. That is why numerical methods have been increasingly employed in recent decades to model the microwelding processes for analyzing the thermal process, fluid flow, stress and distortion, etc. In these modeling efforts, more and more details of welding have been taken into account in the computations [9–13]. Several numerical methods such as finite element method, finite difference method, computational fluid dynamics method, finite volume method, boundary element method, etc., have so far been used to model fusion microwelding processes. Nowadays, the finite element method (especially using commercial finite element analysis packages such as ANSYS and ABAQUS) and the computational fluid dynamics (CFD) method (using commercial CFD packages such as FLUENT and PHOENIX) are more commonly used. In this chapter, the features of thermal process in fusion microwelding will be analyzed first; numerical modeling on the conduction type and the convection type thermal transfer in fusion microwelding will then be introduced respectively. The modeling of resistance microwelding will then be discussed. Finally, a summary and future trends in modeling of fusion and resistance microwelding will be presented.
5.2
Features of thermal process in fusion microwelding
As far as the fusion microwelding is concerned, to enable reliable joining, the heat source must have a high enough energy density. Different types of
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welding heat sources, for example arc, plasma arc, laser beam and electronic beam, have been used in fusion microwelding, in which the specimen is heated to melted state and then solidifies to form a weld. The thermal process of fusion microwelding has the following features: 1. Localized melting region Because the heat source acts on a local region of the specimen, large thermal gradients and uneven temperature fields are formed, which in turn lead to heterogeneity in microstructure and properties, incompatible thermal strains, residual stress and welding distortion. 2. Transient thermal history For fusion microwelding with a moving heat source, the thermal history of a point can be depicted with a thermal cycle curve. When the heat source approaches a specific point, the temperature of that point will rise rapidly; when the heat source is removed, the temperature decreases due to heat conduction. It is clear that the thermal process of heated point is transient, and the metallurgical reactions and phase transformations take place under non-equilibrium conditions. 3. Hybrid heat transfer The melted metal in a weld pool is not stationary; instead, it flows at a high speed. This means in the weld pool both conduction and convection types of heat transfer are present, while outside of the weld pool, conduction is the main mechanism. Also, the surfaces of a workpiece exchange heat with the surrounding environment by convection and radiation. Therefore, all three heat transfer mechanisms are involved in the fusion microwelding process. The aforementioned features make the thermal process in fusion microwelding very complex and bring many difficulties to numerical modeling. For instance, the modeling of such processes requires the use of fine grids and very small time steps, which make the investigation computationally intensive. Due to the rapid advances of computational hardware and software in recent years, numerical modeling of the thermal process in fusion microwelding is now more practical. So far, much work has been done is this area which can be divided into two categories. In the first category, the weld pool is considered as a thermal solid, in which only conduction type thermal transfer takes place. This type of analysis is used when the shape of weld pool is not so important and the main attention is on stress and distortion of the joint. In the second category, the weld pool is considered to be a thermal fluid, in which both conduction and convection types of thermal transfer take place. This type of analysis is used when the weld pool shape and dimensions, weld formation and defects like cracks and porosity in the weld are more important. In Sections 5.3 and 5.4, analyses on conductive and convective heat transfer will be discussed respectively.
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5.3
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Modeling of conductive heat transfer
The temperature distribution during fusion welding changes dramatically with geometrical position and time, and can be expressed as: T = f (x, y, z, t)
(1)
where T is the temperature, x, y and z are coordinates, and t is welding time. In addition, the thermo-physical properties of the base metal are strongly dependant on temperature, and the latent heat is absorbed or released during melting/solidifying and other phase transformation processes. Therefore, the thermal process in fusion welding can be treated as a typical nonlinear transient heat transfer problem.
5.3.1
Governing equations
Based on the basic principles of thermal transfer and conservation of energy, the three dimensional heat conduction equation governing temperature distribution within specimens with heat generation may be expressed as [14]:
ρC p ∂T = ∂T λ ∂T + ∂T λ ∂T + ∂T λ ∂T + Q ( x , y, z , t ) ∂t ∂x ∂x ∂y ∂y ∂z ∂z (2) where Q is the heat generation rate, λ is the thermal conductivity, Cp is the specific heat, and ρ is the density. Three types of boundary conditions are applied in thermal conduction analysis, including: 1. First type boundary condition, in which the temperature at a boundary is specified, i.e. Ts = Ts(x, y, z, t)
(3)
2. Second type boundary condition, in which the distribution of heat flux at a boundary is specified, i.e. λ ∂T = q s ( x , y, z , t ) ∂n
(4)
3. Third type boundary condition, where heat transfer between workpiece and the environment is known, i.e.
λ ∂T = α ( Ta – Ts ) ∂n
(5)
In the above equations, n is the normal outward direction of the boundary surface; qs is the input heat flow from outside on the unit area; α is the
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coefficient of heat transfer at the surface; and Ta is the ambient temperature. In fusing microwelding modeling, the second and the third boundary conditions are usually applied.
5.3.2
Solution of governing equations
The key to solving the conductive heat transfer problem is to find solutions to Eqn. 5.2 to Eqn. 5.5. Analytical solutions can be obtained under very strict assumptions and for some very simplified cases. Please refer to the works from Rosenthal [15] and Rykalin [16] for more information about analytical solutions on welding heat process problems. Assumptions made in analytical solutions result in an error in calculated temperatures especially for the part near to the heat source. In fact, the region around the heat source is most important because the temperatures in this region determine the shape and size of weld pool, and the metallurgical reactions taking place within the weld pool. In contrast, the numerical analysis methods have much higher precision in predicting temperatures in the region around the heat source. As far as numerical methods are concerned, the finite difference method (FDM) and finite element method (FEM) have both been used to solve heat conduction equations. The FEM has been used more and more widely in recent years because it can flexibly fit to the various boundary conditions of complex welding phenomena and can unify the analysis of temperature and the analysis of thermal stress. In many instances, the purpose of computing the temperature is to analyze thermal stress, residual stress, and distortion of welding. As an example, the following will give the solving procedure when using FEM. When solving the nonlinear heat conduction differential equation of fusion welding with FEM, the equation is firstly discretised in the spatial region. Note the shape function as [N], and the node temperature of an element as {T}e, then the temperature in the element can be expressed as [14]: T = [N]{T}e
(6)
By the weighed residual method, the following equation can be obtained:
[ K ]{T } + [ C ] ∂ {T } = {P} ∂t
(7)
in which, the coefficient matrix [K] is heat conduction matrix, {T} is temperature vector, [C] is heat capacity matrix and {P} is heat flow vector. These matrix and vectors are calculated by the following expressions: [K] = ∑ ([K1]e + [K2]e) [C] = ∑ [C]e {P} = ∑ ({P1}e + {P2}e + {P3}e)
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∂ [ N ] T ∂[ N ] ∂[ N ] T ∂[ N ] ∂[ N ] T ∂[ N ] ∂x λ ∂x + ∂y λ ∂y + ∂z λ ∂z dV
[ K1 ] e =
∫
[ K 2 ]e =
∫
[ N ] T α [ N ] dS
∫ [P ] = ∫ [P ] = ∫ [P ] = ∫
[ N ] T ρc [ N ] dV
∆V
∆S
[ C] e =
∆V
e
1
[ N ] T Q dV
∆V
2
3
e
[ N ] T qdS
∆S
e
∆S
[ N ] T α Ta dS
Note that Eqn. 5.7 is nonlinear because the parameters in it (e.g. λ, c, ρ, α) are all temperature dependant. After being discretised in the spatial region, the heat conduction differential equation is then discretised in the temporal region. Within each time step from t to t + ∆t a differential scheme is established for time t + θ∆t, where θ is the weighting coefficient, with a value of 0 ≤ θ ≤ 1. Expressing {T} by the Taylor series expansion equation: {T}(t+θ∆t) = θ{T}(t+∆t) + (1 – θ){T}(t) + o(∆t2)
(8)
∂ {T }( t + θ∆t ) = 1 ({T }( t + ∆t ) – {T }( t ) ) + o ( ∆t 2 ) ∆t ∂t
(9)
where o(∆t2) is a minor term and can be omitted. By substituting Eqn. 5.8 and Eqn. 5.9 into Eqn. 5.7, and expressing {P} in the same way as {T}, the nonlinear differential equation is then transformed into a nonlinear algebra equation as follows:
( ∆1t [C ] + θ[ K ]) {T} θ
θ
( t + ∆t )
(
)
= 1 [ C θ ] – (1 – θ ) [ K θ ] {T }( t ) ∆t
+ θP(t+∆t) + (1 – θ)P(t) θ
(10)
θ
The superscript θ indicates that the matrices [C ] and [K ] are calculated with the temperature T(t+θ∆t) at time t + θ∆t. Equation 5.10 can be abbreviated as [H]{T} = {F}
(11)
By solving the above equations, the temperature field in the specimen can then be obtained.
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Definition of the heat source model
A critical problem in numerical modeling of the heat process in fusion microwelding is the accurate description of the heat source model. Because of the distinct characteristics of both heat source and the interactions between heat source and specimen, different models should be developed for different welding processes. For micro arc welding processes such as micro gas tungsten arc welding and micro plasma arc welding, a two-dimensional Gaussian distribution of heat flux is often employed, which can be expressed as [17]: 2 q ( r ) = q m exp –3 r 2 r
(12)
where qm is the peak heat flux at the center, r is the efficient distribution radius of heat flux, and r is the distance to the center of heat source. Heat source models for laser and electronic beam welding are much more complex, because of their different action mechanisms compared to arc welding. There are typically two types of welding, one is conduction type welding when power density is below some critical value, and the other is keyhole type welding when the power density is greater than the critical value. For conduction type welding, the same model as Eqn. 5.12 can be used. However, for keyhole type welding, a three-dimensional volume type heat source model is necessary to accurately predict the boundary of the weld pool. So far, several volume heat source models have been developed, such as column, cone, semi-sphere, and other rotary type heat sources [18]. In some cases, the planer and volume heat sources are combined together to obtain better temperature results in deep penetration welding. Chang and Na [19] studied the effects of various heat source models on prediction accuracy of laser weld shape in the microjoining of a small structure.
5.3.4
Example
Brockmann et al. [20] numerically studied the temperature field during laser microwelding of thin metal foils, as shown in Fig. 5.1. A thin metal sheet of thickness h, moves with a velocity u is heated by a laser beam of power W from the upper surface. Sheet metal will melt when the beam power is greater than critical value, Wcr. For such laser microwelding, the size of the melted zone is small (100–300 µm). This means that the metal is kept in a horizontal layer by surface tension and curving of the weld pool surface is negligible. The contour K is the boundary of the melted region. The mathematical model that describes the heat transfer in above problem is developed and solved by the explicit finite difference numerical model. Both 2-D (when the thickness of foils is less than radius of laser beam) and 3-D problems are considered.
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Gas (Tg)
Radiation losses Convection losses
x
y
z
u
K
5.1 Principal scheme of heating and melting of the moving sheet by laser. [20]
A Gaussian distribution function for laser power density is assumed. The melting, evaporation, and solidification of the metal are all taken into account. Convection and radiation cooling, and the temperature dependencies of materials properties are also considered. The heat conductivity of melted metal is increased several times higher than the real value and this has shown to be a simple way to take into account the fluid flow (Marangoni effect) in the weld pool. Figure 5.2 demonstrates the differences in predicted temperature fields in a stainless steel sheet when melting is considered or not considered. Obviously, when melting is taken into account, the temperature is lower due to the absorption of latent heat during the melting process. Figure 5.3 shows the calculated temperature field of the underside of the thin sheet. The gray color region has a temperature above the melting temperature of base metal. In this case, the melting region is approximately 250 µm in X direction and 200 µm in Y direction. Therefore, a fully penetrated welding seam can be formed. Figure 5.4 shows the comparisons between the cross section of the resulting weld seam with the calculated temperature field. It can be seen that the calculation agrees pretty much with the real weld. Both the shape and dimensions of the weld are well related to that of the predicted melted zone.
5.4
Modeling of convective heat transfer
As indicated above, both conductive and convective heat transfer mechanisms exist in the weld pool during fusion microwelding. The relative importance
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250
10
500 750
5
Y
129
1750
0
1500
1000 1250
–5 –10 –15 –10
0
10
20
30
40
X (a) 15
250
10
500
5
750 1000
Y
1750 0
1500 1250
–5 –10 –15 –10
0
10
20
30
40
X (b)
5.2 Isotherms for (a) with melting and various properties in liquid and solid state and (b) without melting and constant properties. [20]
of convection and conduction in the overall transport of heat can be evaluated from the value of the Peclet number, Pe, which is defined by:
Pe =
uρCL k
(13)
where u is average velocity of fluid flow, L is the characteristic length (generally taken as the pool radius at the top surface of weld pool), ρ is density, C is specific heat and k is thermal conductivity. When Pe is less than 1, the heat transfer within the weld pool occurs primarily by conduction. When Pe is much higher than 1, the primary mechanism of heat transfer is convection. Obviously, the influence of convective heat transfer should be taken into
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1500
1000
Temperature, °C
2000
500
0 750 500
–250
Y
di
re
ct
0 io
n,
0 µm 250 –250
X
250 µm n, tio c e dir
5.3 Calculated temperature field of the bottom side of a laser heated sheet of stainless steel. The melting region is in gray. [20]
account in numerical modeling to obtain a more accurate thermal history and to study weld microstructure, grain growth, phase transformation kinetics, hot cracking, and porosity, etc., with greater accuracy. Oreper et al. [21, 22] carried out the earliest study in this area, in which the evolution of temperature and velocity fields, solidification rate, and thermal gradient were computed using a two-dimensional model, and the role of conduction and convection in the weld pool during GTAW welding were compared. Zacharia et al. [23, 24] investigated the effect of the surface active element (sulfur) on the weld pool geometry by computational and experimental methods. Betram [25] studied the role of heat and liquid flows on the final solidification microstructure considering the mushy zone behavior. Lei and Shi et al. [26, 27] investigated the heat transfer and fluid flow in the weld pool during GTAW and laser spot welding. Debroy et al. [28, 29] presented a numerical model to simulate heat transfer and fluid flow during steady and transient fusion welding. This model has been used to calculate weld pool geometry, temperature and velocity fields for various materials under moving and stationary heat sources and for laser beam as well as arc welding. In summary, numerical models considering convective heat transfer have been successful in revealing special features in transient fusion welding processes [30]. This section will discuss the general principle and methods in solving convective heat transfer problems in fusing microwelding.
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bw
bw = 150 µm tw = 100 µm
50
–250
–200
–150
–100
–50 0 50 Y direction, µm
100
150
200
0 250
Z direction, µm
100
5.4 Comparison between calculated temperature field and cross section of the corresponding micro-laser welding seam. [20]
5.4.1
Driving forces of fluid flow in the weld pool
When modeling the fluid flow in the weld pool during fusion microwelding, the following driving forces are generally incorporated in the mathematical model: 1. The surface tension gradient driven force, also called the Marangoni stress, is one of the major factors that influences the flow pattern and temperature of liquid metal in weld pool; surface tension directly influences the fluid flow and temperature in a weld pool. The Marangoni stress is determined by the temperature distribution on the surface of a weld pool. For a pure metal, the surface tension decreases with the increase of temperature, i.e., ∂σ < 0 . When the surface active elements like S, O, ∂T Se, and Te are present, the surface tension not only depends on temperature but also on the concentration of these elements. Under certain conditions, surface tension will increase with temperature, i.e., ∂σ < 0 . Fluid flow ∂T directions will be opposite for different signs of temperature coefficient
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2.
3.
4.
5.
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of surface tension, and will therefore lead to different shapes of weld pool. When heat energy is provided by a welding arc, the electric current of the arc will disperse after entering the weld pool. The electromagnetic force arises from the interaction between electrical current and the magnetic field generated by the current itself. The electromagnetic force will push the liquid metal in the center of weld pool flow downward, and then back to the top surface of the weld pool along the fusion line. On the top surface of the weld pool, the flow is from outer edge toward the center in radial direction. The buoyancy force results from the thermal gradient and concentration gradient, which make the density different in the weld pool. The part with higher temperature has lower density, and will flow upward to the surface of weld pool. If the density of solute is greater than the solvent, the part with higher concentration will flow to the bottom of the weld pool. Generally, the buoyancy is a main reason for natural convection; but it is much less significant than electromagnetic force and Marangoni stress. Therefore, it is often precluded in modeling. Different from macro-fusion welding, the droplet transfer is generally not present in fusion microwelding processes. Therefore, it is not necessary to simulate the impinging of droplets on the weld pool and the heat and mass transfer associated with droplet transfer. Pleaser refer to [8] for details on modeling of fusion welding process considering metal transfer. Metal evaporation on the weld pool surface not only influences the temperature of the surface, but also results in a recoiling force that will depress the weld pool surface. Because of the distinct characteristics of heat sources, e.g. laser beam and tungsten arc, metal evaporation has different influences on surface geometry and the temperature distribution of the weld pool.
5.4.2
Governing equations
The mathematical model and corresponding governing equations can be different for various fusion microwelding processes. Even for the same welding process, the model and governing equations can be different when considering different aspects of the problem. Taking the heat and fluid flow in gas tungsten arc welding as an example; to establish the mathematical model, an incompressible, laminar and Newtonian liquid flow is generally assumed in the weld pool. The thermal physical parameters are independent of temperature except for the surface tension, specific heat and thermal conductivity. The driving forces of liquid metal in the weld pool include natural convection, electromagnetic forces, surface
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tension, and buoyancy. Boussinesq assumption is valid, i.e. the density is considered as constant except the density in the buoyancy term. When the heat source (arc) with heat flux of q(r) moves along ξ axis at a constant speed of u0, in a fixed three-dimensional Cartesian coordinate system (ξ, y, z), based on the relationship between the moving coordinates (x, y, z) and the stationary coordinates, x = ξ – u0t, the stationary coordinates can be transformed into the moving coordinates in which the center of the heat source is located at the origin, and x is the distance to the heat source center in the welding direction. The governing equations for this problem are as follows: 1. The equation for mass conservation: ∂ρ ∂ρu ∂ρv ∂ρw + + + =0 ∂t ∂x ∂y ∂z
(14)
2. The equations for the conservation of momentum are: •
in x direction: ∂ ( ρu ) ∂ ( ρU 0 u ) ∂ ( ρuu ) ∂ ( ρvu ) ∂ ( ρwu ) ∂t – + ∂x + ∂y + ∂z ∂x
= – •
∂p + ∂ µ ∂u + ∂ µ ∂u + ∂ µ ∂u + S x ∂x ∂x ∂x ∂y ∂y ∂z ∂z
(15)
in y direction: ∂ ( ρv ) ∂ ( ρU 0 v ) ∂ ( ρuv ) ∂ ( ρvv ) ∂ ( ρwv ) ∂t – + ∂x + ∂y + ∂z ∂x
= – •
∂p + ∂ µ ∂v + ∂ µ ∂v + ∂ µ ∂v + S y ∂y ∂x ∂x ∂y ∂y ∂z ∂z
(16)
in z direction: ∂ ( ρw ) ∂ ( ρU 0 w ) ∂ ( ρuw ) ∂ ( ρvw ) ∂ ( ρww ) + + ∂t – + ∂x ∂y ∂z ∂x
= –
∂p + ∂ µ ∂w + ∂ µ ∂w + ∂ µ ∂w + S z ∂z ∂x ∂x ∂y ∂y ∂z ∂z
(17)
3. The equation for the conservation of energy is
∂ ( ρh ) ∂ ( ρU 0 h ) ∂ ( ρuh ) ∂ ( ρvh ) ∂ ( ρwh ) ∂t – + ∂x + ∂y + ∂z ∂x = ∂ λ ∂T + ∂ λ ∂T + ∂ λ ∂T + Sh ∂x ∂x ∂y ∂y ∂z ∂z
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in which, ρ is the density, h is the enthalpy, λ is the thermal conductivity, T is the temperature, t is the time, µ is the viscosity, p is the pressure, and the u, v w are velocity components in the x, y and z directions respectively. When the enthalpy–porosity technique is used to handle the phase transformation problem, the source term in the momentum equations has the following form: µρ Sx = – u + ( J × B) x K
(19)
µρ S y = – v + ( J × B ) y + ρgβ ( T – Tm ) K
(20)
µρ Sz = – w + ( J × B ) z K
(21)
In Eqns 5.19–5.21, the first term on the right side is the model describing fluid flow in the mushy zone (solid–liquid two-phase zone), which is called Darcy resistance. K is the permeability, which is a measure of the ease with which fluids pass through the porous mushy zone, and is related to the morphology of dendrites. Assuming K is isotropic and only dependant on the liquid mass fraction in the mushy zone, it obeys the Carman–Kozeny equation: fl3 K = K0 2 (1 – f l )
(22)
Where K0 is a constant related to the dendrite dimension and fl is the liquid mass fraction. The temperature dependency of fl can be dealt with linearly as follows: T fl = T l
0 – Ts – Ts 1
T < Ts Ts ≤ T ≤ Tl
(23)
T > Tl
The energy equation is given in the form of enthalpy that can be calculated by the following equation: h = (1 – fl)csT + flclT + ∆H
(24)
in which, cs and cl are specific heats for solid and liquid in mushy zone respectively, and ∆H is the latent heat fraction in enthalpy. The relationship between ∆H and temperature can be expressed as: L ∆H = f l L 0
T > Tl Ts ≤ T ≤ Tl T < Ts
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in which L is the latent heat of fusion of the base metal, and Ts and Tl are solidus and liquidus temperature, respectively. When cs = cl = c, the equation for the conservation of energy can take the temperature as variable and can be written as:
∂ ( ρcT ) ∂ ( ρU 0 cT ) ∂ ( ρucT ) ∂ ( ρvcT ) ∂ ( ρwcT ) + + ∂t – + ∂x ∂y ∂z ∂x = ∂ λ ∂T + ∂ λ ∂T + ∂ λ ∂T + Sh ∂x ∂x ∂y ∂y ∂z ∂z
(26)
And the source term can be expressed as: Sh = –
∂ ( ∆H ) ∂ ( ρu∆H ) ∂ ( ρv∆H ) ∂ ( ρw∆H ) – – – ∂t ∂x ∂y ∂z
(27)
In Eqns 5.19–5.21, there is a term arising from the electromagnetic force. Assuming that the electric field is in a quasi-steady state and the electrical conductivity is constant, the electromagnetic force can be calculated by solving the following Maxwell equations:
∇⋅J=0 J = – σ ∇ϕ 2 ∇ ϕ=0 ∇ × B = µ 0 J
(28)
in which, J is the current density, σ is the electrical conductivity, ϕ is the electrical potential, B is the magnetic flux density, and µ0 is the magnetic permeability. The electromagnetic forces are applied to the source term in the momentum equation by the three components in x, y and z directions, respectively.
5.4.3
Boundary conditions
Boundary conditions are much more complex when fluid flow in the weld pool is considered. Generally, the boundary conditions for the solution region of Eqns 5.14–5.18 include the following aspects: 1. 2. 3. 4. 5.
normal to the free surface boundary tangential to the free surface boundary top surface boundary other surfaces boundary symmetrical boundary.
The symmetrical boundary condition is easy to understand and realized in
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modeling and will not be addressed here; the other four types of boundary conditions are discussed below. Normal to the boundary free surface For the free surface of the weld pool, the following pressure conditions must be satisfied: p = pv + γκ
(29)
where p is the pressure at the free surface in the direction normal to the local free surface, pv is the vapor pressure or any other external pressure acting on the free surface (in GTAW, it can be the plasma arc pressure), γ is the surface tension, and κ is the free surface curvature. When the free surface is assumed to be flat, the second term in Eqn. 5.29 is equal to zero. Tangential to the free surface boundary With the assumption of flat surface of the weld pool, the temperature dependant Marangoni shear stress at the free surface is given by
– µ ∂u = σ s ∂T ; – µ ∂v = σ s ∂T ∂z ∂x ∂z ∂y
(30)
where µ is the viscosity, and σs is the thermal coefficient of surface tension. Top surface boundary The temperature boundary conditions at the top surface of the base metal are determined by both the heat input from heat sources (such as welding arc, laser beam and electronic beam) and the losses due to convection, radiation, and evaporation; the boundary condition can be described as
q ( r ) – q conv – q radi – q evap = – λ ∂T ∂z
(31)
The heat flux q(r) can have different forms when different heat sources are employed, as mentioned in the previous section. The heat losses due to convection, radiation, and evaporation can be written as qconv = α(T – Ta)
(32)
q radi = σε ( T –
(33)
4
Ta4 )
qevap = WHv
(34)
where α is thermal transfer coefficient, Ta is the ambient temperature, σ is the emissivity, ε is the Stefan-Boltzmann constant, Hv is the latent heat for liquid–vapor phase change, and W is the melt mass evaporation rate.
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The electrical boundary condition at the top surface of the base metal is similar to that of heat input, and the current density has Gaussian distribution 2 J ( r ) = 3 I2 exp – 3 r2 π rj rj
(35)
in which rj is the distribution radius of the current density. Other surfaces boundary For the part of the top surface where w = 0, side and bottom surfaces of the specimen where the velocities meet u = v = w = 0, the convective boundary is specified by
α ( T – Ta ) = – λ ∂T ∂n where n is a normal unit vector to the surface.
5.4.4
(36)
Solutions of governing equations
When solving the governing equations 5.14–5.18, different methods can be used such as SIMPLE, SIMPLEC and PISO [31]. The finite differential algorithm named SIMPLE (Semi-implicit method for pressure-linked equation) is very effective in solving coupling fields of velocity and pressure and will be described briefly below. Based on the assumption of static electrical current field, the current and fluid flow are independent of each other. Therefore, the electromagnetic force is calculated first and then used as body force in the momentum equation. Then, the governing equations of velocity and pressure fields are solved in an iterative way. The basic procedure of the SIMPLE algorithm is: 1. give/calculate the velocity, temperature, and pressure 2. calculate the coefficients of momentum equations and obtain a set of approximations on velocity; the pressure gradient term is ignored temporarily 3. substitute the approximated velocity into the pressure equation to calculate an approximated pressure field 4. substitute the calculated pressure results into the momentum equations and obtain a velocity field u*, v* and w* 5 Take u*, v* and w* into the pressure modification equation, obtain p*, and substitute p* to the velocity modification equation to modify the velocity results 6. back to step (2); iterate steps (2)–(6) until convergence is reached 7. increase the time by a time step 8. back to step (1), continue the process until the preset time is reached.
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5.4.5
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Examples
Heat transfer and fluid flow in three dimensional gas tungsten arc welding The heat transfer and fluid flow in gas tungsten arc welding (GTAW) with a moving heat source has been studied by Shi et al. [18], as depicted in Fig. 5.5. The coupled mass, momentum and energy equations are solved by the finite volume method (FVM) with the commercial computational fluid dynamics (CFD) package Phoenics. The base metal is mild steel with the thermo-physical parameters listed in Table 5.1. Two sets of welding parameters used in the modeling are listed in Table 5.2. Figure 5.6 shows the three-dimensional temperature distributions under two welding conditions. It can be found that the temperature at the surface
Welding torch
Welding arc
Symmetry plane
Weld
g Travelin
direction
Y X Base metal
Z Heat affected zone
5.5 Schematic of GTAW process with moving heat source. Table 5.1 Thermal physical property parameters used in computations Property/unit
Value
Density/(kg/m3) Solidus temperature/°C Liquid temperature/°C Viscosity/(kg/m·s) Thermal conductivity in solid state/(J/m·s·°C) Thermal conductivity in liquid state/(J/m·s·°C) Specific heat in solid state/(J/kg·°C) Specific heat in liquid state/(J/kg·°C) Latent heat of fusion/(J/kg) Thermal coefficient of surface tension/(N/m·°C) Thickness of plate/mm Width of calculation region/mm Length of plate/mm Number of grid
7200.0 1472.0 1512.0 0.006 25.1 83.6 702.0 806.74 267.5 × 103 –3.0 × 10–3 2.8 15.0 60 80 × 50 × 15
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Table 5.2 Welding parameters used in computations Parameters/unit
Procedure 1
Procedure 2
Welding current/A Welding voltage/V Traveling speed/(mm/s) Effective heating radius, ra /mm Efficiency of heat source, η
115 25.4 5.6 6.5 70.0%
79 22 5.6 5.5 80.5%
25 289 392 576 759 942 1126 1310 1493 1676 1860 2043 2227 2410 2594
(a) Procedure 1
25 184 342 500 659 817 926 1134 1293 1451 1609 1768 1926 2085 2043
(b) Procedure 2
5.6 Three dimensional temperature fields calculated with different welding procedures.
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Table 5.3 Calculated and experimentally measured width and depth of weld pool Welding parameters
Depth of weld pool/mm Measured/calculated
Width of weld pool/mm Measured/calculated
Procedure 1 Procedure 2
2.5/2.34 1.5/1.12
8.7/7.1 6.5/5.0
of weld pool increases with the welding power input, and the depth and width of weld pool increase correspondingly. The maximum temperature at the weld pool surface is below the vaporization point. The calculated depth and width of weld are compared to the experimental results to validate the numerical model as listed in Table 5.3. It can be found that the calculated depth and width of the weld pool is a little smaller than the experimental results. The differences have been attributed to the errors in estimating the arc efficiency η and effective heating radius ra. Nowadays, these two parameters are usually determined by experiments or experiences in numerical calculations. Figure 5.7 shows the velocity fields under two computational conditions. Obviously, both the melting region and the velocity are increased for increased power input. Heat transfer and fluid flow in laser microwelding He and Debroy et al. [32] developed a transient numerical model to study heat transfer and fluid flow during laser spot microwelding of a stainless steel. Surface tension and buoyancy forces were considered for the calculation of transient convective heat transfer in weld pool. Very fine grids and small time steps were used to achieve accuracy in the calculations. Table 5.4 presents the data used for these calculations. The numerically predicted weld pool cross sections are compared with the corresponding experimental values, as shown in Fig. 5.8. Agreement is required between the calculated and the experimental weld pool geometry and dimensions, and it can be seen that with an increase in the beam diameter, the weld pool becomes wider and shallower. Figure 5.9 shows the computed temperature and velocity fields as a function of time. At the beginning stage, the weld pool expands rapidly in size and the temperature and velocities increase with time. At the end of the pulse, the peak temperature drops and the weld pool shrinks rapidly. The maximum velocity in the weld pool is about 95 cm/s. A two-phase solid–liquid mushy zone exists within the thin region between the solidus and liquidus isotherms. The size of this zone is very small during heating, and the size of mushy zone increases significantly at the end of pulse. As indicated above in Eqn. 5.13, the relative importance of convective and conductive heat transfer in the weld pool can be evaluated from the
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: 1.6667E-82 m/s 3-D Melting of steel
(a) Procedure 1
: 1.6667E-82 m/s 3-D Melting of steel (b) Procedure 2
5.7 Velocity field of fluid flow in weld pool calculated with different welding procedures. Table 5.4 Thermal physical property parameters used in computations32 Property/unit
Value
Density of liquid metal/(g/cm3) Absorption coefficient Effective viscosity/(g/cm·s) Solidus temperature/K Liquidus temperature/K Enthalpy of solid at melting point/(cal/g) Thermal conductivity in solid state/(cal/g) Specific heat of solid/(cal/g·K) Specific heat in liquid state/(cal/g·K) Thermal conductivity of solid/(cal/cm·s·K) Effective thermal conductivity of liquid state/(cal/cm·s·K) Temperature coefficient of surface tension/(Nc/m·K) Coefficient of thermal expansion
7.2 0.27 0.006 1697 1727 286.6 300.0 0.17 0.20 0.046 0.5 –0.43e–5 1.96e–5
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0
0.5 1.5
1.5
1.4
1.4 1.3
16
1.2
K 97
1.3 500 mm/s
z, mm
142
1.2 1.1
1.1 –0.5
0 x, mm
0.5
(a) –0.5
0
0.5
1.5 1.4
16
1.3
97
1.4 K
1.3
1.2
z, mm
1.5
1.2 500 mm/s
1.1
1.1 –0.5
0 x, mm
0.5
(b)
5.8 Experimental and calculated weld pool cross sections for laser power of 1967W and pulse duration of 3 ms under (a) beam radius of 0.428 mm and (b) beam radius of 0.57 mm. [32]
value of the Peclet number, Pe. Figure 5.10 shows the change of maximum Peclet number with time in the weld pool. It can be seen that during most of the welding time, the Peclet number is greater than 1 and convection is the most important heat transport mechanism in the weld pool. This explains why the numerical modeling that takes fluid flow into account is more close to the real welding process.
5.5
Modeling of resistance microwelding
5.5.1
Principles and features
A typical set-up for SSRSW is shown in Fig. 5.11. During welding, an electrode force is applied to achieve good contact at the electrode/workpiece (E/W) and workpiece/workpiece (W/W) interfaces. Then, an electric current passes through the top and bottom electrodes and heats the workpieces by joule heating: Q = I2Rt
(37)
where Q is the heat generation, I the welding current, R the resistance, and t the duration of current (weld time). When the temperature at the W/W interfaces reaches the melting point of the material, a molten nugget begins
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0 x, mm
1.2 0.4
0.2
m
0 –0.4
0.2
1.3
0.4 m y,
m
m y,
0.2
1.3
1500 mm/s 0 –0.4
–0.2
0 –0.4
–0.2
1000 0 x, mm
0.2
1.5
0.4 0.2
mm
1500 mm/s
1.2 0.4
y,
m
m y,
0.2
1.3
1727 1697
1.2 0.4
1.4
z, mm
0.4
0 x, mm (b)
0.2
Mushy zone
1.5 1.4
1697 1000
(c)
1500 mm/s 0 –0.4
–0.2
1697
1.3
1000
1.2 0.4
0 x, mm
0.2
(d)
Modeling of fusion microwelding
(a)
Mushy zone
1.4
2200
z, mm
0.4
z, mm
1.4 1697 1000
1.5
3000
z, mm
1.5
2700
5.9 Computed temperature and velocity fields at different time: (a) t = 1 ms, (b) t = 4 ms, (c) t = 4.5 ms, (d) t = 5 ms. Laser power. [32]
143
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Maximum peclet number
5 4 3 2 1 0 0
1
2
3 Time, ms
4
5
6
5.10 The variation of maximum Peclet number with time. [32] Current
Force
Z Top electrode
Faying surface
Top workpiece
R Nugget
Bottom workpiece
Bottom electrode
5.11 Schematic set-up for SSRSW.
to form and grow. When the welding current is turned off, this molten nugget solidifies to form a spot weld that joins the workpieces together. The electrode force is maintained during the whole process to ensure continuity of the electric current, and is continually applied for a short period after the current is turned off. Unlike LSRSW, no water is used to cool the electrodes because of the limited electrode dimensions of SSRSW. From previous discussions, it is clear that the SSRSW is a complex process in which coupled interactions exist between electrical, thermal, mechanical, and even metallurgical phenomena. The SSRSW is a transient thermal process associated with both material nonlinearity from temperature dependent physical
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properties of workpieces and the geometrical nonlinearity from contacts at the E/W and the W/W interfaces. Due to the complexity of the SSRSW process mentioned above, some challenges exist in numerical modeling work, such as: what kind of analysis procedure should be used to model the coupled interactions between multiple disciplines; how to accurately model the contact resistances at interfaces; how to construct the corresponding finite element mesh and specify appropriately the boundary conditions according to the characteristics of resistance microwelding process, etc. The following will discuss the above aspects respectively.
5.5.2
Analysis procedure
A representative finite difference model is the two-dimensional model by Cho and Cho [9]. This finite difference model solves the coupled thermoelectric problem to simulate the nugget growth process. The advantages of finite difference models lie in their simplicity and moderate demand for computing power. However, finite difference models are, in general, not used to solve mechanical problems. Consequently, the interaction between the mechanical process and the thermal–electrical process in the RSW process cannot be properly accounted for with the finite difference methods. Nied [10] first used ANSYS to simulate the squeeze action and obtained the contact radii at E/W and W/W interfaces. It was found that the contact radius at the W/W interface is about 25% larger than the radius of the electrode tip surface. Assuming the contact areas remain constant, a coupled thermoelectric analysis is then conducted to simulate the nugget growth process. The same analysis procedure was adopted and expanded by Tsai et al. [11, 12]. The analysis procedure used by Nied and by Tsai et al. ignored the influence of mechanical processes on the conduction of welding current and the heat generation process. Syed and Sheppard [13] recognized that the contact areas during the RSW process keep changing due to the interaction between squeeze pressure and thermal expansion in the heated area. They proposed a fully coupled thermal– electrical–mechanical analysis procedure, which involves a significant amount of back and forth iterations for both the thermoelectric and thermomechanical analyses. Such a fully coupled analysis procedure is conceptually rigorous and is a significant step forward in terms of how RSW should be modeled. However, this procedure can be very demanding for computing power and may be cost prohibitive. It is also likely to encounter numerical convergence problems. Browne et al. [33] proposed a more robust procedure, which involves incrementally updating the contact information from the thermal–mechanical analysis using ANSYS to the thermal–electric analysis. They used ANSYS
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for their thermal–mechanical analysis and an in-house finite difference program to perform the coupled thermal–electrical analysis. Similar to Browne’s work, the incrementally coupled thermal–electrical– mechanical analysis procedure has been fully implemented in ANSYS by Chang et al. [34, 35] and Li et al. [36]. Figure 5.12 shows a flow chart of the analysis procedure in which the computations shuttle between the electrical– thermal analysis and the thermal–mechanical analysis by a small time step. Electrical–thermal analysis is used to calculate the joule heat and temperature development during the welding process; the temperature calculated is then imposed into the thermal–mechanical analysis to calculate the deformation, strain and stress. At the next time step, deformed configuration is updated in the electrical–thermal analysis to account for the variation of actual contact area. This process continues until the end of specified welding time. This incrementally coupled analysis procedure has also been implemented using the commercial FEA code ABAQUS [37, 38] in the UNIX operation system.
5.5.3
Modeling of contact resistance
Spot welding is achieved by the ohmic heating from the resistance to electric current flow. The resistance includes two bulk resistances of the electrodes, Apply force
Initial contact state
t = ti = dt t = t + dt
Electric-thermal analysis
Temperature
Elastic-plastic mechanical analysis
New contact state
End
5.12 Flow chart of the incrementally coupled algorithm.
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two bulk resistances of workpieces, two contact resistances at the E/W interfaces and one contact resistance at the W/W interface. Both the electrodes and the workpiece materials have well-defined values of electrical resistivity. However, the contact resistances at the interfaces are not well defined. This creates a critical problem in the modeling of RSW to quantify the contact resistance. Browne et al. [33] developed a special instrument to measure the dynamic resistance during the welding of aluminum alloys; they artificially adjusted the contact resistance values until a reasonable agreement between the predicted and measured dynamic resistance was reached. Others have used the measured static resistance data and extended them to elevated temperatures based on the dependence of materials hardness or yield strength on temperature [9– 12]. Apparently both approaches involved a great deal of trial-and-error and variability from material to material. Feng, et al. [37] proposed a phenomenological contact resistance model on the basis of an earlier model proposed by Greenwood [39]. Essentially, they established a relationship to correlate the contact resistance with contact pressure, interface temperature, and electrical resistivity of the materials in contact. Li et al. [36] first proposed the description of contact resistance in a form similar to the Kohlrausch [40] relation. The theory of this contact resistance model can be summarized as follows: •
•
•
Actual contact resistance consists of the film resistance and constriction resistance. Measured static resistance values represent mostly the film effect, which may be orders of magnitude higher than the constriction resistance and sensitive to surface condition, pressure, and temperature. The welding current breaks down the surface films during the first current cycle and film resistance becomes negligible in comparison to the constriction resistance. Because the film breakdown was associated with the melting and solidification of contact spots, fresh metal-to-metal contacts are established at the E/W and W/W interfaces. Consequently, mathematical characterization of the constriction resistance is the key to an effective model of the contact resistance, which is very dynamic in nature in the presence of high magnitude welding currents. Because of the high magnitude of welding current involved in the welding process, the heating rate at the contact interface can be extremely high. However, attempts to heat up the contact interface above the solidus temperature of the contact materials results in instantaneous collapse of the contact spot and increased contact area. Therefore the thermal gradient can be extremely high in the immediate vicinity of the contact interface and a dynamic equilibrium state can be assumed between the heat generation and heat dissipation in the immediate vicinity of the contact interface. It is assumed that the supertemperature [41] at the contact
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interface remains constant at the solidus temperature of materials in contact. When the bulk temperature rises above the solidus temperature, contact resistance disappears completely. Fundamentally, the conduction of electricity and heat obey the same form of governing differential equation and flow through essentially the same pathway. Therefore an isothermal surface may also be a surface of equal potential [40, 41]. For an infinitesimal element between two such surfaces, there exists a relation between the thermal resistance dW and electrical resistance dR: dW = dR/ρk
(38)
in which ρ and k are the electrical resistivity and thermal conductivity of the material respectively. Rewriting Fourier’s law of heat conduction, the temperature differential between the two isothermal surfaces is then obtained: –dT = qdW = ivdR/ρk = vdv/ρk
•
(39)
where i is the current and v is the electrical potential with reference to the contact interface. It is known that most pure metals obey the Wiedemann–Franz–Lorentz law: ρk = LT
(40)
where, L is the Lorentz constant and T is temperature in Kelvin. It is assumed that the materials involved in RSW also obey this law. Using the Wiedemann–Franz–Lorentz law, Eqn. 5.39 can be rearranged as vdv = –LTdT
(41)
Integrating Eqn. (5.41) over a distance across the contact interface in which the temperature reduces to the bulk temperature, the following relation is obtained
V 2 = L ( TS2 – TB2 )
(42)
where V is the voltage drop due to constriction resistance on one side of the contact member, TS is the contact supertemperature of the contact to maintain solid contact, and TB is bulk temperature of the interface. The fundamentally based contact resistance model described above has been reportedly applied to the analyses of both SSRSW and LSRSW. Agreements were reported between predicted and measured nugget growth data and dynamic resistance [42, 43]. Since this model is independent of experimental measurements and therefore avoids the uncertainty, it has been increasingly adopted in recent publications [44–46].
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5.5.4
149
Mesh structure and boundary conditions
When constructing the finite element mesh to model SSRSW, the axisymmetry of the welding set-up is always considered, and only one half of the model needs to be included in the computational region. Figure 5.13 shows the uphalf of the two dimensional finite element meshes used in simulation of a SSRSW process by ANSYS, in which three types of element are used: a thermoelectric solid element for thermal–electrical analysis; an isoparametric solid element for thermal–mechanical analysis; and a node-to-surface contact element for contact analysis at the E/W and W/W interfaces. A layer of thermoelectric solid element, with its thickness equal to a typical oxide thickness, is constructed to simulate the contact resistance at the interfaces in thermal–electric analysis. A mesh convergence study is generally required to see if the mesh is refined sufficiently. Boundary conditions are required to simulate the interactions between the workpieces, electrodes, and the surroundings during resistance microwelding. As an example, the following lists the boundary conditions used when modeling a SSRSW process with ANSYS [35]. In the thermal–electrical analysis •
The voltage at the bottom end of the lower electrode is set to zero, and the electric current is applied at the top end of the upper electrode.
5.13 Upper half of the finite element mesh used in computations.
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The current is permitted to flow across the E/W and W/W interfaces, but not allowed to flow along the lateral surfaces and centerline of the electrode. The outer surfaces are assumed to be adiabatic; and there is no radial heat flow along the centerline and no vertical heat flow along the contact area of the W/W interface because of the symmetry. The heat transfer across the E/W and W/W interfaces is calculated using the node-to-surface contact elements.
In the thermal–mechanical analysis • • •
Electrode force is applied as an evenly distributed pressure at the top end of the upper electrode. Axial displacements at the bottom end of the lower electrode are constrained. Radial displacements at the central line are restricted.
5.5.5
Examples
Comparative studies on SSRSW and LSRSW Owing to the small thickness of sheets and special process parameters employed, SSRSW has many particular features compared to LSRSW. The finite element method has been used to study the differences between SSRSW and LSRSW [34], and the dominant reason for these differences is analyzed. Figure 5.14 shows the variations in contact diameters at the W/W and E/ W interfaces for SSRSW and LSRSW under conditions listed in Table 5.5. It can be found that the contact diameters at both W/W and E/W interfaces are comparable to the electrode tip diameter in LSRSW. However, contact diameters at both interfaces in SSRSW reduce greatly and are much smaller than the electrode tip diameter. The difference between the contact diameters can be clearly observed in two graphic presentations of the deformed E/W stacks in both LSRSW and SSRSW, as shown in Fig. 5.15. The variation of contact diameter will affect the magnitude and distribution of welding current during welding. The variation in current density distribution at the W/W interfaces during SSRSW and LSRSW is plotted in Fig. 5.16. In both SSRSW and LSRSW, current density is distributed evenly at the first stage. When the nugget forms, the electric current concentrates in the nugget region. The current density in SSRSW is significantly higher than that in LSRSW because of a greatly reduced contact area in SSRSW, although the nominal current density in SSRSW is lower than that in LSRSW. Figure 5.17 shows the predicted nugget shape and size for LSRSW and SSRSW. Clearly, the nugget diameter is close to the electrode tip diameter in
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200 Workpiece/workpiece Electrode/workpiece
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5.14 Variations of normalized contact diameter at the workpiece/ workpiece and electrode/workpiece interfaces for (a) LSRSW and (b) SSRSW (d0 is the electrode tip diameter).
LSRSW while the maximum nugget diameter is only about 30% of the electrode tip diameter in SSRSW. The difference in electrode forces used in SSRSW and LSRSW is thought to be the dominant reason for other differences between SSRSW and LSRSW.
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Plate thickness Welding current (a.c. RMS) Electrode force Welding time Electrode tip diameter Base metal Electrode materials
SSRSW
LSRSW
0.3 mm 0.8 kA 50 N 10 cycles 3.2 mm AISI 1010 mild steel Class II electrodes
0.8 mm 7.0 kA 2000 N 12 cycles 6.0 mm AISI 1010 mild steel Class II electrodes
(b) SSRSW
(a) LSRSW
5.15 Deformed electrode/workpiece stacks in (a) LSRSW and (b) SSRSW.
Studies on the effects of electrode force in SSRSW Electrode force is a critical process parameter in determining the welding quality in SSRSW. The effect of electrode force on SSRSW has been numerically investigated [35]. Details of the welding process parameters are listed in Table 5.6. Figure 5.18 shows the variation of contact radius at the W/W interface. Despite basically the same varying tendency, a great difference exists in the minimum contact radius for different electrode forces. The minimum contact radius for a 50N electrode force is about 50% of the electrode tip radius, while that for 150N is about 85%. Obviously, the actual contact area increases with an increase of the electrode force. Since there is a distinct difference in the contact area under different electrode force, different welding time is required for the nugget to develop; in addition, the final nugget size is different too. Figure 5.19 shows the nugget growth curves under three levels of electrode force. Obviously, the
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Current density, A/mm2
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500 1 0 0.0
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(b)
5.16 Distribution of welding current density at the W/W interface at different weld time (in cycle at 60 Hz) for (a) LSRSW and (b) SSRSW.
higher the electrode force, the longer the time required for nugget to initiate. After initiation at different times, the nugget grows very rapidly and reaches its maximum diameter. The maximum nugget diameter formed under 150N is smaller than that formed under 50N under the same electric current. This has been attributed to the larger contact area and lower current density under higher electrode force.
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(a) LSRSW
(a) SSRSW
5.17 Calculated shape and size of nuggets for LSRSW and SSRSW. Table 5.6 Welding process parameters used in studying effects of electrode force
5.6
Process parameters
Values
Plate thickness Welding current (d.c.) Electrode force Welding time Electrode tip diameter
0.2 mm 1.0 kA 50, 100, 150 N 32–100 ms 3.2 mm
Summary and future trends
Fusion welding is a complicated technology that involves heat transfer, electromagnetic phenomenon, metallurgical process, and fluid mechanics, etc. Thanks to the fast development of computational science and technology in recent decades, numerical modeling on welding processes has made significant progress and demonstrated its value in many aspects, such as assisting in procedures optimization, residual stress and distortion control, and improving peoples’ understanding of physical phenomena. However, there is still much work to do to make the modeling results closer to physical reality. This includes: 1. Establishing a more accurate physical model based on improved understanding of the physical processes related to the welding process.
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P = 50 N P = 100 N P = 150 N 2.0
1.5
1.0
0.5 0.00
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0.03 0.04 Welding time, s
0.05
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5.18 Variation of contact radius at W/W interface under different electrode forces. 1.5
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0.04 Welding time, s
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5.19 Nugget initiation and growth under different electrode forces.
This includes, but is not limited to, research on interaction between laser and base metal, laser and plasma, keyhole behavior, etc. 2. Developing special computational theories and technologies to model more distinct features in welding, such as dynamic modeling of metal loss due to vaporization, tracing algorithms for the deformed surface of the weld pool, computational models to handle solid, liquid and gas simultaneously, etc. 3. Carrying out physical experiments to measure material properties, especially those at high temperatures. On this basis, establishing a
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material library to facilitate numerical modeling jobs and improve solution accuracy. As far as resistance microwelding is concerned, future modeling work should focus on the predictions of residual stresses, microstructure and hardness distribution in welds, so as to correlate the welding process parameters and materials metallurgy with the performance of resistance welds. These numerical works will be of great assistance in selecting the welding parameters, optimizing welding procedure, and realizing on-line monitoring and controlling of welding quality. With in-depth understanding on the physical processes associated with fusion and resistance microwelding processes and the continuous development of computational science and technology, numerical modeling will play a more and more important role in the future.
5.7
Acknowledgements
The author would like to acknowledge Professor Yongping Lei of Beijing University of Technology for providing some of his research results and Professor Yunhong Zhou, the editor of this book, for his constructive suggestions and great patience when preparing the manuscript.
5.8
References
1. K. I. Johnson (ed.), Introduction to Microjoining, TWI, Cambridge, 1985. 2. C. A. Harper, Handbook of Materials and Processes for Electronics, McGraw-Hill, New York, 1970. 3. RWMA, Resistance Welding Manual, 4th edn, Resistance Welder Manufactures’ Association, Philadelphia, PA, 1989, Chapter 14. 4. Y. Zhou, C. Reichert C. and K. J. Ely, in ICAWT ’98: ‘Joining Applications in Electronics and Medical Devices’, Columbus, OH, September 1998, 79–90. 5. Y. Zhou, P. Gorman, W. Tan and K.J. Ely, Journal of Electronic Materials, 2000, 29(9): 1090–1099. 6. H.C.-L. Chu, Proceedings of the 2nd Microelectronic Packaging Technology Materials and Processes, ASM International, Philadelphia, PA, April, 1989. 7. K. J. Ely and Y. Zhou, ‘Micro-resistance spot welding of Kovar, steel and nickel, Science and Technology’, Welding and Joining, 2001, 6(2): 63–72 8. Z. Feng, Processes and Mechanisms of Welding Residual Stress and Distortion, Woodhead Publishing Limited, Cambridge, England, 2005. 9. H. S. Cho and Y. J. Cho, Welding Journal, 1989, 68, (6), 236s–244s. 10. H. A. Nied, Welding Journal, 1984, 63, (4), 1234s–132s. 11. C. L. Tsai, O. A. Jammal, J. C. Papritan and D. W. Dickinson, Welding Journal, 1992, 71, (2): 47s–54s. 12. C. L. Tsai, W. L. Dai, D. W. Dickinson and J. C. Papritan, Welding Journal, 1991, 70, (12), 339s–351s. 13. M. Syed and S. D. Sheppard, ‘Computer simulation of resistance spot welding as a
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coupled electrical-thermal-mechanical problem’, in Modeling and Control of Joining Processes, (ed. T. Zacharia), ASM/AWS, Orlando, FL, 1993, 422–249. C. S. Wu, Numerical Analysis of Thermal Process During Welding, Harbin Institute of Technology Press, Harbin, China, 1990. D. Rosenthal, ‘Mathematical theory of heat distribution during welding and cutting’, Welding Journal, 1941, 20: 220–234. H. H. Rykalin, The Calculation of Thermal Processes in Welding, Mashgiz, Moscow, 1951. J. Li, Q. Guan, Y. W. Shi, D. L. Guo, Y. X. Du and Y. C. Sun, ‘Studies on characteristics of the temperature field during TIG welding with a trailing heat sink for titanium sheet’, Journal of Material Processing Technology, 2004, 147(3): 328–335. Y. W. Shi, Materials Welding Engineering, Chemical Industry Press, Beijing, China, 2006. W. S. Chang and S.J. Na, ‘A study on the prediction of the laser weld shape with varying heat source equations and the thermal distortion of a small structure in micro-joining’, Journal of Materials Processing Technology, 2002, 120(1–3): 208– 214. R. Brockmanna, K. Dickmanna, P. Geshevb and K. J. Matthesc, ‘Calculation of laser-induced temperature field on moving thin metal foils in consideration of Stefan problem’, Optics & Laser Technology, 2003, 35: 115–122 G. M. Oreper, J. Szekely and T. W. Eager, Metall. Trans. B, 17B, 1986, 735–744. G. M. Oreper and J. Szekely, Metall. Trans. A, 18A, 1987, 1325–1332. T. Zacharia, S. A. David, J. M. Vitek and T. DebRoy, Weld. J., 1989, 68, 499s–509s. T. Zacharia, S. A. David, J. M. Vitek and T. DebRoy, Weld. J., 1989, 68, 510s–519s. L. A. Betram, J. Eng. Mater. Technol., 1993, 115, 24–29. Y. P. Lei, Y. W. Shi, H. Murakawa, Y. Ueda, ‘Numerical analysis on the effect of sulphur content on weld pool geometry for type 304 stainless steel’, Mathematical Modelling of Weld Phenomena 4, L. Cerjak and H. Bhadeshia, (eds), Book695, The Institute of Materials, London, 1998, 89–105. Y. P. Lei and Y. W. Shi, ‘Numerical treatment of the boundary condition and source terms on a spot welding process with combing buoyancy – Marangoni driven flow’, Numerical Heat Transfer, Part B, 1994, 26: 455–471. K. Mundra, T. DebRoy, S. S. Babu and S. A. David, Weld. J., 1997, 76, 163s–171s. W. Zhang, J. W. Elmer and T. DebRoy, Mater. Sci. Eng., 2002, A 333, 320–335. W. Zhang, G. G. Roy, J. W. Elmer and T. DebRoy, ‘Modeling of heat transfer and fluid flow during gas tungsten arc spot welding of low carbon steel’, Journal of Applied Physics, 2003, 93(5): 3022–3033. W. Q. Tao, Numerical Heat Transfer, Xi’an Jiaotong University Press, China, 2001. X. He, P. W. Fuerschbach and T. DebRoy, ‘Heat transfer and fluid flow during laser spot welding of 304 stainless steel’, Journal of Physics D: Applied Physics, 2003, 36: 1388–1398. D. J. Browne, H. W. Chandler, J. T. Evans and J. Wen, Welding Journal, 1995, 74(10), 339s–344s. B. H. Chang, M. V. Li and Y. Zhou, ‘Comparisons between small-scale and ‘largescale’ resistance spot welding, Science and Technology of Welding and Joining, 2001, Vol. 6, No. 5, 1–8. B. H. Chang and Y. Zhou, ‘Numerical study on the effect of electrode force in smallscale resistance spot welding’, Journal of Materials Processing Technology, 2003, Vol. 139, No. 1–3, 635–641.
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36. M. V. Li, P. Dong and M. Kimchi, in ICAWT’97: High-Productivity Joining Processes, Columbus, OH, September, 1997, 357–369. 37. Z. Feng, J. E. Gould, S. S. Babu, M. L. Santella and B. W. Riemer, in: Proc. Conf. ‘Trends in Welding Research’, ASM, Pine Mountain, GA, June 1998, 599–604. 38. X. Sun, P. Dong and M. Kimchi, in: Proc. Conf. ‘Seventh International Conference On Computer Technology in Welding’, NIST, San Francisco, CA, July 1997, 447– 457. 39. J. A. Greenwoord, British Journal of Applied Physics, 1967, 17, 1621–1632. 40. F. Kohlrausch, Über das Problem eines elektrisch erwarmten Leiters. Ann. Phy. Lpz, 1900, 1, 312. 41. R. Holm and E. Holm, Electric Contacts: Theory and Application, 4th edn, SpringerVerlag New York Inc., 1967. 42. M. V. Li, P. Dong and M. Kimchi, in: Proc. Conf. ‘Seventh International Conference On Computer Technology in Welding’, NIST, San Francisco, CA, July 1997, 423– 434. 43. M. V. Li, P. Dong and M. Kimchi, in: Proc. Conf. ‘Seventh International Conference On Computer Technology in Welding’, NIST, San Francisco, CA, July 1997, 389– 398. 44. B. H. Chang, D. Du, B. Sui, Y. Zhou, Z. Wang and F. Heidarzadeh, ‘Effect of forging force on fatigue behavior of spot welded joints of aluminum alloy 5182’, ASME Journal of Manufacturing Science and Engineering, 2007, 129 (1): 95–100. 45. B. H. Chang, Y. Zhou, I. Lum and D. Du, ‘Finite element analysis on the effect of electrode pitting in resistance spot welding of an aluminium alloy’, Science and Technology of Welding and Joining, 2005, Vol. 10, No. 1, 61–66. 46. J. C. Feng, Y. R. Wang and Z. D. Zhang, ‘Nugget growth characteristic for AZ(31)B magnesium alloy during resistance spot welding’, Science and Technology of Welding and Joining, 2006, Vol. 11, No. 2, 154–162.
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6 Sensing, monitoring and control M M A Y E R, University of Waterloo, Canada
6.1
Introduction
As with all fields of engineering, the feedback methodologies of sensing, monitoring, and control are also applied extensively in the field of welding and joining. Feedback is inherent to the concept of quality in engineering; it is used to assure the quality of joints by comparison of the actual with the desired information about the process result. Feedback is used to better understand mechanisms of joint formation, guiding new joint design, modeling, material selection, and process improvement. This can enable inventions of new joining processes and improved materials for joining. Finally, feedback is used for process control, resulting in more efficient production of joints. The application of the concepts of sensing, monitoring, and control to a process is illustrated in Fig. 6.1. A process usually receives input parameters (variables) and produces output responses. A process consists of a set of mechanisms that interconnect the variables and produce the responses. This set of mechanisms might not be known (black box). Feedback signals can either be obtained in real-time from ‘within the process’ (in-process), or thereafter (post-process). It is understood here that sensing, monitoring, and control are methods different from but complementary to the existing more time-consuming analytical methods, such as cross-sectioning of joints, microstructure and material analysis, and others. A control method is different from pure sensing and monitoring as it closes the loop to the input side of the process. The control action involves a set of rules or an algorithm to determine a change of input parameters (adaption) based on monitoring data. In the following, feedback methodologies are discussed in more general with respect to microjoining and nanojoining. Definitions of frequently used expressions are given, and basic measurement methods are described. Signal types useful for joining processes are classified. Example applications are summarized, and an outlook on the future of sensing, monitoring, and control in microjoining and nanojoining is given. 159 WPNL2204
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Microjoining and nanojoining Process mechanisms Input parameters (variables)
Post-process
Adaption
In-process
Output values (responses)
Control (algorithm)
Sensing/monitoring
6.1 Concept of process, sensing, monitoring, and control. Process mechanisms illustrated by arrows interconnected with another and connecting process variables (settings) to process responses (results).
6.2
Definitions and methods
According to the Merriam-Webster on-line dictionary, a sensor is ‘a device that responds to a physical stimulus (as heat, light, sound, pressure, magnetism, or a particular motion) and transmits a resulting impulse (as for measurement or operating a control)’. Sensors are a subset of transducers. It ‘leads across’ (latin: trans ducere) from an input to an output. The resulting output impulse is the ‘sensor value’ or ‘measurement data’ transmitted to a display or recording device. In most industrial application, the sensor value is a voltage recorded by, for example, a data-acquisition (DAQ) electronic board inside a computer. A time series of measurement values is usually called a ‘signal’. A voltage recording device can be characterized by its range and resolution. A range of 0 to 10 V means that the only signal voltages measurable are voltages larger than or equal to zero and smaller than or equal to 10 V. The resolution of a DAQ board depends on the analog-to-digital (AD) converter element. For example, if a commercial DAQ board has a 12-bit AD converter, there are 212 value steps available in the measurement range. If, for example, a range of 10 V is selected, the corresponding resolution (step size) is 10 V/ 212 = 2.44 mV. The corresponding resolution for a 16-bit card would be 0.153 mV. The precision (repeatability) of a recording device is its ability to reproduce a measured value. The precision of recorded signals suffers from noise arising from a variety of different sources. Noise causes the recorded signal to vary
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by random values around a long-term average value. Accuracy and precision are not the same. The accuracy of a recording device is its ability to provide measurement signals the average of which over a long enough time is close to the correct (‘true’) value. Usually, a calibration procedure is carried out to improve and optimize the accuracy of a given device. The measured value is adjusted by means of a calibration factor to the value of using a known signal level (standard). An uncalibrated device can be inaccurate but still be very precise. On the other hand, a device can be very accurate and at the same time not very precise, in which case it requires a large number of measurement values (a large sample size) to obtain an accurate average. Monitoring is the act of repeated sensing and recording values taken from a process, as, for example, a production process. In most cases it is used with a control action. The production manager sooner or later analyzes the monitoring data and decides (and this is a control action) if production goes on or is stopped to introduce changes to the process. Two basic monitoring types are sampling and screening. Sampling usually applies to parts being manufactured, and is characterized by the fact that not all parts are sensed. If production is organized in lots, each consisting of a certain number of parts, sampling would only measure values from a fraction of all parts of the lot. This can save time and money and still assure the desired quality level. In contrast, screening measures all parts, making sure only acceptable pieces reach the customer. To improve the efficiency of sampling or screening monitoring, various standards and statistical methods are applied [1]. A control method is a monitoring method combined with a predictive control action that induces changes (‘controls’) to at least one process input. In combination with statistical methods, control becomes statistical process control (SPC), and, if applied correctly, increases the efficiency of a production process. Control methods can be divided into run-to-run methods and realtime methods. In the first a post-process (‘after the run’) measurement is used to define the control action, in the latter an in-process measurement is typically used. A very basic control method might be to run production while monitoring the company’s bank account. The monitoring action would be to check if the profit materializing on the account falls under a predefined value, caused, for example, by a problem with a customer who turned unhappy and stopped paying. The correct control action in such a case would be to fix the problem in a way that is predictive of the customer starting to pay the bills again. More advanced control methods use control charts [2] with moving averages, upper and lower specification limits, and tables describing what to do and when to do it, including corrective actions that are predictive of the process getting back under control again. Examples for real-time control methods include proportional-integral-derivative (PID) methods as used for temperature
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control or motion trajectory control of a robot arm. Artificial neural networks can be applied for real-time control [3].
6.3
Signals from joining processes
Out of the six different kinds of signals [4] – mechanical, thermal, magnetic, electric, optical, and chemical, all except chemical and magnetic signals seem particularly suitable for characterizing joining processes. Mechanical signals include all destructive tests, e.g. tensile tests, that actually break the joints. Many joining processes involve heat, therefore temperatures or their differences measured close to the process are signals significant to the process result. An example for an electric signal describing a metallic joint is the joint’s contact resistance. The class of signals most apparent to humans are the optical signals. Important information about the joint quality can be obtained in most cases just by looking at the joint. Optical signals can contain information about the joint geometry which includes its dimensions and location, and about the joint surface, e.g. its color or roughness. Optical signals can be obtained from non-visible radiation such as X-rays or ultrasound, by use of intermediate devices. Some more specific details about mechanical, thermal, and optical (geometrical) signals are given in the following three subsections.
6.3.1
Mechanical signals
Mechanical signals include stress, strain, force, and pressure. However, these signal types are interrelated, and the type of calibration procedure determines if the signal has a unit of stress, force, or pressure. Strain gages have many applications to measure mechanical signals. They are robust and can be arranged in Wheatstone bridge configurations. Consequently they are tailored to a specific application, correcting for interfering inputs such as temperature [5]. Piezoelectric sensors produce a charge (a transient voltage) when experiencing a stress change and are usually used in combination with a charge amplifying signal conditioning circuit. Piezoresistors based on silicon are resistors that change their resistance upon a stress change. They are used as strain gages with high gage factors. When applying such sensors to microjoining or nanojoining, the mounting of the sensor is an engineering challenge. Film deposition with subsequent patterning can be used provided these processes are compatible with the restrictions due to the joining application.
6.3.2
Thermal signals
Thermocouples, thermistors, and resistive temperature detectors (RTDs) are sensor types used to measure temperature. Thermocouples are a widely used
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type of temperature sensor. They work on the basis of the thermoelectric effect (Seebeck effect), which causes a metal subjected to a thermal gradient to generate a voltage. As the magnitude of this effect depends on the type of metal, two wires of dissimilar metals can be used by connecting their ends to one another, creating two points of contact of dissimilar metals. One of the wires is cut, and the two new ends are connected to a voltmeter. The measured value is proportional to the temperature difference between the two points of contact of the two metal wires. Therefore, a thermocouple measures temperature differences. Thermistors and RTDs consist of a resistor and are based on the temperature dependence of its resistance. They are more precise and accurate than thermocouples and can measure absolute temperature. A thermistor is based on a semiconductor, an RTD on a metallic resistor. A popular RTD is the ‘Pt100’, a platinum thin film resistor with a nominal resistance of 100 Ω at 0 °C. In contrast to that of other metals, the resistance change of Pt is substantially proportional to temperature in a range of at least –200 °C to +700 °C. The accuracy of a Pt-100 depends on its class. A class B Pt-100 is accurate to less than 1 K in a temperature range from –200 °C to +300 °C.
6.3.3
Geometrical signals
Geometrical signals describe either the position or the dimension of a part as it is seen by, for example, a camera. When using tools fixed to motors to hold, form, or deform materials during a joining process, the motor position as, for example, obtained using an encoder, can yield a signal about the degree of plastic deformation. In combination with cameras and machine vision algorithms, contrast differences can be evaluated on images taken, and features can be compared with previously stored information. The machine vision algorithm yields position or dimension of a feature identified on the image. Placement and dimensional values can be measured considerably more accurately than the pixel size of the image would suggest. Methods based on machine vision are often applied to closed-loop feedback control systems, as for example reported in [6].
6.4
Applications
Examples of microjoining and nanojoining applications of sensing, monitoring, and control are described in this section. They are classified in research and development applications and production process control applications. They are taken from the fields of microresistance welding, microelectronics wire bonding, and surface-mount technology (SMT) soldering.
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6.4.1
Microjoining and nanojoining
In-process monitoring
In-process signals are obtained either from sensors or from an electrical control circuit. If these in-process signals correlate with a quality response, they can be used for quality monitoring. A basic methodology to do this is to first define a reference range in a way that if the signal remains inside the range, the quality response is acceptable. After each monitoring action, the signal is checked with respect to the reference range. If it is outside the range, the production or the process is interrupted. More sophisticated versions of this methodology use formulas or other models to deduce a quality value from the monitoring signals. In [7] for example, the dynamic resistance across electrodes in a resistance spot welding (RSW) process is monitored. The dynamic resistance signal is closely related to the nugget formation in the weld. Characteristic values of this signal are evaluated and brought into relation with the final weld strength by training a neuronal network. After the nodes of this network are weighted by a correct training procedure, the network can be used in real-time to predict the weld strength using the characteristics from the dynamic resistance monitoring signal. A number of welds need to be carried out first together with the destructive testing to obtain data to train the network. In another example [8], machine mounted sensors and analytical theories (a set of analytical models) are used to monitor the RSW process. Analytically determined in-process parameters and responses are compared with target values that are references for good weld quality. The analytical values include contact area, dynamic resistance, joule heating rate, total heat produced during welding, thermal expansion (electrode displacement), nugget geometry, and temperature distribution. The monitoring method enables an improved defect weld detection compared to off-line test coupon methods. In microresistance welding, material combinations different from steel are very common. The dynamic resistance is monitored during µRSW of 200 µm thick Ni sheets [9]. The process is called small scale resistance spotwelding (SSRSW). The dynamic resistance differs significantly from that observed in larger scale steel sheet RSW. Apparently, a different set of process mechanisms are taking place in Ni sheet RSW. Reported process stages in chronological order are asperity heating, surface breakdown, asperity softening, partial surface melting, nugget growth and expulsion. Wire bonding processes with ultrasound are used to electrically connect microcircuit devices with fine wires to substrate metallizations or leadframes [10]. The wire diameters are in the range of about 12.5 µm to 50 µm. For power microcircuits, larger diameter wires are used. As most microcircuits are interconnected with this process, the number of such wire joints manufactured is probably by far the highest of all types of joints made annually. The ultrasonic vibrations are produced by a stack of piezo-electric
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disks attached to a horn that amplifies the vibration amplitude which is highest at its tip. There, a capillary tool is fixed that guides the fine wire to the bond location. The bonding process uses ultrasonic vibration and normal clamping force for typically a few tens of ms until the wire is metallurgically connected to the bond location. A material combination common in wire bonding is Au and Al. In the thermosonic version of the process, a ball is formed at the end of the gold wire, and subsequently pressed and bonded to an Al pad of the microcircuit. To assist future developments of improved bonds with improved materials, testchips are produced with CMOS technology [11]. Such testchips contain custom designed test pads with stress and temperature sensors co-integrated next to them (in situ). These microsensors have been used to identify the mechanisms responsible for successful joining [12, 13, 14]. The temperature microsensors are RTDs fabricated from Al lines with the minimum possible line width, about 1 µm in older CMOS technologies. The Al lines are designed in a way to minimize the average distance to the testpad [15, 16] resulting in typical resistances in the range of 10 to 100 Ω. The materials used for the stress microsensors are the p+ and n+ diffusions which are usually used for the CMOS transistors. When used as resistors, these diffusions change their resistance upon a change in the stress state. This is the piezoresistive effect. Compared to metallic strain gages, diffused silicon exhibits gage factors that are one to two orders of magnitude larger. Signals from such piezoresistive microsensors recorded during wire bonding are applied in industry in wire bonder characterization and trouble shooting. Piezoresistors are placed below [13, 17] and next to [18, 19, 20] test bonding pads. An example of the ultrasonic stress signal measured by the sensor presented in [13] is reproduced in Fig. 6.2. Several applications for wire bonders with such microsensors are described in [21], including the calibration of the ultrasound and force parameter, the characterization of the bond head vibrations, and the longterm testing of the process parameters. Wire bonding is a fully automated high speed production process. Production rates of more than 50 bonds (25 wires) per second can be demonstrated on modern equipment [22]. Naturally, monitoring techniques play a key role in assuring high enough device yields. A device yield of, for example, 99.9% means that 0.1% of the devices or 1000 parts-per-million (ppm) are defective. Assuming a device with 300 wire loops, each loop with two bonds, such a high device yield can only be obtained if the defect rate of the wire bonds is lower than 1000/600, i.e. lower than 1.67 ppm. This corresponds to a wire bond yield of higher than 99.9998%. To achieve such high yields, modern wire bonders apply several monitoring technologies. A group of them are the electrical non-stick detection methods. After the first bond of a wire loop, a pre-defined d.c. or a.c. voltage is applied between wire and stage and the current is measured. If the bond connection is not established, the measured
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3
Tool lift off Signal break-off
Microsensor signal (mV)
2
1
0
–1
–2
Start of ultrasound End of ultrasound
–3 0
4
8 Time (ms)
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6.2 Reflow soldering: dimensional monitoring of packaged chip. Displacement measurement system described in [24].
current is outside a specified range, the bonder stops, and the operator assists the machine. After the second bond of the loop, a similar monitoring step is carried out. Finally, after breaking the wire, the measurement is carried out a third time to monitor if the breaking process was successful. In case it was not, the wire is still connected to the substrate. For this third test, the monitored value needs to be in another specified range to assure that the wire has been broken. The soldering of microelectronic components onto substrates such as printedcircuit boards (PCBs) has become a mainstream technology of second level electronic packaging [23]. In soldering, the components are thermomechanically stressed as they are heated up to 40 K above the melting temperature of the solder. With the advent of directives and legislations addressing environmental aspects of microelectronics, the thermomechanical stresses increase due to the replacement of lead-based solders with higher melting temperature alternatives. One of the directives issued by the European Union is the Restriction of Hazardous Substances (RoHS) directive, leading to the replacement of tin-lead solders with, for example, Sn-Ag-Cu (SAC) alloy solders that have an about 30 K higher melting temperatures than the Sn-Pb eutectic. Too high levels of thermomechanical stresses are undesirable as they promote the formation of cracks and delaminations. In the case of a plastic encapsulated microcircuit (PEM), moisture can be present inside the polymer molding compound. Possible failure mechanisms are observed during soldering.
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The soldering process vaporizes the moisture fast enough to prevent its diffusion outside of the polymer bulk, resulting in large deformation and cracking of the package or in delamination of the polymer from the semiconductor surface. The monitoring approach adopted in [24] involves the application of a touching probe acting as a real-time displacement measurement system to measure the package deformation at one point of the surface. The system consists of two cantilever beams with strain gages mounted in a Wheatstone bridge arrangement, touching the package at two points opposite each other, as illustrated in Fig. 6.3. The objective is to obtain fundamental insights into defect mechanisms in the package during reflow. The monitoring method delivers time resolved deformation signatures of packages for the applied reflow profile. The signatures contain information about the occurrence of delamination or cracking within the package that can be confirmed by post-reflow inspection using scanning acoustic microscopy (SAM). Using this method, the time and temperature during defect formation can be obtained.
6.4.2
Process control
Automatic wire bonders have machine vision systems for recognizing the positions of the chips and substrates with respect to the bonding tip. Such alignment systems are able to achieve bond placement accuracies of 2 µm or better at the three sigma level. An alternative use of the bonder’s pattern recognition system is to check and correct the positioning of bonded balls on the pads. After bonding a set of bonds, a specified subset is checked and the placement is corrected if needed before bonding the next set. This method can significantly increase the placement accuracy, however it reduces the throughput. The idea of in-process deformation monitoring of bonding wires being welded to substrates can lead to higher yields in production. Such methods Upper cantilever beam with strain gage Microcircuit Polymer
Defect
Lead Lower Displacement
Probe tip
6.3 Ultrasonic microsensor signal of tangential stress on pad during ball bonding. High pass filter. From [13].
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are motivated by the assumption that the bond quality correlated with the degree of ultrasonically induced deformation. One example uses the realtime z-position signal of the bondhead, and modulates the ultrasonic signal of the ultrasound generator. These signals together with the bonding force signal are shown in Fig. 6.4. The real-time horn position signal can in some cases be obtained by a proximity sensor next to the horn [25, 26]. In production a defined deformation endpoint detection mechanism is provided, as shown schematically in Fig. 6.5. This deformation control method results in more uniformly shaped wire bonds. The method can lead to a significant reduction of the pull force variations for an example wedge bond process [27] if a constant force parameter is assured. If this is not the case, the same wire
Wire deformation
1
2
3
Force parameter profile
Ultrasound parameter profile
Time
6.4 Wire deformation, force, and ultrasound parameter profiles during wedge bonding. Numbers 1–3 denote cleaning, bonding, and heel weakening phases needed for successful wedge bonding. Adapted from [27].
Deformation endpoint
Time
0 Ultrasound
6.5 Closed-loop wire deformation endpoint detection and subsequent stop of ultrasound.
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deformation occurs for bonds with different strength [28]. This method is currently available on modern bonders. A method similar to deformation monitoring uses a laser dopplerinterferometer to measure the in-process attenuation of the tool tip amplitude of a wedge–wedge bonder [29]. This attenuation is caused by the increasing bond strength. The implementation of this method in production has not been reported. The actual friction force is measured by this method. The increase of the frictional force corresponds to the reduction of the vibration amplitude. From [30] it follows that this monitoring value correlates with the actual bond strength only if the bonding force is constant. For a specialized transducer system, the amplitude reduction during bonding can be monitored using a piezoelectric sensor placed on a correct position on the transducer [31, 32]. An in-process method to directly measure the shear force of a bonded ball similar to that of the standard shear test, uses a horizontal bond head motion instantly after the bonding took place, to shear the ball from the pad [33–36]. In this bond head shear test, the capillary still holds the ball as during bonding but due to the horizontal bond head motion shears it away from the pad. During this shearing action, the electromotive force is available from the motor control board and stored on the bonder. After the shearing, the bond head moves back to the original bonding position, rebonds the sheared ball and continues with the next bonds. If used during production, this method reduces the throughput even if only a fraction of the bonded balls are tested. However, this shear force method is believed to deliver reliable in-process ball bond quality values provided the bonding force is constant, and has been demonstrated to detect a decrease in material quality [33].
6.5
Future trends
Driven by performance and cost optimization of commercial products, continuous efforts in research and development will expand the range of applications of sensing, monitoring, and control, leading to further progress in all fields of microjoining and nanojoining. In particular, the efforts will be focused on new sensors providing in situ and in-process feedback. For example, embedded sensors and microsensors are expected to be applied in the future for microjoint reliability monitoring. Methods to record real-time reliability data during accelerated aging or cycling will be developed. For microelectronics wire bond reliability, the focus might be on the development of integrated sensors measuring changes of the contact zone stress measured next to bonded balls. Simultaneously, the resistance of ball bonds could be measured using a four-wire method as reported in [37]. Correlations between microsensor signals recorded during bonding and the in situ signals recorded during reliability testing could be investigated as well as the correlation of these
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signals with conventional responses. A different type of reliability feedback is expected from such sensors, adding transient in situ signals to the conventional pass or fail information. The expected progress in microjoint reliability could include improved understanding of failure mechanisms and subsequently methods to control and prevent such mechanisms. Progress could be enhanced by the reduction of the sample size required in reliability tests and the acceleration of such tests. Nanosensors are promising candidates for novel applications and could bring billions of dollars or revenue in a few years. Potential applications are in the fields of defense, semiconductor, and biomedical engineering. Security devices based on nanosensors will revolutionize the detection of bio-threats and hazardous materials resulting in huge medical benefits. Perhaps there will be a real life tricorder to sense biological signals such as diseases from a location remote to the infected person or animal. Similarly, implantable nanosensors might be able to give us an early warning when we are about to catch a cold or if another illness is about to break out, allowing us to meet the threat early by appropriate measures. Since nanosensors are little, their sensitivity might be reduced, a fact that might be compensated by implanting arrays of nanosensors or have large numbers of them floating in the bloodstream. Other new applications could include high-performance but low-power mobile computing devices. A review of reports on nanosensor and nanowelding research leads to speculations about how these two fields might be connected. Potential applications of sensing and monitoring of nanojoining processes include the fields of process research and development, nanowelding equipment characterization, and nanowelding production monitoring. Popular among the nanocomponents are carbon nanotubes (CNT) as they attract great attention from the international research communities due to their unusual material parameters and potential applications. Nanowelding of single-walled carbon nanotubes (CNTs) is reported in [38]. The e-beam irradiation of a special high-voltage scanning electron microscope (SEM) and its heating effect were used to create just enough defects for these bonds to form without damaging their electrical properties. An example nanowelding process reported in [39] involves 1.6 nm diameter carbon nanotubes (CNTs) grown on a [111] Si substrate. A scanning probe microscope (SPM) with a nano-tip is used to position the tip within nanometer range of the CNT and then apply voltage between tip and CNT on the Si. In an atmosphere with 30–50% humidity, conductive atomic force microscopy (AFM) oxidation can be achieved, forming Si oxide by the electrical discharge. The oxide grows around the CNT and fixes it to the Si substrate. Such discoveries pave the way to efficiently handle and fix materials such as DNA molecules to a Si substrate and enable controlled fabrication of molecular circuits and nanotube networks. For such future projects, the nanowelding processes need to be characterized and optimized. To this end,
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a nanosensor providing feedback of the temperature produced during discharge might be useful. Such a sensor has been reportedly used in an SEM [40]. In this report, integrated circuit technology is combined with focused ion beam (FIB) machining to produce integrated Au-Ni thermocouples with contact areas of 200 nm × 100 nm down to 50 nm × 30 nm. One conclusion is that the sensitivity of the smaller samples is about 3 times less, whereas the signal-to-noise ratio is about 10 times smaller. Nevertheless, the smaller sensor is applied successfully to measure the temperature change caused by the electron beam scanning over the sensor with a speed of 20 µm/s. The maximum temperature increase monitored is 70 K. Such high temperatures need to be taken into account when handling bio-materials. In [41], an alternative design for a temperature nanosensor is reported, which can be applied as an actuating structure, too. The design is based on, for example, a Au/Cr-to-Ni thermocouple placed on a microtip normally used in a scanning thermal microscope (SThM), a device based on the AFM concept. The thermocouple has a junction diameter of less than 100 nm and a sensitivity of more than 10 µV/K. The primary application of the nanodevice is writing and reading in data storage applications. In the writing mode, a few mA current pulses are injected through the device which has an electrical resistance of approximately 100 W, resulting in power pulses of 10 to 100 µW. However, a voltage of 5 Vp-p can be applied for several ms without destroying the device, resulting in powers in the hundreds of mW range. Such heating powers might find applications in the nanowelding field.
6.6
References
1. D. C. Montgomery, Introduction to Statistical Quality Control, 4th edn., John Wiley & Sons, Inc., New York, 2001. 2. W. A. Shewhart, Economic Control of Quality of Manufactures Product, Van Nostrand, Toronto, The Bell Telephone Laboratories series, 1931. 3. R. L. Mahajan, ‘Process modeling, optimization and control in electronics manufacturing’, Chapter 5 in Y. C. Lee and W. T. Chen (eds)., Challenges in Electronics Manufacturing, Chapman & Hall, London, pp. 185–220, 1998. 4. R. Pallás-Areny, J. G. Webster, Sensors and Signal Conditioning, John Wiley & Sons, Inc., New York, 1991. 5. E. O. Doebelin, Measurement Systems: Application and Design, 4th edn., McGrawHill, New York, 1990. 6. M. C. Hellinga, J. P. Huissoon, H. W. Kerr, ‘Identifying weld pool dynamics for GMA fillet welds’, Science and Technology of Welding and Joining, 4(1), pp. 15–20, 1999. 7. Y. Cho, S. Rhee, ‘New technology for measuring dynamic resistance and estimating strength in resistance spot welding’, Meas. Sci. Technol., 11, pp. 1173–1178, 2000. 8. Herman A. Nied, Stanley J. Godwin, Robert K. Cohen, Robert V. Klint, Hsin-Pang Wang ‘Resistance spot welder process monitor’ US Patent number: 4596917, Issue date: Jun. 24, 1986.
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9. W. Tan, Y. Zhou, H. W. Kerr, S. Lawson, ‘A study of dynamic resistance during small scale resistance spot welding of thin Ni sheets’, Journal of Physics D Applied Physics, Vol. 37, pp. 1998–2008, 2004. 10. G. G. Harman, Wire Bonding in Microelectronics, 2nd edn., McGraw-Hill, NewYork, 1997. 11. M. Mayer, ‘In situ monitoring of microjoining processes’, Proc. Inter-University Research Seminar IURS 2005, Aug. 15–16, Tsinghua University, Beijing, China, pp. 97–102, 2005. 12. J. Schwizer, ‘In-Situ Wire Bond Prozessuntersuchung mit integrierten piezoresistiven Mikrosensoren’ (in German), diploma thesis, ETH Zurich, 1999. 13. M. Mayer, J. Schwizer, O. Paul, D. Bolliger, H. Baltes, ‘In-situ ultrasonic stress measurements during ball bonding using integrated piezoresistive microsensors’, Proc. 1999 Intersociety Electron. Pack. Conf. (InterPACK99), pp. 973–978, 1999. 14. M. Mayer, ‘Microelectronic bonding process monitoring by integrated sensors’, Ph.D. thesis, No. 13685, ETH Zurich, Zurich, 2000; also: Hartung-Gorre, Konstanz, Germany, 2000. 15. M. Mayer, O. Paul, D. Bolliger, H. Baltes, ‘In-situ calibration of wire bonder ultrasonic system using integrated microsensor’, Proc. 2nd IEEE Electr. Packaging Technol. Conf. EPTC’98, pp. 219–223, 1998. 16. M. Mayer, O. Paul, D. Bolliger, H. Baltes, ‘Integrated temperature microsensors for characterization and optimization of thermosonic ball bonding process’, IEEE Trans. Comp. Packaging Technol., vol. 23, no. 2, pp. 393–398, 2000. 17. J. Schwizer, M. Mayer, D. Bolliger, O. Paul, H. Baltes, ‘Thermosonic ball bonding: friction model based on integrated microsensor measurements’, Proc. 24th IEEE/ CPMT Intl. Electronic Manufacturing Technology Symposium IEMT’99, Austin, Texas, Oct. 18–19, pp. 108–114, 1999. 18. J. Schwizer, M. Mayer, O. Brand, H. Baltes, ‘Analysis of ultrasonic wire bonding by in-situ piezoresistive microsensors’, Proc. Transducers ‘01/Eurosensors XV, pp. 1426– 1429, 2001. 19. J. Schwizer, M. Mayer, O. Brand, H. Baltes, ‘In situ ultrasonic stress microsensor for second bond characterization’, Proc. Intl. Symp. Microelectronics IMAPS 2001, pp. 338–343, 2001. 20. J. Schwizer, Q. Füglistaller, M. Mayer, Michael Althaus, O. Brand, H. Baltes, ‘MEMS system with multiplexer for in situ and real-time wire bonding diagnosis’, Proc. SEMI Technical Symposium, Advanced Packaging Technologies I, SEMI Singapore, pp. 163–167, 2002. 21. J. Schwizer, M. Mayer, O. Brand, Force Sensors for Microelectronic Packaging Applications, Springer Science+Business Media, Series: Microtechnology and MEMS, 2005. 22. M. Barp, D. Vischer, ‘Achieving a world record in ultra high speed wire bonding through novel technology’, Proc. Electronics Manufacturing Technology Symposium IEMT 2002, pp. 342–347, 2002. 23. R. R. Tummala, E. H. Rymaszewski, A. G. Klopfenstein (Eds), Microelectronics Packaging Handbook, Part I, 2nd edn., Chapman & Hall, New York, 1997. 24. M. G. Pecht, A. Govind, ‘In-situ measurements of surface mount IC package deformations during reflow soldering’, IEEE Trans. Comp. Packaging Manuf. Technol., Part C, vol. 20, no. 3, pp. 207–212, 1997. 25. F. Farassat, ‘Method and apparatus for the production and quality testing of a bonded wire connection’, US Pat. Appl. 2004221653, Published on November 11, 2004.
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26. F. Farassat, ‘Wire bonding ultrasonic control system responsive to wire deformation’, US Pat. 5314105, Issued May 24, 1994. 27. F. Farassat, ‘Entwicklung und Erprobung eines Regelsystems zur Verbesserung der Verbindungsqualität beim Ultraschallbonden’ (in German), Ph.D. thesis, Technical University Berlin, 1996. 28. Y. Zhou, X. Li, N. Noolu, ‘A footprint study of bond initiation in gold wire crescent bonding’, IEEE Trans. Comps. Packg. Technol., Vol. 28. No. 4, pp. 810–816, 2005. 29. U. Draugelates, K. H. König, ‘Erhöhung der Verfahrenszuverlässigkeit beim Ultraschallbonden durch eine prozeßintegrierte Regelung der Werkzeugschwingung’ (in German), Verbindungstechnik in der Elektronik, 1, pp. 24–30, 1995. 30. M. Mayer, J. Schwizer, ‘Ultrasonic bonding: understanding how process parameters determine the strength of Au-Al bonds’, Proc. Intl. Symp. Microelectronics IMAPS’02, pp. 626–631, 2002. 31. M. J. Hight, R. V. Winkle, J. R. Dale, ‘Ultrasonic bonding apparatus’, US Pat. 4040885, Issued on August 1977. 32. L. W. Chan-Wong, S. S. Chiu, S. W. Or, Y. M. Cheung, ‘Piezoelectric sensor for measuring bonding parameters’, US Pat. 6279810, Issued on August 28, 2001. 33. J. Medding, M. Mayer, ‘In situ bond head shear test for wire bond production process control’, Technical Program of SEMI Singapore’04, Singapore, 2004. 34. J. Medding, M. Mayer, ‘Ball bond process optimization using in situ ball shear force’, Proc. Electronic Packaging Technology Conference EPTC’03 (IEEE), Singapore, pp. 781–784, 2003. 35. J. Medding, M. Mayer, ‘In situ ball bond shear measurement using wire bonder bond head’, Intl. Manufacturing Technology Symposium (IEMT’03), San Jose, USA, pp. 59–63, 2003. 36. M. Mayer, J. Medding, ‘Method for determining optimum bond parameters when bonding with a wire bonder’, US. Pat. Appl. 2004079790, Published on April 29, 2004. 37. B. Krabbenborg, ‘In-situ monitoring of bond degradation in power ICs under highcurrent stress’, Proc. IEEE Reliability Physics Symposium Proceedings, pp. 238–247, 1998. 38. ‘Nanowelding – Creating Tiny Junctions’, http://www.rpi.edu/web/News/research/ tip_sheets/jan_03/nanowelding.html, website of Renssealer Polytechnical Institute, retrieved on May 7, 2007. 39. X. Duan, J. Zhang, X. Ling, Z. Liu, ‘Nano-welding by scanning probe microscope’, J. Am. Chem. Soc., 127 (23), 8268–8269, 2005. 40. D. Chu, D. T. Bilir, R. F. W. Pease, K. E. Goodson, ‘Thin film nano thermocouple sensors for applications in laser and electron beam irradiation’, Proc. 12th Intl. Conf. TRANSDUCERS, Vol. 2, pp. 1112–1115, June 8–12, 2003. 41. D. W. Lee, T. Ono, M. Esashi, ‘Fabrication of thermal microprobes with a sub-100 nm metal-to-metal junction’, Nanotechnology, vol. 13, no 1, pp. 29–32, 2002.
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7 Assembly process automation and materials handling Y M C H E U N G and D L I U, ASM Assembly Automation Ltd, Hong Kong
7.1
Introduction
The yield and quality of a product can be improved by using automated assembly equipment resulting in the elimination of any inconsistency introduced by human factors. A fully automatic production line can be created by careful selection of equipment and control of the quality of incoming materials. The equipment requirements for the assembly process should be well understood so that they can be translated into procurement specifications. A production line can be arranged in either batch process mode or in-line process mode. For batch process operation, the allocation of the equipment for certain assembly processes may be more flexible, but it needs more management resources to ensure appropriate allocation of the equipment and high utilization. On the other hand, for in-line process operation the production line is built to include all the required assembly processes so that the layout of the production line is optimized for a given product. The management resources for scheduling equipment allocation will therefore be minimal. The operator just has to load raw materials to the input port at the front end of the production line and collect the end products from the output port at the back end. It is easier for the production supervisor to manage and maintain the throughput of given products on an in-line process operation. It is the objective of this chapter to discuss the considerations for implementing an automatic assembly process production line. Since batch mode operation has been adopted in most traditional production lines for product assembly processes, it is quite established and we shall not discuss this mode of operation here. However, we shall discuss the basic considerations for constructing an in-line process production line and the general requirements of the assembly equipment used for the in-line process in Section 7.2. ASM IDEALine is used as an example of an electronic packaging and assembly production line for in-line processes. Another major consideration for running an automatic assembly production line is the way the processing materials are handled. Section 7.3 will discuss material handling requirements for 174 WPNL2204
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automatic assembly equipment. Optimizing the material flow can help improve the assembly yield and throughput of the production line. When considering the type of equipment for an automatic assembly production line, the process control capability provided by the equipment and the requirements of the assembly processes should be understood in order to ensure a high production yield. Most of the processing parameters for assembly processes will be discussed in Section 7.4 and it will be helpful for the reader to understand the various roles of the assembly processes and be able to select the correct equipment for the task. The detail design procedures for building customized assembly equipment to perform specific assembly processes, namely active alignment and laser spot welding for optoelectronic packaging applications is discussed in Section 7.5. This case study shows how important the assembly specifications of the final product are in contributing to the design requirements and specification of processing capability for assembly equipment.
7.2
Assembly equipment for an in-line process
In order to construct a fully automatic assembly production line for an inline process dedicated to a given product, the yield, up-time and reliability of the equipment must be very high in order to facilitate the operation. The speed of materials flow and the product identity in the line are monitored and controlled by a cell controller which synchronizes all the process blocks in the line. The process data associated with the assembled product will also be downloaded from the cell controller to the related process equipment. With the help of the cell controller the batch or log identification of the product as well as up-to-date known good unit information and processing materials (leadframes, substrates and the dice) will be delivered to the corresponding assembly process equipment along the production line. In addition to the requirements of the assembly process, another factor affecting the configuration of a dedicated in-line process production line is the ‘line balancing’ of the assembly equipment. A balanced production line means that the throughput of each group of process equipment is basically the same. For a balanced production line, the throughput of each process block (for which a group of equipment performs the same process) is basically the same as throughout of the production line. Defining a configuration to maximize the throughput of a balanced in-line production line is one of the major considerations for optimizing the capital investment when building the production line. ASM IDEALine is an example of an in-line process production line used for packaging and assembly of electronic integrated circuits (ICs) and one of the most common configurations is shown in Fig. 7.1. A cell controller links up all the assembly equipment in the production line through SECS communication (an industrial standard for communication between equipment).
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Electronic packaging assembly process flow Epoxy curing
Wire bonding
Moulding
Post-mould cure
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Die attach
Trim and form
Buffer station
Moulding machine
Cell controller
Die attach bonder
Epoxy curing oven
Post-mould curing oven
Trim and form machine
Wire bonder
7.1 The flow chart of electronic packaging assembly process and the layout of ASM IDEALine for an electronic packaging in-line process assembly line which include a cell controller, a die attach bonder, an epoxy curing oven, a buffer station, four gold wire bonders, another buffer station, a moulding machine, a post-mould curing oven and a trim and form machine (courtesy of ASM Pacific Technology).
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The cell controller controls the speed of material flow along the assembly line. The operator loads the wafers and leadframes into the die bonder. The diced wafer mounted on the wafer ring will be transported automatically from the magazine onto the wafer expander of the wafer table by the robotic arm. In the meantime, the leadframes will be picked and placed into the workholder by the vacuum pick arm of the leadframe loader. The leadframe will then be held by an indexer and transported to the epoxy dispensing stage by a motorized indexer. An appropriate amount of epoxy will be dispensed on the die pad of the leadframe. The vision alignment optics will search for the fiducials or reference marks on the leadframe so that the leadframe will be aligned and the epoxy will be dispensed to the right location. With the help of the ejecting pins of the die detachment module, the target die on the wafer table will be picked up by the bondhead of the die bonder. The die will be transported and bonded to the leadframe by the high-speed bondhead of the die bonder. The pickup of the die from the wafer and placement of the die on to the substrate is performed accurately with the help of vision alignment optics. The bonding accuracy is better than 20 µm. A barcode reader mounted on top of the wafer table is used to scan the serial number on the wafer. The die bonder sends a request and downloads the test map of the corresponding wafer from the cell controller. Based on the information from the test map, only known good dice will be picked and bonded onto the leadframe. After the die-attaching process, the leadframes with bonded chips will be transported to a curing oven which usually consists of multiple heating and cooling zones. For in-line process configuration, snap-cured epoxy is commonly used to reduce the curing time of the epoxy. These leadframes are temporary stored in a buffer station after epoxy curing. The buffer station sends one strip of leadframe at a time to next available wire bonder for wire bonding. In most cases, the time needed for wire bonding one unit is much more than the time for die attachment. One die bonder can be linked up to four to six wire bonders with their throughput balanced. Automatic wire bonding is an extremely challenging process [1, 2]. Gold wires of diameter ~ 20 µm are bonded onto small bond pads of size < 50 µm on a silicon chip. In order to achieve ultra-fine bond pad pitch that is less than 50 µm, ultra-low loop that is less than 60 µm and high bonding speed that is less than 60 ms per wire, the wire bonder will require a highly accurate vision alignment system, very fast and precise motion control for the X–Y table and bondhead, highly accurate and consistent bond force and impact force control for the bond-tip, and a robust and reliable ultrasonic transducer, etc. After wire bonding, the leadframe with wire bonded chips is sent to another buffer station ready for the moulding process. The multiple leadframes with wire bonded chips are loaded into the moulding machine in batches. Usually, multiple mould presses, each with multiple cavities, are used in a
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moulding machine to increase its total throughput. The moulded leadframes are stacked up in a carrier and sent to the post-mould curing oven in batches. After the post-moulding process, the moulded leadframe is sent to the ‘trim and form’ machine for the separation and trimming process. Another buffer station next to the post-mould curing oven is frequently used to balance the throughput between the oven and the separation machine. The ‘trim and form’ processes are performed by strong mechanical presses. The devices on the leadframe are separated from the rest of the leadframe and shaped to the right geometry by a set of high-speed mechanical presses. If a strip tester is used at the end of the line, the leadframe will be partially cut before sending to the tester. Complete separation and trimming of the units will be done after testing. Serviceable units are separated out and bad units are scrapped. Depending on the choice of output carriers for endproducts, the devices can be unloaded and output to JEDEC trays or plastic tubes. The quality of the assembly materials is very often related to their cost. One needs to balance the cost of materials against their qualities in order to obtain a good yield at a reasonably low cost for the end product. This requires the assembly equipment to be able to identify and skip the known defective units during the assembly process. For an in-line process production line, multiple equipment for various assembly processes is interconnected to each other. A common scheme across all equipment to identify known bad units will be needed. The layout of assembly equipment clearly defines the material loading zone(s), the path of material flow, the main assembly areas and the unloading port of the finished product. Mechanically, the docking heights of the equipment have to be aligned so that the equipment can be linked to form an in-line process line. Electrically, a common communication method is selected for all equipment in the line. Communication between equipment can be achieved by simple multiple I/O signals or the SECS/GEM industrial standard implemented by RS232 or TCP/IP communication.
7.3
Material handling in assembly equipment
Other than the process requirements, the materials-handling requirements for assembly processes are also a major consideration in the design of the assembly line. The handling sequences of the assembly processes determine how the materials flow in the assembly equipment as well as in the production line.
7.3.1
Layout of processing modules
The layout of the functional modules in assembly equipment defines the footprint of the equipment as well as the flow of materials and its handling
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method. The machine layout defines which assembly process will be handled in parallel and which process will be handed in sequential order. For pickand-place equipment, it dictates the speed and the accuracy of the placement process by the equipment. High-speed pick-and-place equipment with good repeatability and accuracy can be achieved by careful machine design and optimization.
7.3.2
Orientation of processing parts
The orientation of the assembly parts has to be well defined in the assembly line. If it is not it will cause confusion and falsely identify known good units in the assembly equipment. Special attention is needed for multiple units on the same carrier of the assembly part. It is particularly important if both good and bad parts go through the assembly line at the same time. They can be identified only by recognizing the numbers or marks on their carriers together with the known good unit information (say, a wafer test map) from the cell controller. In the in-line process line, known good unit information will pass to equipment downstream of the production line through their communication link. The location of a known good unit in the test map is usually given by referring to a reference location or mark on the carrier at a specific orientation. If the orientation of the carrier of the assembly part does not match with that given in the test map, fault identification of known good units will result. Some assembly materials can be handled without a carrier but some cannot. If the carrier is used to transport the assembly materials along the workholder, there should be a reference mark on the carrier so that it is identifiable by a sensor (either optical or mechanical) on the equipment so that its orientation can be determined. In addition, one should add non-symmetrical features to the pockets on the carrier to eliminate the miss-orientation of the parts.
7.3.3
Vision alignment technique
The vision alignment technique is a very powerful technique to align assembly parts to a very high accuracy. A magnified video image of the assembly part on the workholder is grabbed by a CCD or CMOS camera using a video frame grabber under illuminating optics. The digitized image is then analyzed by some algorithm of a pattern recognition program in a PC. The displacement of the part from the cross-hair of the camera can be determined absolutely by image-processing software and optical calibration of the digitized image. In general, the higher the optical magnification and higher pixel density of the camera, the better the resolution obtained by this technique. The choice of optical components and configuration of the light source will
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affect the image quality and the resolution of the detectable features. With the proper choice of optical components and the sub-pixel image-processing technique, the resolution for vision alignment down to 1 µm is achievable in some precise equipment. The position error or offset of the assembly part from its target location can be determined by the vision alignment system. This offset (position error) can be compensated by the last motion step of the parts holder. The ultimate placement accuracy of the handler is determined by the resolution calculated from the algorithm for vision alignment determined from the algorithm of the image processing software, the accuracy of the calibration factor for the vision-to-motion conversion scheme, and the image quality of the reference marks on the assembly part and fiducials on the target substrate.
7.3.4
Kinematic design and motion control
In some cases, visual alignment of parts may not be possible due to space or time considerations. In these cases, high placement accuracy of the parts can be achieved only by kinematic design and careful error budgeting of all the related motorised axes of the workholder and the pick-and-place arm. For example, a stiff indexing arm driven by a linear motor with a high-resolution linear optical encoder located in close proximity to the moving target can greatly reduce placement error. The resolution of an optical linear encoder varies from 5 nm to 10 µm. Encoders with a resolution of 0.1 or 0.2 µm are readily available at reasonable prices. Nevertheless, the selection of encoder is usually based on the requirement of the motion performance as well as the type of bearing being used for that motion axis. A high-resolution optical encoder with resolution down to 50 nm or below may be used only for a motion axis controlled by a high-speed controller and constructed with high-quality linear bearings such as linear air bearings or precise fracture bearings.
7.3.5
Interchangeability of application tools
Another consideration in the design of a workholder is its application range. When a product changes, the application tool may need to be replaced. The interchangeability of tool sets in respect to the common mounting interface should be considered in the early stages of design since it will affect the layout and the definition of the reference surfaces for the module. It is desirable to have a common mounting interface for various sets of tools to handle similar but not identical products. In addition, tool-less design for product conversion is very often a requirement for the equipment used in cleanroom environments.
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Pick and place processes
Most of the assembly processes involve more than one assembly part. Two or more components can be joined together. Pick and place processes are very common for assembly parts. Pickup and placement processes will be discussed separately since their requirements are very different. A thin-die pickup process will be used as an example to demonstrate optimization of this pickup process. The thermoplastic tape placement process is used as an example to elaborate the considerations of this specific placement process. The implementation of the pickup and placement processes for assembly equipment will be affected by: • • •
the sizes and dimensions of the process materials from previous assembly processes the physical states (gel or solid) of the process materials from previous assembly processes the process requirement of the next assemby process.
7.3.7
Example of thin-die pickup process
Well controlled pickup motion sequences together with a special tool are needed for detaching very thin (down to 50 µm) silicon chips from dicing tape in die-attach processes. The processing window for the die-detachment process has to be well understood before making an appropriate tool for this pickup process. Finite element analysis (FEA) of the stress distribution on the die in the die-detachment process shows that the arrangement and the locations of the ejecting pins underneath the die are critical for a successful pickup process [3]. A 2-D finite element model is constructed to calculate the stress and the strain induced on a 50-µm thin die by the bending moment and the interfacial peeling strength between the die and the dicing tape when the upward displacement load of the ejecting pins is introduced. The optimized locations of the ejecting pins can be obtained from this FEA model if the thickness of a die and the adhesion strength of a dicing tape are known. The result shows that if the outer ejecting pins are located at 0.5 mm from the adjacent edge of a 50 µm die, the interfacial peeling stress induced by the upward load of the ejecting pins will be strong enough to overcome the adhesion strength of the die on UV dicing tape before the die reaches its critical strain due to the bending moment induced by the upward load of the ejecting pins. Hence, the die can be detached from the dicing tape without cracks.
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Example of controlled placement processthermoplastic adhesive tape application process
Subsequent assembly processes will be likely to proceed right after the placement of the part, therefore a well-controlled placement process is needed to prevent damage to the assembly parts and to ensure that parts will be placed at the target location accurately. Bonding processes such as die attach, flip chip bonding, thermoplastic tape bonding, etc., are susceptible to void formation or interfacial delamination failures. For example, in thermoplastic adhesive tape application processes, a very soft thermoplastic tape will be picked up and placed on a heated flat surface of the substrate. If a tool having a flat contact surface is used to pick and place the tape, there will be voids between the thermoplastic tape and substrate surface. To eliminate these voids a special tape application tool with a compliant convex surface driven by a two-step motion profile will be needed [4]. The schematic of this compliant convex collet and the squeezing action of the special collet are shown in Fig. 7.2. The voids trapped inside the contact interface can be squeezed out by compressing this special collet. Figure 7.3(a) shows the voids trapped inside the bonding interface between the thermoplastic tape and the flat substrate surface if a flat application tool is used. The voids are eliminated by the collet having a compliant convex surface (Fig. 7.3(b)). The convex collet holding the tape will move down to a level where the central convex surface is just touching the surface of the substrate. The vacuum suction holding the tape will then be turned off. The collet goes down, further compressing the tape and squeezing the voids trapped on the interface out sideways until the collet is flattened totally. Note that the surfaces of the substrate and the adhesive tape need to be cleaned thoroughly otherwise voids trapping foreign materials will be found.
Holder Convex rubber collet Small vacuum hole Thermoplastic tape Squeezed direction Die or substrate
7.2 The diagram illustrates the convex contact surface of the special collet which can be used to squeeze the voids trapped inside the bonding interface. The centre region of convex surface will touch down first. Compressive action of the collet will squeeze the voids out sideway until the compliant collet is fattened.
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Trapped foreign material Trapped air bubbles
Flat substrate
Placed tape
(a)
(b)
7.3 (a) Voids form at the interface between the thermoplastic tape and flat substrate surface if a flat application tool is used, (b) a void free interface can be obtained if a tape application tool with compliant convex surface is used (courtesy of ASM Pacific Technology).
7.4
Control of process parameters for assembly processes
The specifications of assembly equipment are driven by the process requirements it is going to perform. It should provide the largest process window for a given assembly process. Most of the assembly processes can be realized at a specific pressure, temperature and time under certain controlled environments and some may require alternative energy sources such as ultrasonic vibration or laser radiation. The exact suitability of the equipment providing the control of these process parameters is essential to yield and productivity. For most cases, two or more processing parameters need to meet their desirable values at the same time in order to perform a specific assembly process. Some assembly processes involve physical interaction of the processing materials on the substrates. The process window for all parameters fulfilling their requirements in the same process condition may be quite narrow. Therefore, the assembly equipment needs to have the capability of controlling all related parameters precisely for realizing a given assembly process. Some common processing parameters for assembly processes are discussed below.
7.4.1
Accurate timing
Accurate timing control is a prerequisite in all assembly processes. The timescale for most of the assembly processes ranges from sub-milliseconds to tens of seconds. For example, the temporal profile of a laser pulse used for laser spot welding is important and needs to be fine-tuned in accordance
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with the optical properties of the surface coatings and the thermal properties of the base materials of the welded parts. The resolution of the timing control for the temporal profile of the laser pulse should be in the order of 0.1 ms for a pulse duration of about ten milliseconds. Another example is the gold wire bonding process. For an advanced gold wire bonder, the cycle time for bonding a wire is less than 70 ms. Firing time of the ultrasonic transducer is as short as 3–5 ms depending on the excitation frequency of the ultrasonic transducer being used for the gold wire bonding process. The throughput of an ultra-fast pick-and-place machine can achieve 20,000 units per hour (UPH) which means that the cycle time for a complete pick-and-place process is less than 180 ms.
7.4.2
Temperature profiling
A number of assembly processes are realized at elevated temperatures, for example, the thermosonic wire bonding process, eutectic soldering process, die attachment by thermoplastic tape adhesives, solder reflow process, etc. The work chuck of the equipment needs to maintain a uniform temperature distribution over all bond sites for the bonding process being performed at various locations on the substrate. The temperature distribution over the whole working surface of the work chuck should be kept to within +/– 3 °C to +/ – 5 °C. The eutectic soldering process occurs at a very narrow and specific temperature window therefore the stability and the fluctuation of the temperature can affect the bonding quality of the assembly. Of course, the melting temperature of eutectic solder may shift away from its nominal value if the composition of the metallization is changed on the bonding interfaces for some reason. Hence, the quality of the bonding materials is also important for the eutectic soldering process. In the case of a solder reflow oven, the controllability of the rate for ramp up and cool down of the temperature at various heating zones for the reflow profile is one of the major concerns since solder joint reliability will be affected by the solder reflow profile [5].
7.4.3
Accurate placement
Precision positioning is a basic requirement for a part handler used in an assembly process. The placement accuracy is mostly defined by the requirements of the specific assembly process. For example, ordinary die attach equipment needs to place a die to within 20 µm of its target location. However, if it is a flip chip process, the placement accuracy has to be improved and the placement error should be within 5–10 µm. The bonding accuracy of a wire bonder for bonding 45 µm bond pad pitch (BPP) is less than 3 µm
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error! In an automatic machine, a visual alignment system with a very high spatial resolution should be used to detect the target location and the present position of the part. Proper material selection and careful kinematic design are needed to compensate thermo-mechanical deformation of the pick arm in order to achieve high placement accuracy.
7.4.4
Force and pressure
After moving the part to the assembly position, appropriate clamping force will be applied to the part to provide enough interfacial pressure along the interface for assembly processes. The required compressive forces are quite critical for the assembly processes such as thermo-compression bonding, thermosonic wire bonding, injection moulding, etc. The stiffness of the fixture should be high enough to withstand the applied force without causing significant structural deformation. The accuracy of applied force should be in the range of 0.2% to 1% of its full range depending on the needs of the pressuresensitive processes. An external load cell is used for force calibration, which measures the actual force delivered by the system. The linearity and accuracy of the force actuator can be measured and calibrated by using this external load cell. Precise motion and vibration control are needed to eliminate unwanted impact force for very delicate bonding processes such as thermosonic wire bonding and flip chip bonding processes. A fine pitch thermosonic wire bonding process requires a bond force on the order of ~ 10 gf and its repeatability should be better than 0.5 gf. On the other hand, thermo-compression bonding process of thermoplastic adhesives may require over 60 kgf per stroke if multiple units are pressed at the same time. The interfacial pressure required for this bonding process is in the range 0.1 to 1.0 MPa. The planarization of a sizable press should be adjusted to within 1:1000 to 1:10000 with respect to the working surface to ensure a uniform pressure distribution.
7.4.5
Application of ultrasonic energy
Ultrasonic energy is one of the major energy sources for assembly processes, and has a profound effect on softening the process materials. Ultrasonic assisted processes such as ultrasonic metal welding, ultrasonic plastics welding, ultrasonic wire bonding, ultrasonic paper cutting, ultrasonic cleaning, etc., are commonly adopted in automatic assembly equipment. According to the process requirements, the designer needs to select the ultrasonic excitation frequency and the ultrasonic transducer, design the geometry of the mechanical amplifier, and design the mounting interface and the process tool for a given assembly process.
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The ultrasonic excitation frequency can range from 20 kHz to 2 MHz and the power range from less than a watt to a few kilowatts. A small fraction of a watt of ultrasonic power is enough to make a wire bond and it takes up to a few kilowatts of ultrasonic power to weld two metal parts together. The ultrasonic transducer is a device which converts the electrical energy to mechanical energy in the ultrasonic frequency domain. A mechanical amplifier is basically a mechanical resonator that resonates at the operating frequency and magnifies the amplitude of the vibration of the transducer at anti-node location. The selection of material used for the mechanical amplifier is important since the acoustic velocity and acoustic damping factor of the materials can change the resonating behaviour of the amplifier. Modal analysis is a basic CAE tool for the design optimization of an ultrasonic mechanical amplifier. The transducer is integrated to the mechanical amplifier by a very strong mechanical coupling. With a proper design of the mounting interface, the process tool can be mounted onto the mechanical amplifier with minimum power attenuation. The absorption of the ultrasonic energy by the process part greatly depends on the materials properties of parts being made and the interfacial pressure of the bonding interface.
7.4.6
Application of laser radiation
High-intensity laser radiation can be used for laser marking, laser welding, laser drilling, laser cutting, laser cleaning, etc. Both continuous-wave (cw) and pulse mode lasers can be used for various assembly processes. Pulse mode laser can deliver high peak radiation power (up to megawatts) and laser pulses that facilitate the ablation process during the laser marking, cutting, drilling and cleaning processes. High energy continuous-wave lasers can be used for laser welding processes to join metallic or plastic parts. A pulse mode laser having a pulse width at a few milliseconds to tens of milliseconds can be used for laser spot welding. However, there is no single laser that can work for all kinds of materials since their radiation linewidth is usually quite narrow and the absorption of the laser radiation depends very much on the optical absorption coefficient of surface materials and the surface metrology of the process part. Laser radiation can be delivered by bulk optics or optical fibres. It would be easier to position and install the laser head in assembly equipment if optical fibres are used to deliver the laser radiation. However, the laser power delivered to the part may be limited by the power breakdown of the optical fibre. If an invisible high-power laser source is used, a visible guiding laser light through the same optical fibre will be useful for machine setup. However, one should always remember that the focal point of the visible guiding laser is not at the same depth as the focal point of the invisible laser source.
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On the other hand, one needs to pay special attention to the selection of a power/energy meter for calibrating the radiation power of a high-power laser source. Depending on the type of detector (e.g., silicon detector, GaAs detector or pyroelectric sensor, etc.), the optical power/energy meters are good only for certain spectral ranges and power/energy density ranges. If the power meter is a pyroelectric sensor of broad spectral range, its temporal response is usually slow.
7.4.7
Application of reducing gases
In some high-temperature bonding processes, for example copper wire bonding, eutectic die attach process, etc., reducing gases are used to prevent the oxidation of the copper leadframe and bonding materials, for example copper wire and solders. In addition, the reducing atmosphere in the process environment can affect the wettability of the eutectic solder on the leadframe. One can change the kind of reducing gas and/or its composition to increase the surface energy of the bonding surface so as to improve the wettability of the solder on the leadframe. The inlet and outlet of the reducing gases should be arranged in a way that the gases are flowing uniformly across the bonding surface of the leadframe. In order to keep the temperature uniform inside the heating chamber, one should control the gas flow and prevent excessive leakage of the reducing gases inside the chamber.
7.5
Case study: design of customized assembly equipment – an alignment laser welder for a fibre pigtailed TO-can laser diode package
The considerations for designing assembly equipment are mostly governed by the requirements of the assembly processes. An active alignment laser welder for fibre pigtailed TO-can laser diode package will be used as an example to illustrate how to design customized assembly equipment based on the special requirements for assembly processes. This equipment, as described below, is used in batch assembly processes. It can be easily converted to the assembly equipment used in the in-line assembly processes if bridges are made to link it to upstream and downstream equipment. However, the major concern is whether its throughput can match the rest of the production line since active alignment and laser welding are slow assembly processes.
7.5.1
Assembly processes analysis
An active alignment laser welder basically performs two major assembly processes: (i) active alignment process to obtain the maximum optical coupling of the laser light from the laser diode of the package to the single-mode fibre
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(SMF), and (ii) laser spot weld of the metallic housings to join them so that the positions of the parts given by the active alignment are permanently fixed. One needs to understand the process requirements for assembling the target package in order to define the design specifications for the active alignment laser welder. Figure 7.4 shows the sectional view of the target optoelectronic package which is a fibre pigtailed TO-can laser diode. It is a relative low-cost design. A commercial 5 mW and un-cooled TO-can laser diode is used as the laser source [6]. In this device, the laser light emitted from the laser diode will be focused by a ball lens and coupled to a ferruled SMF inside the metallic housing. The bottom assembly containing the laser diode needs to be carefully aligned to the top assembly where a ferruled SMF is furnished so that the maximum laser light can be obtained from the other end of the fibre. All the metallic housings of the package are made of stainless steel AISI 304L. These metallic housings will be joined together by laser spot weld so that the maximum optical coupling resulted from the active alignment of SMF can be secured. The configuration of the package illustrated in Fig. 7.4 shows that a TOcan laser diode can be inserted and welded to a bottom metallic case prior to the final assembly process as described below. A ring with flange facing
X–Y metallic case
Sleeve
Single-mode fibre Ball lens
Z metallic case Laser diode
Lens cap
Bottom metallic case
Photodiode
TO-can package
7.4 The sectional view of a fibre pigtailed TO-can laser diode package.
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upwards, called a Z metallic case, is inserted into this bottom metallic case from the top. Before they are welded together, the Z metallic case can slide up and down along the cylindrical axis of the bottom case to allow axial active alignment. The SMF housed inside a precise ferrule will be inserted into an X-Y metallic case before they are placed onto the Z metallic case. The active alignment along the X-Y direction can be achieved by adjusting the X-Y metallic case laterally on the top flange of the Z metallic case. The objective of the active alignment process is to direct the laser beam from the laser diode into the SMF and to achieve maximum optical coupling efficiency. From the results of numerical simulation, the maximum optical coupling efficiency is expected to be 16–18% for this package [7]. The maximum optical coupling efficiency can be achieved when the tip of the SMF is aligned to a position such that the propagation mode of a laser beam in SMF at its input face is the best matched with the spatial propagation mode of a Gaussian laser beam at its focused beam waist [7]. The requirements for active alignment The required active alignment resolution can be derived from the plots of the optical coupling efficiency versus the offset of the lateral alignment. The result obtained from simulation is helpful for one to understand the expected coupling efficiency during the active alignment procedure. Figure 7.5 shows simulated curves of the optical coupling efficiency versues the alignment offset along the lateral and axial directions for a given optical construction.
0.15
Coupling ef
ficiency
0.20
2.2
0.10 0.05
2.0 ce
0.00 –20 La –10 te ra lo f
)
m
(m
an
1.8 0 fse
10
t(
icr
20
on
)
Ba
to s-
n
le
ll-
m
e
br
-fi
st di
1.6
7.5 The simulated optical coupling efficiency of the given optical design versus the lateral offset as well as the ball lens to fibre distance along the axial direction. The line joining the central peaks or valleys of the simulated curves shows the changes of coupling efficiency along the axial direction.
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The line joining the central peaks and valleys of the curves of optical coupling efficiency versus lateral offset in Fig. 7.5 is basically the coupling efficiency of the SMF along axial direction on the optical axis of the laser diode and the ball lens. It is found that the lateral alignment tolerance for maintaining 95% of the coupling efficiency is less than ~+/– 1 µm [7]. It implies that the step size for the lateral X/Y alignment has to be in tens to hundred nanometres (~ < 0.1 µm) resolution. On the other hand, the alignment tolerance along the axial direction for coupling efficiency higher than 95% can be within ~+/– 30 µm from its peak position. Therefore, the minimum step for axial alignment can be a fraction of a micron (~< 1 µm). It is desirable to have resolution at 50 nm and 0.2 µm for the encoders on the alignment stages along lateral X/Y axes and axial Z axis, respectively, for the motion platform conducting the active alignment of a given device. With the consideration of the placement consistency of the laser diode in the ball lens capped TO-can and the TO-can onto the bottom metallic case as well as the insertion repeatability of this assembly onto the work chuck, the estimated search range for lateral alignment along the X/Y-direction can be as large as ~ +/– 200 µm and the search range for axial alignment along Zdirection can be ~ +/– 150 µm. Even if the axial alignment is already at its peak, it will be very time consuming to do the active alignment along the lateral directions for the search window at 400 µm × 400 µm with the incremental step size at only 50 nm. The chances of the fibre tip locating at +/– 1 µm from the peak value inside a search window of 400 µm × 400 µm is as small as 25 ppm! A smart searching algorithm and procedure are needed to reduce the time required for peak searching. Therefore, a coarse alignment procedure is introduced prior to the fine active alignment procedure to reduce the time needed for the alignment process. The goals for the coarse alignment are (i) locating the focused laser spot of the laser beam and (ii) positioning the laser diode such that this focused laser spot is within +/– 20 µm laterally and +/– 50 µm axially from the tip of the SMF. The subsequent fine active alignment can be done in this much smaller search range. These two-step alignment procedures can reduce the total alignment time from more than two minutes to less than 30 seconds [8]. An innovative method to locate the position of the focused laser spot of the laser beam is done by introducing a multi-mode optical fibre (MMF) of core size at ~ 50 µm as optical probe to detect the laser light at a coarse alignment stage [9]. The numerical aperture (NA) of a MMF is 0.2–0.3 which is much bigger than that of the SMF at ~ 0.1 since the core size for a typical SMF is ~ 10 µm. Due to large NA, the light reception angle will be larger for the MMF. Therefore, a bigger incremental step can be used for searching the laser spot.
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The intensities of the laser beam coupled into the MMF at a different level from the focused laser spot are shown in Fig. 7.6. The curves show distinctive peak features when the MMF is scanning at 0.15 mm above and below the focusing plane. The coarse alignment is able to resolve the peak to within +/– 50 µm along the axial direction. Since the optical resolving power of the MMF is lower than the SMF, the width of laser spot as detected by scanning the MMF laterally across the beam will appear bigger. The width of laser spot as detected by the MMF at the focusing plane is ~ +/– 30 µm as shown in Fig. 7.6. Therefore, the coarse alignment along lateral direction should be able to locate the peak to within +/– 20 µm. After completing the coarse alignment procedure, the assembly containing the laser diode will be moved to perform the fine active alignment by using the ferruled SMF. The coarse alignment procedure positions the laser spot to within +/– 20 µm laterally from the optical axis of the SMF and +/– 50 µm 2000
MMF optical coupling intensity (arb. units)
1800
At focus Focus plus ~ 0.15 mm Focus less ~ 0.15 mm
1600 1400 1200 1000 800 600 400 200
–100
–80
–60
–40
0 –20 0 20 Lateral axis (microns)
40
60
80
100
7.6 The optical intensity of the laser beam coupled onto the multimode fibre MMF when the laser diode scans laterally across the MMF. The line connecting the solid cycles is the data obtained from the scanning on a focusing plane that contains the focused beam waist of the laser beam. The line connecting the empty squares is the data obtained from the scanning 0.15 mm above the focused plane. The line connecting the empty diamonds is the data obtained from the scanning 0.15 mm below the focused plane.
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axially from the tip of the SMF. The fine active alignment procedure will be performed on the stage containing X-Y and Z motion axes. The search window for the fine alignment is now reduced to +/– 20 µm laterally and +/– 50 µm axially. A few iterations for the active alignment procedure will be needed to locate the peak coupling intensity of the laser beam. Each iteration includes the following steps: (i) search the peak in the laterally direction, (ii) move the laser diode by the X-Y stage to the new position where the maximum optical coupling is obtained, (iii) move the laser diode on the Z stage up and down and search for the peak coupling intensity along the axial direction, and (iv) position the laser diode at the new height where the maximum optical coupling is obtained. Repeat the procedure until the maximum optical coupling intensity obtained from current iteration is different from the one before by less than 3%. The requirements of laser spot weld Laser spot welding is a very common technique adopted to join components in optoelectronic packages [10]. The isometric view of the welding part and the direction of the laser beam are shown in Fig. 7.7(a). Three laser processing heads are mounted on the same plane and arranged 120 degrees apart from each other as shown in Fig. 7.7(b). The laser processing heads point inwards aiming at the target of the same centre and make a 45 degree angle from the vertical axis. Three laser spots project on different locations (of 120 degrees apart) at target interface along the perimeter of the metallic case for each laser shot. The laser spots are elliptical at the target surface since the laser beams project on the surface at 45 degrees from normal. The sectional view of the welding part and the locations of target interfaces and the firing sequence of the laser welds are indicated in Fig. 7.7(c). The alignment of the laser beams is such that all three beams focus at the surface of the welding part and the beams point at the same centre of a circle. The laser welds will be done at the two junctions, firstly between the bottom metallic case and the Z metallic case, and secondly between the X-Y metallic case and Z metallic case. The firing sequence of the laser welds is important since there will be post-weld shift (PWS) after the laser spot weld [11–13]. The PWS will change the relative positions of the welded components and hence it will upset the maximum coupling efficiency resulting from the active alignment procedure. The first laser shot will be done at the junction between the bottom metallic case and the Z metallic case since the coupling efficiency is less sensitive to PWS along axial direction for this optical construction. The 2nd and 3rd laser shots will be done at the junction between the X-Y metallic case and the Z metallic case.
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120 degrees
Laser beam
(a)
(b)
45 degrees X–Y metallic case
Junction for 2nd and 3rd laser shots Z metallic case Bottom metallic case
Junction for 1st laser shot
(c)
7.7 The isometric view of the welding part and the direction of the three laser beams are indicated in (a). The plane view of welding parts and the arrangement of the laser process head at 120 degrees apart are shown in (b) and the sectional view of welding parts and the pointing directions of the laser beams making a 45 degree angle from the vertical axis of the package are shown in (c). The firing sequence of the laser welds and their target welding interfaces are also indicated in the diagram.
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The assembly yield of this package depends very much on the PWS of the welding at this junction. The 3rd laser shot serves as a security weld which is used to strengthen the joint along this interface. This security weld will be done at a location about 30 degrees from the 2nd weld and it can be achieved by turning the package 30 degrees about its central vertical axis. Three equal energy laser pulses 120 degrees apart of typical width 10 ms in which 0.5 ms are for ramp up, 9 ms at plateau and 0.5 ms for ramp down. The laser beams are directed and focused at the junction between the bottom metallic case and X-Y metallic case. The laser welding experiment shows that the optimized power density for each laser pulse is ~ 0.16–0.19 MW/ cm2 for laser spot welding of the Z metallic case and bottom metallic case. At 0.17 MW/cm2 the tensile breaking load for the joint is up to 32 kgf. Figure 7.8 shows the energy partition, the location and the size of the laser spot at the target surface of the welding junction. The peak power of each laser pulse is around 0.73 kW and the size of the elliptical laser welding spot is 0.6 mm by 0.84 mm. The location of the laser spot for the lap-fillet weld joint is 60% of the spot projected onto the bottom edge of the Z metallic case and 40% projected onto the side wall of the bottom metallic case. For the second laser shot, three laser spots are carefully aligned to irradiate at the junction between the base flange of the X-Y metallic case and the top surface of the Z metallic case. The power density of each laser pulse at the target surface is in the range of 0.2–0.3 MW/cm2. A typical pulse width is ~ 8 ms in which 0.5 ms are for ramp up, 7 ms at plateau and 0.5 ms for ramp down. The peak power for each pulse is ~ 0.4 kW. The joint between the XY metallic case and the Z metallic case is a fillet weld joint. The location,
45 degrees
Laser Z metallic case 60%
0.50 mm
Interface
Interface 40%
0.34 mm Bottom metallic case
0.60 mm
7.8 The recommended energy partition, the location and size of the laser spot at the target surface for welding the Z metallic case and the bottom metallic case.
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energy partition and size of the laser spots at the target surface for welding the X-Y metallic case and the Z metallic case are shown in Fig. 7.9. The laser welding spot for the fillet weld joint is elliptical and 0.42 mm by 0.6 mm. The location of the laser spot for the fillet weld joint is 45% of the spot projected onto the base flange of the X-Y metallic case and 55% projected onto the top surface the Z metallic case. Both the welding depth and shear breaking load of the joint increase with the power density of the laser pulse. However, the laser power density at the welding junction has to be optimized. A weak welding joint will result if the laser power is too low. Voids inside the weld will be formed during rapid solidification if the laser power is excessively high, say, > 0.4 MW/cm2. In addition, in order to reduce the PWS induced by the laser weld, it is more desirable to have a less powerful laser pulse for welding this joint. To enhance the strength of the structure, an additional laser security weld at a higher laser power can be applied. This security weld can be done at a location rotated about 30 degrees away from previous weld. A shear breaking load up to 28 kgf can be achieved if the junction is welded together by the laser pulses of power density at 0.2 MW/cm2 followed by a security weld at 30 degrees apart having laser pulses at 0.3 MW/cm2 power density. Post-weld shift (PWS) Precision active alignment is needed for coupling the laser beam from a laser diode into a SMF. Minute changes in physical dimensions will decrease the optical coupling efficiency. Due to the PWS induced by laser welding, an active realignment along lateral directions may be needed after welding between the Z metallic case and the bottom metallic case. This joint is
45 degrees Laser X–Y metallic case 45%
0.27 mm Interface
55%
0.33 mm Z metallic case
0.42 mm
7.9 The recommended energy partition, the location and size of the laser spot at the target surface for welding the X–Y metallic case and the Z metallic case.
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welded first since the drop in optical coupling efficiency due to the axial misalignment induced by PWS is less sensitive to the lateral misalignment. In order to reduce the magnitude of the PWS induced by welding the XY metallic case and the Z metallic case, a preload is applied to compress the X-Y metallic case onto the Z metallic case during laser welding. A high compressive pressure building up at the contact interface minimizes the PWS displacement of the parts during the solidification of the weld right after laser spot welding. Figure 7.10 shows the measured PWS versus the compressive preload applying to the X-Y metallic case on top of the Z metallic case [11]. It is found that a minimum 2 kgf (or equivalent to 1.2 MPa at the interface) compressive preload is needed to prevent excessive PWS. An alternative practice to minimize the negative impact of PWS is the introduction of a so called, ‘laser hammering’ procedure [14]. However, this procedure may be quite time consuming. An excessive PWS may also result if the laser welding spots deviate from their nominal welding locations. Figure 7.11 shows the measured PWS versus the position difference ∆ of the laser spots. The inset in Fig. 7.11 explains that the position difference ∆ is the distance between the extreme positions of the laser spots at the target interface. The line connecting the solid circles represents the measured PWS versus position difference ∆ of the spots at ~0.8 kgf (or 0.48 MPa) compressive preload. The line connecting the open circles represents data obtained from the condition having a 3 kgf (or 1.8 MPa) compressive preload. As shown in Fig. 7.11, the PWS increases with the position difference ∆ among the laser spots. 4
Resultant PWS (µm)
3
2
1
0 0
2
4 Load (kg)
6
8
7.10 The measured post-weld-shift (PWS) is plotted as a function of the compressive preload applying to the X–Y metallic case on top of the Z metallic case during the laser weld.
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3 ∆ Resultant PWS (µm)
2.5 2 1.5 1
0.5 0 0
40 80 120 Spot position difference ∆ (µm)
160
7.11 The measured post-weld-shift (PWS) versus the position difference ∆ of the laser spots at ~0.8 kgf (solid circle) and 3 kgf (open circle) compressive preload on the X–Y case during laser weld. The position difference ∆ of the laser spots is the distance taken between two extreme spots as shown in the inset of the diagram.
It is desirable to ensure the laser processing heads are well aligned and the position difference ∆ of the laser spots does not exceed 80 µm. A PWS of less than a micron can be achieved if (i) a compressive preload of > 2 kgf is applied and the interfacial pressure is larger than 1.2 MPa between the welding parts and (ii) the locations of laser welding spots deviate from their nominal positions by less than +/– 40 µm.
7.5.2
Prototype active alignment laser welder
In a prototype active alignment laser welder, the good TO-can laser diode assembly inserted to the pockets of the metallic carriers is fed from the magazine on one side of the equipment. Each assembly contains a TO-can laser diode having a bottom metallic case pre-welded to it with a Z metallic case loosely attached and placed on top of the bottom metallic case. The magazine on the other side contains empty X-Y metallic cases (i.e. the ferruled SMF are not present) on the metallic carriers. This empty X-Y metallic case is picked and transported to a holder in the middle of the machine where fine active alignment is carried out. A ferruled SMF is inserted into X-Y metallic case at this location. A holder clamps and secure the X-Y metallic case containing the ferruled SMF in position preparing for the fine alignment. The laser diode assemblies in the carrier are fed into the machine on the input track by a motorized indexer. A vacuum pick arm will pick a TO-can
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laser diode assembly from the carrier. The whole assembly, that is the TOcan laser diode housed inside the bottom metallic case and the Z metallic case loosely attached on it, will be inserted into the socket on the work chuck. The leads for the electrical connection of the laser diode will contact the electrodes of the socket. A constant electric current will be provided to power up the laser diode. Figure 7.12 is a schematic diagram of the motion platform for active alignment. The work chuck will be mounted on this platform which consists of a theta stage and a Z stage mounted on top of the X-Y stage. When the laser diode assembly is inserted in the work chuck on the platform, the Z metallic case will be clamped by a holder which is mounted directly on the X-Y stage. The height of the Z metallic case measured from the X-Y stage is fixed by this holder. The base of the bottom metallic case as well as the TO-can laser diode will be clamped and locked to the work chuck on top of the theta and Z stage. The Z stage can move the work chuck up and down providing position adjustment along the axial direction of the laser diode. The theta stage can turn the assembly to different angular orientations. The work chuck containing the TO-can laser diode assembly can be moved up and down along the Z axis by the Z stage and laterally along the X-Y axes by the X-Y stage. The X-Y metallic case is held by a holder on top of this motion stage at the active alignment location. The holder can move up and down to bring the ferruled
X–Y metallic case containing the ferruled SMF fixed by a holder
Z metallic case
Welding spot Holder
Bottom metallic case containing the laser diode
Work chuck and theta stage
Z stage
X–Y stage
7.12 The schematic diagram shows the sectional view of the active alignment motion platform of the work chuck. The theta stage and Z stage are built on top of the X–Y stage.
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SMF/ X-Y metallic case assembly onto the top surface of the flange of the Z metallic case for fine active alignment. Coarse alignment along the lateral direction can be achieved by moving the work chuck to the location where a fixed MMF is mounted. At this location, the MMF will be used to detect the light emitted from the laser diode. The work chuck containing the laser diode assembly on the X-Y stage will scan across the MMF for lateral coarse alignment. The work chuck will arrive at a position where the maximum light intensity will be received by the MMF. Coarse lateral alignment can bring the beam waist to within +/– 20 µm from the optical centre of the MMF. Since the Z metallic case is clamped and fixed by a holder mounted on the X-Y stage, the bottom metallic case can be slid freely inside the Z metallic case before laser welding. Coarse alignment along the axial direction can be achieved by moving the laser diode assembly up and down on the work chuck by the Z stage. The distance between the laser diode and the input face of the MMF can be adjusted accordingly. The optimal height of the bottom metallic case can be obtained at the peak optical coupling intensity for the MMF. As a result, coarse alignment along the axial direction positions the focused beam waist to within +/– 50 µm from the input face of the MMF. After completing the coarse alignment procedure, the laser diode assembly on the work chuck will be moved to the fine alignment location by the X-Y stage. The result from the coarse alignment will be used to determine the displacement of the X-Y stage and the height of the Z stage of the work chuck when the laser diode assembly has transported to the fine alignment location. At this location, the X-Y metallic case containing the SMF will then be moved down and land on top of the assembly containing the Z metallic case and the TO-can laser diode. As obtained from the coarse alignment procedures, the position of the focused beam waist before fine active alignment is within +/– 20 µm laterally from the optical axis of the SMF and within +/– 50 µm axially from the tip of the SMF. The work chuck is mounted on an X-Y stage which is built on a table lifted by an air bearing and having motion resolution at 50 nm. The active alignment procedure will be performed by cooperative X, Y, Z motions of the X-Y stage and Z stage. The X-Y metallic case is held fixed on top of the Z metallic case. The Z metallic case is fixed in height but can be moved laterally by the X-Y stage. The height of the laser diode inside the Z metallic case can be changed by the Z stage below the work chuck. An iterative procedure is needed for the active alignment to maximize the coupling efficiency of the laser output into the SMF. The alignment procedure can be considered as successful if the resulting optical coupling efficiency changes by no more than 3% as compared to the previous alignment result. A 25 W air-cooled pulsed mode Nd:YAG laser welder (Miyachi ML2150A) is used to perform the laser spot weld for these stainless steel AISI
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304L metallic cases. The laser beam is delivered via three graded index optical fibres of 0.4 mm core and three laser processing heads. These laser processing heads are arranged at 120 degrees from each other. The laser beams point at the same centre of the welding target and project onto the target surface at 45 degrees. The laser processing head is of 35 mm diameter and has an optical magnification of 10:7 and working distance of 87 mm. All laser processing heads are mounted on the same framework so that they can be moved up and down together. The laser processing heads are aligned so that all three beams focus on the same horizontal plane and point at the same central vertical axis of the welding target. The laser spots will be fired at the junctions of the corresponding welding interfaces. The laser pulse energy from each laser processing head should be adjusted to within 5% variation. The welding between the Z metallic case and bottom metallic case will be done first. The desirable power density of each laser pulse for the laser weld of this junction is ~ 0.17 MW/cm.2 After this welding is done, the distance between the ball-lens to fibre tip will then be fixed. For most of the time, the active alignment along the lateral direction may need to be done again in order to retain the maximum optical coupling efficiency. Before performing the welding for the X-Y metallic case and the Z metallic case, a preload will be applied gradually on top of the X-Y metallic case compressing it against the Z metallic case. A compressive preload at a minimum of 2 kgf is needed to produce a compressive interfacial pressure (> 1.2 MPa) between the base of X-Y metallic case and top of Z metallic case in order to minimize the PWS after welding. The laser processing heads mounted on the same framework will be elevated and the laser beams will point to the interface between the X-Y metallic case and the Z metallic case. The positions of the welding spots are located at a position less than +/– 0.04 mm from their nominal values as shown in Fig. 7.9. After this laser weld is performed, the X-Y metallic case and Z metallic case will be joined. The holder clamping the top X-Y metallic case will then open and release the package. The package is turned by 30 degrees about its central axis for performing a security weld. The desirable power densities of the laser pulses are ~0.2 MW/cm2 and ~0.3 MW/cm2 for the normal weld and security weld of this junction, respectively.
7.6
Future trends
Automatic assembly equipment should be built in accordance with the requirements of each dedicated assembly process. A set of specialized equipment connected in-line will be a favourable choice for the production of a specific product if the reliability of the equipment is high and the product conversion rate is low. The diversity and complexity of assembly processes will be challenges for all equipment suppliers. The process window
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of highly specialized equipment needs to be optimized so that high production yields and high throughputs for specific assembly processes can be achieved. The basic system platform of the equipment may deviate from the standard design if a highly customized product is needed. The disadvantage of some highly customized equipment is usually its low production volume. The manufacturer may not have a chance to perfect the design for some immature product. In addition, the process window for the designated assembly process may need further fine tuning as well. Mature assembly equipment needs to be tested by a volume production run so that the yield, throughput, down time, MTBA, MTBF, etc., can be verified. Equipment suppliers may prefer to build the equipment based on their standard system platform so that the reliability and stability of all electrical components such as power supply, motor, motor drivers, sensors, motion controller, etc., can be very high. In addition, if the equipment is built from the same system platform, in principle, the human-to-machine interface (HMI) across the equipment can be very similar and it is easier for an operator to accustom to the new equipment. On the other hand, end users are more inclined to use highly customized equipment that fits their needs. Mass customization of specific assembly equipment will be the solution as well as the challenge for equipment suppliers. A production line composed of in-line customized assembly equipment facilitating specific assembly processes becomes the ideal solution for volume production of specialized products such as miniature CMOS cameras, optoelectronic packages, smart sensors, RFID products, bio-medical devices, etc.
7.7
References
1. George G. Harman, Wire Bonding in Microelectronics – Materials, Processes, Reliability, and Yield, 2nd edn., McGraw-Hill, New York, 1997. 2. Yu Hin Chan, Jang-Kyo Kim, Deming Liu, Peter C.K. Liu, Yiu Ming Cheung, and Ming Wai Ng, ‘Wire Bondability of Au/Ni Bond Pads: Effects of Metallisation Schemes and Processing Conditions,’ Proceedings of International Symposium on Electronic Materials and Packaging 2002, pp. 69–76, 2002. 3. Arthur C.M. Chong and Cheung Y.M. ‘Finite Element Stress Analysis of Thin Die Detachment Process,’ Proceedings of the 5th International Conference on Electronic packaging Technology (ICEPT) 2003, pp. 44–51, 2003. 4. Cheung Y.M. and Arthur C.M. Chong, ‘New Proposed Adhesive Tape Application Mechanism for Stacking Die Application,’ Proceedings of the 5th International Conference on Electronic packaging Technology (ICEPT) 2003, pp. 309–315, 2003. 5. John H. Lau and Yi-Hsin Pao, Solder Joint Reliability of BGA, CAP, Flip Chip and Fine Pitch SMT Assemblies, McGraw-Hill, New York, 1997. 6. For example, Mitsubishi InGaAsP-MQW-FP laser diode model number: ML725C8F. 7. Cheung Y.M. and Yiu C.H. ‘Simulation of the Alignment Sensitivity on the Coupling Efficiency of a Ball-Lens Capped TO-Can Laser Diode Source into a Single-Mode Fiber’, Proceedings of International Symposium on Electronic Materials and Packaging 2002 (EMAP2002), pp. 197–203, 2002.
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8. Yasushi Fujimura, Yoshiki Kuhara, and Naoyuki Yamabayshi, ‘Optical-axis Alignment Method, Optical-axis Alignment Device, Inspection Method of Optical Devices, Inspection Device of Optical Devices, Method of Producing Optical Module, and Apparatus of Producing Optical Module’, US Patent No. 5,666,450, 9 September 1997. 9. Yiu Ming Cheung, Peter Chou Kee Liu and Ching Hong Yiu, ‘Apparatus and Method for Active Alignment of Optical Components’, US Patent No. 6,886,997, 3 May 2005. 10. Min Kyu Song, Seung Goo Kang, Nam Hwang, Hee Tae Lee, Seong Su Park and Kwang Eui Pyun, ‘Laser Weldability Analysis of High-Speed Optical Transmission Device Packaging’, IEEE Transctions on Components, Packaging, and Manufacturing Technology – Part B, vol. 19, no. 4, pp. 758–763 (1996). 11. Yiu C.H. and Cheung Y.M. ‘Measurement of Post Welding Shift for TO-Can Pigtiail Package’, Proceedings of The 5th Pacific Rim Conference on Lasers and ElectroOptics 2003, pp. 585, 2003. 12. Cheung Y.M. and Yiu C.H. ‘Transient characteristics of post-welding shift of TOcan pigtail package’, Proceedings of 5th International Conference on Electronic Packaging Technology 2003, pp. 221–224, 2003. 13. Yiu C.H. and Cheung Y.M. ‘Investigation into Transient Characteristics of PostWelding Shift for Optoelectronic Pigtail Packages’, Proceedings of the 55th Electronic Components and Technology Conference, pp. 259–265, 2005. 14. Soon Jang, ‘Automation manufacturing systems technology for opto-electronic device packaging’, Proceedings of the 50th Electronic Components and Technology Conference 2000, pp. 10–14, 2000.
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8 Microelectronics wire bonding I L U M, M M A Y E R and Y Z H O U, University of Waterloo, Canada
8.1
Introduction
Wire bonding is the most utilized technique for forming electrical interconnections to integrated circuits in the microelectronics industry with many trillion wires bonded annually [1]. Wire bonding is widely accepted because of its flexibility and ease of use compared to competing methods. In the wire bonding process a small diameter metal wire (usually 25 µm diameter Au) is bonded with a tool firstly to a metal layer on the microchip (usually Al) and then to a metal layer on the package (leadframe or substrate) thus forming the interconnection. The wire material Au is preferred for its excellent ductility and resistance to corrosion and oxidation. The most important variant of wire bonding is thermosonic ball bonding in which the stage is heated up to temperatures of 100 °C to 250 °C. The thermal energy is combined with ultrasonic vibration and a normal bonding load to form the bond. This is a refinement of the process that was first introduced as thermocompression wire bonding in the 1950s at Bell Labs in which only thermal energy was combined with a normal bonding load [1]. In the 1960s ultrasonic welding was described [2] and applied to the wire bonding process. By adding ultrasonic energy, bonding times are shortened and lower temperatures can be used which prevents damage to the devices. These two benefits have resulted in the wide implementation of the thermosonic ball bond process in industry. Since that time the process has seen many improvements related to the equipment and in particular the throughput rate. A significant change in technology is the introduction of higher ultrasonic frequencies [3, 4]. Before the 1990s the ultrasonic frequency used was generally about 60 kHz. The reason for the selection of the original frequency was not so much through research but the fact that it worked [1]. Current modern autobonders utilize an ultrasonic frequency in the range of 100 kHz to 140 kHz. The advantages of such higher frequencies are a faster bond formation and higher quality bonds at lower temperatures. 205 WPNL2204
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As wire bonding is a serial process, it requires a large investment in equipment to cope with the throughput in chip manufacturing. In some cases, the throughput of one chip bonder is distributed to eight wire bonders. Consequently, the industry has put substantial efforts into the development of alternative processes to replace wire bonding. One of these processes is flip-chip bonding in which the electrical contact pads of the integrated circuit are soldered directly to a set of corresponding terminals on a substrate. However, wire bonding remains an important and probably the leading interconnection production technology in microelectronics packaging. The concept of “wire” to electrically connect two points in space has proven irreplaceable in other industries also. The wire bonding industry serves a multi billion dollar market which includes bonding equipment and consumables such as bonding tools and bonding wire. Most of the production of wire bonds is performed with fully automated wire bonders and processes that are robust enough to run unattended for several hours. Section 8.2 describes the process of thermosonic ball bonding and discusses other variants of the wire bonding technique. In Section 8.3 the main process parameters ultrasonic power, bonding force, bonding time, and bonding temperature are described and their effects on the bonding quality is discussed. In Section 8.4 the physicochemical mechanisms responsible for bond formation are described. Wire bonding quality control as it is used in production environments is discussed in Section 8.5. Equipment and fixturing for the wire bonding process are described in Section 8.6. Finally, future directions in wire bonding are discussed in Section 8.7.
8.2
Wire bonding process
8.2.1
Basic description
In a typical thermosonic ball bonding process, a 25 µm diameter Au wire is bonded onto Al metallized bond pads and a Ag plated leadframe finger. Figure 8.1 shows ball bonded Au wires on a Au metallized substrate. A spool of 25 µm diameter Au wire is shown in Fig. 8.2. Both ultrasound and thermal energy are used to form the bond. A capillary tool typically made of Al2O3 ceramics is used to provide the normal bonding load to the wire. The capillary is clamped into the horn as shown in Fig. 8.3. A stack of piezoelectric transducer discs is attached to the base of the horn and used to produce ultrasound. This is discussed in further detail in Section 8.3.1. The ultrasound propagates as a longitudinal wave along the length of the horn and as a transverse wave along the capillary as shown in Fig. 8.4. Vibration nodes appear in the horn and may appear in the tool. The number and location of nodes depends on both frequency and tool geometry. An oscillating tangential
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Ball bonds
8.1 Top view of Au ball bonds on Au metallized substrate.
8.2 Spool of 25 micron diameter Au wire mounted on an autobonder.
displacement results at the tip of the tool. The thermal energy is usually provided by heating cartridges in the stage which holds the chip and substrates or leadframes (bond materials). The bonding process can be described with the aid of Fig. 8.5. The process steps are: 1. A free air ball (FAB) is formed on the end of the wire protruding from the capillary tip. This is done with an electric arc in an electrical flameoff (EFO) process. The molten Au forms a ball due to surface tension before it solidifies. An air column drags the wire up and ball with it until the ball is held by the capillary tip where the capillary hole is small enough to stop the ball from moving through. The capillary is moved to and brought down to the first bond site (pad).
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Clamps
Wire Capillary
x, y
Ultrasound transducer Horn
Chip
Support structure
Wire clamp
Ultrasonic horn Capillary
Electrode
8.3 Schematic of bonding equipment setup showing horn with clamped capillary. Lower image showing photograph. Longitudinal wave
Ultrasound transducer
Transverse wave
8.4 Illustration of ultrasound propagation along horn and capillary.
2. The capillary pushes the ball to the pad. A normal bonding force is applied and after a set amount of time ultrasonic vibration is switched on for a set amount of time, typically between 5 and 20 ms, forming the bond. 3. The capillary is then raised, with the ball bond (first bond) remaining on the bond pad, and the wire threading through the capillary. 4. The capillary moves to the second bond site (lead) along a special trajectory forming a wire loop.
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Heat 2
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8.5 Illustration of ball bonding process steps.
5. The capillary is brought down to the second bond site. The crescent bond (second bond) is formed using a similar set of parameters as are used for the first bond. The parameter values, however, can differ significantly. 6. The capillary rises to a preset height where it waits until the wire clamps close and hold the wire tightly but without plastically deforming it. The clamps then rise to form a tail by breaking the wire at the weakest point which normally is next to the crescent bond. The wire end is called the ‘tail’ protruding from the capillary tip.
8.2.2
Process variants
Ball bonding accounts for most of today’s wire bond production. However, there are significant variations on the wire bonding technique depending on the specific application and different wire materials and metallization materials combinations used. Three of them are listed in Table 8.1 and described
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Table 8.1 Summary of the main wire bonding techniques/materials Procedure
Standard wire material
Main applications
Shielding gas
Thermosonic ball Ultrasonic wedge
Au Al
None None
Cu ball bonding
Cu
Microelectronics Power devices, automotive Microelectronics
95%N-5%H
below. Commonly used wire materials are Au, Al, and Cu which all have in common high ductility. The material combinations that can be bonded are usually ductile metals, e.g. Au wire on Al metallization or Cu wire on Ag metallization. Ultrasonic wedge–wedge bonding was the first ultrasonic wire bonding technique developed [1]. In this process there is no ball formed on the end of the wire, which is usually Al, and the wire itself is pressed sideways with a tool against the bond location. Normally, wedge bonding is performed at ambient temperatures unless Au wire is used. Figure 8.6 shows a schematic explaining the process steps. The tool is called a wedge and is shown in Fig. 8.7. It is typically made of tungsten carbide. Thick wires with diameters up to 500 µm or higher are used to connect power devices with this process variant. In ball bumping, balls are bonded and the wire is broken right at the neck of the ball. No loops are formed. The wire portion remaining on the bond site is called a ball bump. Several bumps can be placed on top of each other. One application of ball bumping is to form studs for flip chip interconnects. A second application aims to improve second bond quality by adding a bump onto the second bond. In a third application a bump is placed on the chip followed by connecting it with the substrate by placing the first bond on the substrate, and the second bond on the bump previously bonded on the pad of a chip. This allows for extremely low loops but still assures that the wire stands off the chip surface as the bump serves as a spacer. The bumping process follows the first three steps of the thermosonic process but after the ball is bonded onto the substrate the clamps close and the wire is broken leaving a bonded ball on the substrate. Sometimes, the capillary is moved horizontally when located above the bonded ball to weaken the wire before it is broken. Figure 8.8 shows bonded ball bumps. Ultrasonic ribbon bonding is a technique similar to ultrasonic wedge bonding. However, the round wire is replaced with a ribbon and a different bonding tool is used. Compared to round wires, wide and thin ribbons allow bonding of larger cross sections and the creation of larger contact areas. Compared to round wire, the plastic surface extension of the ribbon during bonding is achieved by increased force and ultrasound parameters. This
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Tail
Tail
8.6 Illustration of wedge-wedge bonding process steps.
Ultrasonic horn
Wedge tool
Wire clamp
8.7 Wedge tool used in ultrasonic wedge-wedge bonding.
makes it easier to create highly reliable bonds with round wire. However, the ribbon bond connection has a large wire-chip contact interface which is particularly suitable for power applications where the wire-chip interface
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70 µm
8.8 SEM of ball bumps.
carries large currents [5]. Furthermore, a ribbon can transmit higher frequency signals due to its high circumference to cross-sectional area ratio compared to that of round wire.
8.3
Process parameters
The major parameters in today’s wire bonding processes are ultrasonic power, bonding force, bonding time, and bonding temperature. They are described in the following sections.
8.3.1
Ultrasound
Ultrasound is defined as sound with a frequency above 12 kHz and is thus inaudible to the human ear. Ultrasound has many applications in industry such as medical imaging or ultrasonic cleaning. In modern wire bonding, ultrasonic vibration is fundamental to the process. The design of an ultrasonic system for wire bonding consists of an electronic power supply, control system, piezoelectric ultrasonic transducer, horn, and clamping mechanism to clamp the bonding tool. Piezoelectric material creates a force when voltage is applied to it. An ultrasound generator delivers a.c. current to the piezo-ceramics. A longitudinal ultrasonic wave is produced by the expanding-contracting stack of piezo-ceramics. The wave extends along the horn as shown in Fig. 8.4. The ultrasonic vibration is transmitted from
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the horn to the tool which needs to be clamped tightly to the horn. The ultrasonic wave is transverse in the tool. The horn is designed to amplify the vibration amplitude to larger values at its end. The frequency of the a.c. voltage applied to the piezo-stack is matched to a longitudinal mechanical resonance of the horn. Typical wire bonders use a fixed ultrasound frequency between 60 kHz and 140 kHz that cannot be changed. Vibration nodes may exist in the bonding tool depending on the applied frequency and tool geometry. The ultrasonic oscillating tangential displacement at the tool tip is fundamental in the ultrasonic bonding process. The vibration amplitude of the bonding tool can be measured directly with a laser interferometer or a capacitor microphone [6]. The vibration amplitudes of both unloaded and loaded tools are proportional to the power setting of the ultrasonic generator. The technical control parameter of ultrasound generators is either the electrical power, current, or voltage amplitude delivered to the piezo-ceramics. Modern generators usually control current as it is proportional to the vibration amplitude. Current control compensates for impedance variations observed among a group of otherwise identical bonders, resulting in better parameter portability from bonder to bonder. Increasing ultrasonic power increases bond strength [7]. However, too high ultrasonic powers lead to bonded balls that are badly deformed. Ultrasound contributes important mechanisms to modern wire bonding. One is the “oscillating tangential displacement”, and the second, “ultrasonic softening”, is a by-product of the ultrasound and it is not conclusive as to its requirement in bond formation. A third is the “ultrasonically enhanced interdiffusion” observed when bonding dissimilar materials [8]. The oscillating displacement of the tool tip causes sliding friction between wire and bond area which is a fundamental requirement in the ultrasonic wire bonding process. The friction aids in the cleaning of the surfaces and the wearing down of surface roughness, allowing subsequent intimate metal/ metal contact and bonding. While applying ultrasound the yield point of a metal apparently decreases, an effect referred to as “ultrasonic softening” [9]. Applied ultrasound has a similar effect in lowering the yield point as applied thermal energy. However, the ultrasound softening effect can be achieved at much lower applied energy levels as compared to thermal energy. This higher efficiency of lowering yield stress is suggested to be due to the attenuation of ultrasonic energy only at the points that affect deformation such as dislocations and vacancies [10] whereas thermal energy is uniformly distributed across the bulk material. This softening occurs throughout the frequency range of 20 kHz up to 1000 kHz. Usually the material shows no residual softening once ultrasound is removed. With higher applied ultrasonic powers the apparent yield stress decreases more. However, at very high power levels a residual hardening effect occurs [10]. The material can thus be permanently hardened.
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In opposition to this observation, a high enough level of ultrasound can result in a permanent softening of an Al wire used in ultrasonic wedge– wedge bonding as reported in [11]. In this case the ultrasound causes a recrystallization of the bonding wire near the bond interface. Measurements show that the recrystallized microstructure has a decreased microhardness. Possibly, the ultrasound causes frictional heating and annealing of the wire.
8.3.2
Bonding force
Ultrasonic current control signal
Force control signal
A force normal to the bond site is used to clamp the wire to the bond site to assure intimate contact during the application of ultrasound (bond force, bonding force, clamping force). All modern bonders have at least two stages in the bonding force profile [12]. These two stages are typically controlled using two parameters which are termed impact force and bond force. Some bonders use the parameters contact velocity and bond force instead. The contact velocity in combination with the touchdown detection determines the actual impact force. The two bonding forces are shown in the bonding force profile plot in Fig. 8.9. Standard industry practice is to select the impact force to be about 1.5 to 2 times the value of the bonding force. Generally the impact forces result in most of the ball deformation and during the subsequent ultrasound application little additional deformation of the ball occurs [13]. The normal bonding force is chosen at a sufficiently high level so that enough friction power is available for sufficient wear and bonding. Too high forces decrease bonding by limiting the amplitude of interfacial sliding and
Bond time Impact force Bond force
Ultrasound time
Ultrasound amplitude
Time
8.9 Parameter settings and control profiles during a wire bond.
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therefore friction power of the ball [14, 15]. A too large decay of the tool tip amplitude due to the higher friction force can be compensated by increasing the ultrasound parameter [16].
8.3.3
Bonding time
Bonding usually is completed after 10 to 20 ms of ultrasound for Au ball bonding on Al metallization but depends on other parameters and bonding materials. Further increases in bonding time do not increase the amount of bonded area when Au wire is used [14, 15]. As proposed by [17] the bonding progresses with the total sliding distance made by the oscillating relative displacement. This distance mainly depends on the relative ultrasonic amplitude and the ultrasonic frequency. While the bond strength develops there is little ultrasonic deformation of the ball. Once a certain bond strength is achieved, the additional stress field from the ultrasonic force causes the wire to yield again and deform more [14, 15]. Shorter bonding times are preferred for increased productivity in industrial applications. However, when bonding on materials such as copper with its tenacious oxide layer, a longer bond time may be useful. The longer bond time allows for more ultrasonic cleaning and frictional energy to be delivered to the surfaces to wear off the oxide [18].
8.3.4
Bonding temperature
The additional energy provided by elevated temperatures assists in the bond formation and increases the bonding throughput. In some cases, increased temperatures aid in eliminating surface contaminants [19]. However, elevated bonding temperatures may introduce an oxide layer on many materials such as copper which reduces bondability [20]. When bonding on Ag plated Cu lead frames, typical bonding temperatures are between 220 °C and 240 °C. When bonding on polymer based substrates, the bonding temperatures used are limited by the polymer and are usually about 120 °C to 150 °C. Temperatures that are too high may damage the components being bonded. Temperature microsensors can be used to measure the average temperature very close to the bonding interface [21, 22]. The microsensor in [21] consists of aluminium lines around a test bond pad. Changes in temperature cause a change in the electrical resistance of the aluminum lines. By calibrating known temperatures to the measured resistances the temperatures can be measured. The temperature change measured by such a microsensor typically is not more than 1 K. The response time is shorter than 1 ms. In [23] a finite element model is developed to estimate the average contact temperature if the microsensor temperature is known. Not more than 10 K average temperature
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increase is expected at the contact zone. However, momentary flash temperatures at asperity contact points can be considerably higher than the average contact temperature [24], a fact which has not been experimentally verified in wire bonding.
8.4
Mechanisms of bond formation
Although there is recent research [25–30] describing some aspects of some mechanisms, there is still no full explanation of any of the wire bonding mechanisms, nor is there any analytical mathematical model or “universal wire bonding formula” explaining any of the bonding mechanisms. However, a better understanding of the mechanisms will lead to the next level of process improvements and solutions for the current bonding challenges. For example, optimized values for ultrasonic frequency could be suggested, optimized time-dependent process profiles for ultrasound and force parameters could be calculated, or improved on-line feedback based real-time process control systems could be developed [31, 32]. It is widely accepted that the wire bonding process is a solid state process. In order to create a bond between two metals the surfaces must be relatively contaminant free [33]. In Section 8.4.1–8.4.3, an overview is given of mechanisms proposed in the literature to account for ultrasonic wire bonding. They include ultrasonic heating, plastic deformation, and ultrasonic wear.
8.4.1
Ultrasonic heating
One of the very first bonding mechanism theories proposed was that of interfacial heating. It is postulated that the interfacial movement caused by ultrasonic energy caused rubbing, heating and then even melting of the interfaces resulting in bonding [34]. A large quantity of experimental evidence shows that although there is a temperature rise, that bond temperatures do not approach the melting point of the materials. For example bonding performed at liquid nitrogen temperatures shows that no bubbles in the liquid nitrogen are observed, indicating that temperature rise is very low [7, 35]. Therefore, it is concluded that the observed temperature rise does not create the bonding but is rather a by-product of the bonding process. A TEM study which measured diffusion distances of vacancies to the grain boundaries indicates that bond temperatures are no more than about 250 °C [36]. Furthermore, examination of the bond interface in the same TEM study indicates no evidence of melting. Melting may be present in wire bonds made under non-optimal conditions of too high ultrasonic power producing overdeformed bonds. In this case the friction energy is large and temperature rise may be substantial. In particular if bonding on Sn – a material melting at a relatively low temperature –
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melting may be evident when studying the bond interface [37]. However, at optimal bonding parameters there is no melting observed. Even when bonding on Sn which has a low melting point, under normal conditions there is no melting. The evidence suggests that melting is not required for successful ultrasonic wire bonding.
8.4.2
Deformation
Another mechanism to explain the bond formation in ultrasonic bonding is the deformation mechanism similar to that described as cold welding. This theory postulates that the interfacial extension due to the lateral deformation of the wire results in displacement of contaminants, bringing new metal surfaces to the interface which will readily bond with each other. Reference [38] reports on a numerical analysis of the interfacial contact process in thermocompression wire bonding. It is concluded that wire height reduction by a factor of 0.5 is required to produce a strong peripheral bond. The result is comparable to the deformation required for cold pressure welding. Reference [35] reports on the similarity of the deformation in thermocompression and thermosonic bonding and suggests that the only difference in the two processes is the form of energy used to cause plastic flow. However, this is a simplified view of the effects of ultrasonic energy as studies show that bonds made with the same amount of deformation with and without ultrasonic energy demonstrate different bond strengths for both Au ball bonding [39] and Al wedge bonding [26]. This shows that deformation alone cannot be used to explain the thermosonic bonding process completely.
8.4.3
Wear by microslip
The contact mechanics governing the wire bonding process can be approximated to a certain extent by the classical contact mechanics between two ideal elastic spheres [40] or between an elastic sphere and an elastic plane [41]. When two surfaces in contact, are subjected to both normal force (N) and tangential force (S), gross sliding on a macroscopic scale initiates when the tangential force exceeds a critical value, µsN, where µs is the coefficient of static friction. If the tangential force is less than the critical value, no apparent displacement occurs. However, the tangential force will set up strains so that a minute tangential displacement of one body relative to the other occurs at some parts of the interface [41]. Such a partial relative displacement is called “microslip” between the two sliding partners. Shear traction q without slip would rise to infinity which obviously is not attainable. Therefore the material slips to relieve the stress. This microslip at the periphery is in the order of 0.25 to 2.5 µm [42].
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Reference [40] provides an equation to calculate the size of the microslip annulus observed at the periphery of the contact zone, 1
S 3 a′ = a 1 – µ N s
(1)
where a′ is the inner radius of the annulus and a is the outer radius which is also the contact radius. There is no slip occurring over the circle of radius a′ known as the stick region. Figure 8.10 shows a schematic of the stick and microslip regions in a circular contact undergoing microslip. The slip annulus grows inwards with increasing tangential force up to the point of gross sliding at which point the microslip annulus has grown to the center of the contact circle. In reference [41] elastic spheres are pressed against elastic plates and loaded with oscillating tangential forces. In the absence of gross sliding, a fretted annulus due to microslip occurs on the surface of the plate. As the magnitude of the tangential force increases the inner radius of the annulus grows inwards up to the point of sliding as predicted by Equation 8.1.
Transition line
Relative motion Tangential force (ultrasonic power)
a
B
a Gross sliding regime
a′ Microslip annulus
a
a′
A
Stationary Microslip regime Contact diameter Normal force, N
8.10 Illustration of sticking and sliding in a circular contact under constant normal and increasing tangential load, showing transition from microslip to gross sliding.
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With the oscillating relative motion of the contacting surfaces fretting wear occurs [43]. Fretting wear is common in industry for clamped connections [44, 45]. Wear of material [46] can be quantified according to:
t=d H 1 K PV
(2)
where t is the time required, d is the depth of material worn, P is the mean or nominal pressure, H is the hardness of the material, K is the wear coefficient constant, and V is the sliding velocity. With the removal of surface material, fresh underlying material is exposed promoting bonding. Reference [47] reports that the ultrasonic energy results in a tangential force being applied in a reciprocating manner. At low vibration amplitudes, microslip initiates at the periphery of the contact zone. Microslip displaces contaminants revealing clean metal. Bonding is the result of adhesion built upon asperity bridges which are created when the clean metal bodies come into contact under the combination of normal and tangential forces. Microslip occurs in a similar way in both ultrasonic ball bonding and ultrasonic wedge bonding as studied with bond footprints in [25, 26], demonstrating that the underlying wear mechanism is significant for both of the main ultrasonic wire bonding techniques. Bond footprints are shown on micrographs of contact zones after removal of the bonded wire or after an unsuccessful bonding attempt that results in some imprint or wear on the surface. Evidence of microslip is observed on such footprints in the form of surface morphology changes or wire residues (“bonding”) remaining on the surface. In this sense, “density of bonding” is the same as density of wire material remaining on the surface. In [25] the large plastic ball deformation and complex bonding tool geometry are taken into account to explain the microslip phenomena observed in the evolution of ball bond footprints. Figures 8.11(a) to (d) illustrate the footprint morphology transitioning from microslip to gross sliding as ultrasonic power increases. The highest normal stress at the ball/substrate interface is located on a ring at about the chamfer diameter and capillary shoulder as shown in Fig. 8.12(b). The lower normal stress at the center area leads to less wear as described by Equation 8.2 compared to the areas under the chamfer diameter and capillary shoulder. The bond formation by wear in ultrasonic wedge bonding is expected to be similar to that described for thermosonic ball bonding as only the wire and bonding tool materials and geometries of the two methods are different from each other. Experimental results of a footprint study of ultrasonic Al wedge bonding at ambient temperatures are explained by a bond development model in [26]. The wedge tool tip is shown in Fig. 8.13(a) and (b). Figure 8.14 depicts a footprint development model for wedge-bonding and illustrates the footprint morphology transitioning from microslip into gross sliding
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(a)
(b) Microslip
(c)
(d)
Gross sliding
8.11 Illustration of ball bond footprint change for increasing ultrasonic power. Shaded areas indicate fretting. Dashed circle indicates the capillary chamfer diameter. Bonding density indicated by cross hatching density (a) and (b) partial fretting, (b), (c), and (d) ultrasonically enhanced deformation, (c) and (d) gross sliding, (c) extensive bonding, (d) substantial amount of wire residues on footprint. Wedge
Capillary
Increasing compressive stress
Wire
(a)
(b)
8.12 Illustration of normal stress distribution at contact interface for (a) wedge bonding and (b) ball bonding.
with an increasing tangential force similar to that observed in ball bonding. Due to the presence of the stress maximum at the center of the tool as shown in Fig. 8.12(a), in wedge bonding the center area of the footprint is mainly bonded as compared to the mainly unbonded central area in ball bonding. A common observation with all wire bond footprints is the onset of microslip at interface locations of low normal stress such as the periphery. Furthermore, microslip can result in strong wear especially when accompanied by large normal stresses as described by Equation 8.2. With increased normal bonding forces the threshold ultrasonic power required for transition from microslip WPNL2204
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(a) Front radius = 25
30° Hole φ = 50 Back radius = 15
100 µm 100 µm
(b) Width = 102
Bond length = 50
8.13 Micrographs of wedge tool tip with geometry definitions (a) side, (b) bottom view. Increasing ultrasonic power
(a)
(b)
(c)
Microslip Gross sliding
8.14 Illustration of wedge bond footprint change for increased ultrasonic power. Shaded areas indicate fretting. Bonding density indicated by cross hatching density (a) partial fretting and little bonding, (b) and (c): gross sliding and ultrasonically enhanced deformation, (b): moderate bonding, (c): substantial amount of wire residues on footprint.
into gross sliding also increased. This is obvious from Equation 8.1 and the law of friction. Without melting occurring, some other method of interfacial contaminant dispersal is required for joining. In wire bonding the ultrasonic vibration causes an oscillating tangential displacement at the wire/substrate interface resulting in wear and removal of the surface contaminants and microdeformation of surface roughness. With increased ultrasonic power there is increased displacement at the interface and subsequent increased
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cleaning of the surfaces. With cleaner surfaces more intimate contact between clean metal surfaces can occur and the result is more bonding.
8.5
Quality control
The bond quality right after bonding is assessed, for example, by visual inspection as standardized in [48]. The bond connections themselves are mainly controlled using destructive sampling tests. There are two main standard tests for wire bonds measured right after the bonding: the pull test and the shear test. The value of the pull test for advanced bonding processes is discussed in [49]. The applicability of current quality assurance methods to future fine pitch bonding processes are under debate [50].
8.5.1
Pull test
A hook as used on a pull tester is shown in Fig. 8.15. It is used for testing the strength of either the first or the second bond for both ball bonding and wedge-wedge bonding. In this test a hook is placed under the wire loop, pulled upwards with a controlled speed, and the force required to break the bond is recorded. The pull test is standardized in [51] and [52]. According to the latter, the required minimum pull force is 3 gf (1 gf = 9.8 mN) for a process with a 25 µm diameter wire. The pull force value is variable depending on where the hook is placed as shown by the force vectors in Figure 8.16. The hook is placed closer to the bond to be tested. For a sufficiently strong wedge bond the wire breaks in the heel region of the bond. When testing a moderately strong Au ball bond, the wire breaks in the neck region of the ball. Therefore, the strength of some ball bonds cannot be
Substrate
Hook
8.15 Pull test hook.
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fully characterized with the pull test. However, a lower quality limit can be assured and the pull test remains in use for assuring the first and second bond quality of many production processes.
8.5.2
Shear test
A close-up of a shear tester is shown in Fig. 8.17. In this test a shear ram (shear tool, shear chisel) is used to shear through the bond or the wire at a fixed height from the surface of the bond pad (typically about 3 µm). The shear test is standardized in [53]. It measures the maximum force required to shear the bond. A sufficiently bonded ball is sheared through the ball material leaving a layer of the ball material bonded onto the substrate. A poorly bonded ball leaves little material behind when sheared. It effectively delaminates Force measured by pull tester
Force at 1st bond
Force at 2nd bond
1st bond Chip
Wire loop during pull test
2nd bond
Metallization
Substrate
8.16 Force balance during pull test.
Substrate
Shear tool
8.17 Shear tool (shear ram).
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from the interface. The industry minimum strength required depends on the material combination used, the diameter of the bond connection, and the standard deviation of the shear force results. The shear strength is calculated as the shear force divided by the nominal bond area. Alternatively this test can be used to shear through a wedge bond to test the shear strength of the wedge bond itself. However, this is not a standardized test.
8.6
Equipment and fixturing
8.6.1
Autobonders
Figure 8.18 shows an advanced automatic ball bonder. Autobonders are fully automated bonders that when provided with a bonding program will perform continuous bonding without operator intervention except when production stoppages arise. Bond programs will include all items required to perform
8.18 Photograph of autobonder.
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the bonding including bond positions and parameters and material handling. Each bond can have its own specific parameters. An advanced vision system comprises a camera mounted to the bondhead as shown in Fig. 8.19 and software used for programming and chip alignment before bonding. A computer monitor displays both the computer interface and vision system information. A microscope is fixed to the bonder and used for viewing the bonds. Autobonders have automated material handling systems which move material from the input stage through the bonding area and into the output (storage) stage. During the bonding the material can be held in place by vacuum and an upper clamping plate. For production, material is usually provided in strips with several dies attached on them. In automatic ball bonders it is the bondhead that moves to perform bonding while the material remains stationary. However, in some automatic wedge bonders, the material may be rotated since the wire loop is required to be parallel to the horn.
8.6.2
Semi-automatic bonder
An example of a semi-automatic bonder is shown in Fig. 8.20. Such bonders are used for research and development and small quantity production. A bonding stage is used to clamp and heat the material and perform bonding on. Clamping can either be by vacuum or spring loaded clamping tabs and is adjusted manually. The bonding stage is placed on the x–y table which is moved by the operator via a mechanical micromanipulator. By observing through the microscope and operating the micromanipulator the operator locates the desired bond position and triggers an automatic bonding procedure by pressing a button. In contrast to the auto bonders the bonding tool remains fixed in x–y space. The process parameters are selected by buttons or dials on the machine and may be shown on an LCD display. The number of
Capillary Camera
Electrode Window Clamping plate
8.19 Front view of autobonder bondhead.
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parameters available with a semi-automatic bonder is much smaller than that of the parameters available with an autobonder. They are limited to ultrasonic power, bonding time, bonding force, search height, and loop height.
8.7
Current directions
8.7.1
Cu wire
A recurring trend has been the replacement of Au wire with Cu wire [54–57]. There are many benefits of Cu wire such as lower cost and a more than 30% higher electrical conductivity than Au. However, there are challenges with bonding of Cu wire. Cu wire unlike Au requires a forming gas to successfully form free-air balls as shown in Fig. 8.21(a). Typically a 0.2 l/min. flow rate of 95%N-5%H is used for 25 µm diameter Cu wire and delivered directly to the tail prior to EFO. In ambient atmosphere the ball formation will result in oxidized and misshaped balls as shown in Fig. 8.21(b). Cu wire also readily oxidizes in ambient atmosphere as opposed to Au. The wire becomes unbondable a few days after exposing the wire spool to air. Wire spool
Alignment aid
Micromanipulator
Horn
Stage with heater
8.20 Photograph of semi-automatic bonder.
(a)
(b)
8.21 Free-air ball formed with Cu wire (a) with 95%N-5%H forming gas, and (b) in ambient air.
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A main challenge when bonding with Cu wire is the higher hardness of Cu compared to Au which can lead to defects such as chip cratering [58, 59] in processes that work fine with Au wire. Another challenge is the tail formation which can cause frequent production stops if it is inconsistent. The different microstructural and material properties of Cu compared to Au may be the reason for a larger variability in the breaking force required to break the wire after the wedge bond [60]. The tail bond controls the length of tail available for the FAB formation. Weak tail bonds will lead to the wire lifting off of the substrate before the tail bond breaking stage and the tail will recede into the capillary resulting in stoppage of the process due to the absence of the tail.
8.7.2
Bonding on low-k substrates
Low-k materials can lack the mechanical stability to survive the wire bonding process. Low-k materials are becoming more common as insulation layers under the metallization layers [61–63] in microchips. The variable k is used for the dielectric constant of an insulating material. A lower k-value means better insulation properties with air being the best insulator with a k-value equal to one. As the trend of signal speeds carried by the interconnection wires increases a better insulating material (lower k-value) is required to eliminate the cross talk between wires. The standard insulation is SiO2 with a k equal to 4.1 [63]. New lower k materials being used with k-values less than 3 (typically about 2.7) [63] possess much lower thermo-mechanical stability than SiO2. In order to use the wire bonding process on low-k materials, damage to the low-k material is prevented either through adjustment of the bonding process or the structure of the low-k material, or – most promising – adjusting both.
8.7.3
Insulated bonding wires
As miniaturization combined with increased number of signal connections on a microchip continues unabated into the future, insulated (coated) bonding wires offer a viable solution to meet the challenges [64]. The main feature of insulated bonding wires is the insulation on the wire which prevents wires from short circuiting when touching and enables complex package designs. Insulated bonding wire has undergone extensive reliability testing and passes [65]. The existing wire bonding infrastructure can be used for bonding insulated wire with no major equipment upgrade required to use this new technology. As shown in Fig. 8.22(a) a spherical FAB is formed with insulated wire which is essential for reliable bonding. A FAB formed with insulated wire shows characteristic “striping” in the FAB, however the bottom half of the ball does not show the “striping” and the ball bond passes reliability tests [65]. Figure 8.22(b) shows a stitch bond made with insulated bonding wire.
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The stitch bond process parameters are adjusted from the bare wire parameters in order to form a stitch bond with insulated bonding wire due to the requirement of the removal of the insulating coating. However, the parameters remain within the bonding parameters window [65]. Insulated bonding wire is beneficial for both current and future bonding applications. In current applications the use of insulated bonding wires can increase product yield by eliminating wire shorts resulting from wire sweep during the encapsulation stage. Ultra high density single chip applications such as fine pitch and multi-row (as shown in Fig. 8.23 where the touching wires do not short circuit) wire bonding densities can be further increased by using insulated bonding wires. The performance of multi-chip packages can
(a)
(b)
8.22 (a) Free-air ball formed with insulated bonding wire (b) stitch bond made with insulated bonding wire. [64] Ball bonds
8.23 Multi-row bonding using insulated bonding wires showing wires crossing and touching without short circuiting. [64]
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be improved by using insulated bonding wires to make direct chip-to-chip connections, bypassing substrate routing and vias, and thus improving high speed performance [65]. Most promising is the enabling of wire bonding for new applications such as complex multi-tiered stacked die and system in package (SiP).
8.7.4
Finer pitches
With the trend of wire densities becoming larger on chips, the ball bond pitches become smaller. As capillaries with finer and finer tips become available, bond pitches are moving below 45 µm [66, 67]. In such processes, the bonded ball shapes become less spread out sideways, minimizing the ratio of footprint to wire diameter. In addition to the increased requirements of bond-to-bond parameter and positional repeatability, a current challenge of fine pitch ball bonding is reduced reliability, sometimes affected by the marks left on the pads after probing the chips. The ball bond reliability is reduced by interfacial cracks that form during the longterm test due to interdiffusion of atoms. Such cracks often cover too much of the interface before the 1000 h at 175 °C required by, for example, a JEDEC standard. Efforts to improve the reduced bonding performance for a 35 µm bond pad pitch process include the reconsideration of the capillary design and bonding wire formulation [66, 67].
8.8
References
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9. B. Langenecker, ‘Effects of Ultrasound on Deformation Characteristics of Metals,’ IEEE Transactions on Sonics and Ultrasonics, Vol. SU-13, No. 1, pp. 1–8, Mar. 1966. 10. B. Langenecker, ‘Effect of Sonic and Ultrasonic Radiation on Safety Factors of Rockets and Missiles,’ AIAA Journal, Vol. 1, No. 1, pp. 80–83, Jan. 1963. 11. U. Geissler, M. Schneider-Ramelow, K. D. Lang, H. Reichl, ‘Investigation of Microstructural Processes During Ultrasonic Wedge/Wedge Bonding of AlSi1 Wires,’ Journal of Electronic Materials, Vol. 35, No. 1, pp. 173–180, Jan. 2006. 12. R. G. McKenna, R. L. Mahle, ‘High Impact Bonding to Improve Reliability of VLSI Die in Plastic Packages,’ Proc. Electronic Components Conference, pp. 424–427, 1989. 13. J. Brunner, B. Chylak, ‘Optimization of the Wire Bonding Process On Low-k Pad Structures,’ Proc. IMAPS Device Packaging Conference, Mar. 2006. (Kulicke and Soffa website http://www.kns.com). 14. M. Mayer, ‘Microelectronic Bonding Process Monitoring by Integrated Sensors,’ Ph.D. thesis, No. 13685, ETH Zurich, Zurich, 2000. 15. M. Mayer, Microelectronic Bonding Process Monitoring by Integrated Sensors, Hartung-Gorre, Konstanz, Germany, 2000. 16. Z. N. Liang, F. G. Kuper, M. S. Chen, ‘Concept to Relate Wire Bonding Parameters to Bondability and Ball Bond Reliability,’ Microelectronics and Reliability, Vol. 38, No. 6–8, pp. 1287–1291, Jun.–Aug., 1998. 17. J. L. Harthoorn, ‘Joint Formation in Ultrasonic Welding Compared with Fretting Phenomena for Aluminium,’ Proc. Ultrasonics Intl. Conf., pp. 43–51, 1973. 18. N. Noolu, I. Lum, Y. Zhou, ‘Roughness Enhanced Au Ball Bonding of Cu Substrates,’ IEEE Transactions on Components and Packaging Technologies, Vol. 29, No. 3, pp. 457–463, Sept. 2006. 19. J. L. Jellison, ‘Effect of Surface Contamination on the Thermocompression Bondability of Gold,’ IEEE Transactions on Parts, Hybrids, Packaging, Vol. PHP-11, No. 3, pp. 206–211, Sept. 1975. 20. G. Hotchkiss, et al., ‘Probing and Wire Bonding of Aluminum Capped Copper Pads,’ Proc. 40th Annual IEEE International Reliability Physics Symposium, Dallas, pp. 140–143, 2002. 21. M. Mayer, O. Paul, D. Bolliger, H. Baltes, ‘Integrated Temperature Microsensors for Characterizaton and Optimization of Thermosonic Ball Bonding Process,’ IEEE Transactions on Components and Packaging Technologies, Vol. 23, No. 2, pp. 393– 398, Jun. 2000. 22. S. Shivesh, M. Gaitan, Y. Joshi, G. G. Harman, ‘Wire-Bonding Process Monitoring Using Thermopile Temperature Sensor,’ IEEE Transactions on Advanced Packaging, Vol. 28, No. 4, pp. 685–693, Nov. 2005. 23. M. Mayer, A. Zwart, ‘Ultrasonic Friction Power in Microelectronic Wire Bonding,’ Proc. THERMEC 2006, Vancouver, Canada, pp. 3920–3925, July 2006. 24. M. F. Ashby, J. Abulawi, H. S. Kong, ‘Temperature Maps for Frictional Heating in Dry Sliding,’ Tribology Transactions (USA), Vol. 34, No. 4, pp. 577–587, Oct. 1991. 25. I. Lum, J. P. Jung, Y. Zhou, ‘Bonding Mechanism in Ultrasonic Gold Ball Bonds on Copper Substrate,’ Metallurgical and Materials Transactions A: Physical Metallurgy and Materials Science, Vol. 36, No. 5, pp. 1279–1286, May 2005. 26. I. Lum, M. Mayer, Y. Zhou, ‘Footprint Study of Ultrasonic Wedge-Bonding with Aluminum Wire on Copper Substrate,’ Journal of Electronic Materials, Vol. 35, No. 3, pp. 433–442, Mar. 2006.
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27. Y. R. Jeng, J. Y. Chen, ‘On the Microcontact Mechanism of Thermosonic Wire Bonding in Microelectronics: Saturation of Interfacial Phenomena,’ Tribology Transactions, Vol. 48, No. 1, pp. 127–132, Jan./Mar. 2005. 28. M. Mayer, J. Schwizer, ‘Ultrasonic Bonding: Understanding How Process Parameters Determine the Strength of Au-Al Bonds,’ Proc. of SPIE – The International Society for Optical Engineering (IMAPS 2002), Vol. 4931, pp. 626–631, 2002. 29. M. Mayer, J. Schwizer, ‘Thermosonic Ball Bonding Model Based on Ultrasonic Friction Power,’ Proc. Electronic Packaging Technology Conference EPTC’03 (IEEE), Singapore, pp. 738–743, 2003. 30. H. Gaul, M. Schneider-Ramelow, K. D. Lang, H. Reichl, ‘Predicting the Shear Force of a Wire Bond Using Laser Vibration Measurements,’ Proc. Electronics Systems Integration Technology Conf., Dresden, pp. 719–725, 2006. 31. B. Gobel, A. Ziemann, ‘Quality Control for Wire Bonding,’ US Patent number: 4984730, Issue date: Jan. 15, 1991. 32. H. J. Hesse, D. Holtgrewe, ‘Quality Control Method,’ US Patent number: 6308881, Issue date: Oct. 30, 2001. 33. ASM Handbook Volume 6 – Welding, Brazing, and Soldering, 1993. 34. V. H. Winchell, H. M. Berg, ‘Enhancing Ultrasonic Bond Development,’ IEEE Transactions on Components, Hybrids, and Manufacturing Technol., Vol. CHMT-1, No. 3, pp. 211–219, Sept. 1978. 35. G. G. Harman, K. O. Leedy, ‘An Experimental Model of the Microelectronic Ultrasonic Wire Bonding Mechanism,’ Proc. 10th Annual Proc. Reliability Physics, Las Vegas, USA, pp. 49–56, Apr. 5–7, 1972. 36. J. E. Krzanowski, N. Murdeshwar, ‘Deformation and Bonding Processes in Aluminum Ultrasonic Wire Wedge Bonding,’ Journal of Electronic Materials, Vol. 19, No. 9, pp. 919–928, Sept. 1990. 37. J. Lee, M. Mayer, Y. Zhou, ‘The Feasibility of Au Ball bonding on Sn Plated Cu,’ Journal of Electronic Materials, accepted for publicaton 2007. 38. Y. Takahashi, S. Shibamoto, K. Inoue, ‘Numerical Analysis of the Interfacial Contact Process in Wire Thermocompression Bonding,’ IEEE Transactions on Components, Packaging, and Manufacturing Technol., Part A, Vol. 19, No. 2, pp. 213–223, Jun. 1996. 39. Y. Zhou, X. Li, N. Noolu, ‘A Footprint Study of Bond Initiation in Gold Wire Crescent Bonding,’ IEEE Transactions on Components and Packaging Technologies, Vol. 28, No. 4, pp. 810–816, Dec. 2005. 40. R. D. Mindlin, ‘Compliance of Elastic Bodies in Contact,’ American Society of Mechanical Engineers – Transactions – Journal of Applied Mechanics, Vol. 16, No. 3, pp. 259–268, Sept. 1949. 41. K. L. Johnson, ‘Surface Interaction Between Elastically Loaded Bodies Under Tangential Forces,’ Proceedings of the Royal Society of London, Series A, Vol. 230, No. 1183, pp. 531–548, Jul. 1955. 42. K. L. Johnson, Contact Mechanics, Cambridge University Press, London, 1985. 43. E. Rabinowicz, Friction and Wear of Materials, John Wiley and Sons, Inc., New York, 1965. 44. N. Ohmae, T. Tsukizoe, ‘The Effect of Slip Amplitude on Fretting,’ Wear, Vol. 27, No. 3, pp. 281–294, Mar. 1974. 45. M. Odfalk, O. Vingsbo, ‘An Elastic-Plastic Model for Fretting Contact,’ Wear, Vol. 157, No. 2, pp. 435–444, Sept. 1992. 46. Wear Control Handbook, ASME, 1980.
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47. G. K. C. Chen, ‘The Role of Micro-Slip in Ultrasonic Bonding of Microelectronic Dimensions,’ Proc. International Microelectronic Symposium, October 30–31, 1972, Washington D.C., pp. 5-A-1-1–5-A-1-9, 1972. 48. ‘Lower Power Wire Bond Inspection,’ Section 3.2.1, Method 2010.10, Standard, as published in MIL-STD 883E, 22 Mar. 1989. 49. V. Sundararaman, D. R. Edwards, W. E. Subido, H. R. Test, ‘Wire Pull on Fine Pitch Pads: An Obsolete Test for First Bond Integrity,’ Proc. 50th Electronic Components and Technology Conference, pp. 416–420, 2000. 50. L. Levine, J. Brunner, S. Haggenmüller, ‘Guidelines for Quality Must Change,’ Proc. EPP Europe, pp. 50–52, Jan. Feb. 2003. 51. ‘Test Method for Pull Strength for Wire Bonding,’ Standard, SEMI G73-0997, 1997. 52. ‘Bond Strength (Destructive Bond Pull Test),’ Notice 4, Method 2011.7, Standard, as published in MIL-STD 883E, 22 Mar. 1989. 53. ‘Test Methods for Destructive Shear Testing of Ball Bonds,’ Standard, ASTM F126906. 54. J. Chen, D. Degryse, P. Ratchev, I. DeWolf, ‘Mechanical Issues of Cu-to-Cu Wire Bonding,’ IEEE Transactions on Components and Packaging Technologies, Vol. 27, No. 3, pp. 539–545, Sept. 2004. 55. M. Sheaffer, L. R. Levine, B. Schlain, ‘Optimizing the Wire-Bonding Process for Copper Ball Bonding, Using Classic Experimental Designs,’ IEEE Transactions on Components, Hybrids and Manufacturing Technology, Vol. CHMT-10, No. 3, pp. 321–326, Sept. 1986. 56. D. Degryse, B. Vandevelde, E. Beyne, ‘FEM Study of Deformation and Stresses in Copper Wire Bonds on Cu LowK Structures During Processing,’ Proc. 54th Electronic Components and Technology Conference, Vol. 1, pp. 906–912, 2004. 57. J. E. J. M. Caers, A. Bischoff, J. Falk, J. Roggen, ‘Conditions for Reliable BallWedge Copper Wire Bonding,’ Proc. IEEE/CHMT European International Electronic Manufacturing Technology Symposium, pp. 312–315, 1993. 58. C. W. Tan, A. R. Daud, ‘Bond Pad Cratering Study by Reliability Tests,’ Journal of Materials Science: Materials in Electronics, Vol. 13, No. 5, pp. 309–314, May 2002. 59. K. Toyozawa, K. Fujita, S. Minamide, T. Maeda, ‘Development of Copper Wire Bonding Application Technology,’ IEEE Transactions on Components, Hybrids and Manufacturing Technology, Vol. 13, No. 4, pp. 667–672, Dec. 1990. 60. J. Beleran, A. Turiano, D. R. Calpito, M. Stephan, D. Saraswati, F. Wulff, C. Breach, ‘Tail Pull Strength of Cu Wire on Gold and Silver-Plated Bonding Leads,’ Proc. Semi Technology Symposium, Semicon Singapore, 2005. (Kulicke and Soffa website http://www.kns.com). 61. C. L. Yeh, Y. S. Lai, ‘Transient Analysis of the Impact Stage of Wirebonding on Cu/ Low-K Wafers,’ Microelectronics Reliability, Vol. 45, No. 2, pp. 371–378, Feb. 2005. 62. D. Degryse, B. Vandevelde, E. Beyne, ‘Mechanical FEM Simulation of Bonding Process on Cu LowK Wafers,’ IEEE Transactions on Components and Packaging Technologies, Vol. 27, No. 4, pp. 643–650, Dec. 2004. 63. J. Tan, Z. W. Zhong, H. M. Ho, ‘Wire-Bonding Process Development for Low-k Materials,’ Microelectronic Engineering, Vol. 81, No. 1, pp. 75–82, Jul. 2005. 64. Website http://www.microbonds.com/ 65. R. J. Lyn, J. I. Persic, Y. K. Song, ‘Overview of X-Wire Insulated Bonding Wire Technology,’ Proc. IMAPS 39th International Symposium on Microelectronics, San Diego, USA, Oct. 2006.
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66. B. Chylak, S. Kumar, G. Perlberg, ‘Optimizing the Wire Bonding Process for 35-µm Ultra-Fine-Pitch Packages,’ Proc. Technical Program, Semicon Singapore, 2001. 67. B. Chylak, S. Tang, L. Smith, F. Keller, ‘Overcoming the Key Barriers in 35 µm Pitch Wire Bond Packaging: Probe, Mold and Substrates Solutions and Trade-Offs,’ Proc. IEEE/CPMT International Electronics Manufacturing Technology (IEMT) Symposium, pp. 177–182, 2002.
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9 Solid-state diffusion bonding A S H I R Z A D I, University of Cambridge, UK
The aim of this chapter is to provide a background to solid-state diffusion bonding of materials and/or components for which the conventional welding methods have proved unsuccessful. The basic aspects of solid-state diffusion bonding are explained in Section 9.1 and its advantages and limitations, in comparison with other welding/joining methods, are outlined in Section 9.2. Main bonding parameters and a typical laboratory bonding rig are described in Section 9.3. Section 9.4 contains a brief reference to the theoretical aspects of solid-state diffusion bonding, including a chronological table of the various modelling approaches made to date. The detrimental effects of surface oxide, as the major obstacle in solid-state diffusion bonding, and the methods used to overcome this problem are detailed in Section 9.5. The latest development in solid-state diffusion bonding, based on using liquid gallium to modify stable surface oxides, is described in Section 9.6. Finally a summary of solid-state diffusion bonding is given in Section 9.7.
9.1
Basic definition
Welding techniques are generally classified into two categories: liquidphase welding processes (e.g. arc welding) and solid-state welding/bonding processes (e.g. solid-state diffusion bonding and forge welding). In the former case, bonds are established by the formation and solidification of a liquid phase at the interface, while, in the latter case, the applied pressure has a key role in bringing together the surfaces to be joined within interatomic distances. Diffusion bonding, as a subdivision of both solidstate welding and liquid-phase welding, is a joining process wherein the principal mechanism is interdiffusion of atoms across the interface. However, only the basic aspects of solid-state diffusion bonding are discussed in this chapter. The International Institute of Welding (IIW) has adopted a modified definition of solid-state diffusion bonding, proposed by Kazakov.1 234 WPNL2204
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Diffusion bonding of materials in the solid-state is a process for making a monolithic joint through the formation of bonds at atomic level, as a result of closure of the mating surfaces due to the local plastic deformation at an elevated temperature which aids interdiffusion at the surface layers of the materials being joined. Although precise details are not known about the actual methods used by early blacksmiths and craftsmen, it is evident that solid-state welding has been used for more than a thousand years.2 Solid-state diffusion bonding is based on keeping the components to be joined at a high temperature and under a moderate interfacial pressure for a relatively long time. The bonding temperature depends on the material, and is normally between 50 and 70% of the melting point of the materials in Kelvin scale. The higher the bonding pressure, the shorter the bonding time; however, the pressure is limited by the amount of plastic deformation permissible in the components. Bonding times are much longer than those in ultrasonic and friction welding processes. Diffusion bonding can be used to join metals to ceramics and composites. Figure 9.1 shows high-precision aluminium components joined using solid-state diffusion bonding. The following terms are also used to refer to variants of the solid-state diffusion bonding process: • • • • •
diffusion welding isostatic bonding hot press bonding auto-vacuum welding and thermo-compression welding.
9.1 Aluminium-based microelectronic components joined using solidstate diffusion bonding and without any plastic deformation.
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9.2
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Solid-state bonding in comparison with other joining methods
Achieving high integrity joints with minimal detrimental effects on the parent material in the bond region and also the possibility of joining of dissimilar materials are the most promising features of solid-state diffusion bonding. Figure 9.2 shows the bond line of two dissimilar superalloys which were joined by solid-state diffusion bonding. The microstructures and compositions of the bonded superalloys, at and around the bond line, are identical to those of the parent alloys.3 Accordingly, interest in solid-state diffusion bonding has been growing in the last 40 years. The combination of diffusion bonding and superplastic forming for producing complex structures (e.g. honeycomb structures and hollow fan blades used in jet engines) has also played a part in increasing the use of this process as a commercial joining method. A number of advantages and limitations of solid-state diffusion bonding are listed below although not all apply to the diffusion bonding of all materials. Advantages: 1. The process has the ability to produce high quality joints so that neither metallurgical discontinuities nor porosity exist across the interface. An ideal diffusion bond is free from flaws, oxide inclusions, voids and loss of alloying elements.
50 µm
9.2 Optical micrograph of a dissimilar joint. The microstructures and compositions of the bonded Inconel 718 (lower piece) and a cobaltbase superalloy PWA647 (upper piece) are identical to those of the parent superalloys.3
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2. With properly controlled process variables, the joint can have strength and ductility equivalent to those of the parent material. 3. The joining of dissimilar materials with different thermo-physical characteristics, which is not possible by other processes, may be achieved by solid-state diffusion bonding. Metals, alloys, ceramics and powder metallurgy products have been joined by solid-state diffusion bonding. 4. Fairly high precision products can be manufactured without subsequent machining because of the solid-state nature of the process and relatively low bonding pressures required. This means that good dimensional tolerances for the products can be attained. 5. Solid-state diffusion bonding of components with intricate shapes or cross sections is possible. 6. Apart from the dimensions of the bonding machine, there is no limitation on the number of joints which can be made simultaneously. 7. Excluding the heating stage, the time required for bonding is independent of the size of bonded area. 8. Apart from the initial investment, the consumable costs of solid-state diffusion bonding are relatively low as no expensive solder, electrodes, or flux are required (although the capital costs and the costs associated with heating for relatively long times may be high). 9. Solid-state diffusion bonding is free from ultraviolet radiation and gas emission so there is no direct detrimental effect on the environment, and health and safety standards are maintained. 10. Automation is possible by computer controlling the basic variables of the solid-state diffusion bonding, e.g. time, temperature, and pressure. Limitations: 1. Great care is required in the surface preparation stage. Excessive oxidation or contamination of the faying surfaces will decrease the joint strength drastically. Solid-state diffusion bonding of materials with a stable oxide layer is very difficult (see Sections 9.5 and 9.6). Production of thoroughly flat surfaces and also precise fitting-up of the mating parts takes a longer time than with conventional welding processes. 2. The initial investment is fairly high and production of large components is limited by the size of the bonding equipment used. 3. Application of the force and heat in a vacuum or protective environment makes on-site working difficult. 4. The viability of this process for mass production is questionable, particularly because of the long bonding times involved. 5. The lack of reliable data about joint behaviour, e.g. fatigue life or fracture toughness, etc., and also inadequate methods for non-destructive testing
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the joint quality currently are problems, although it is expected they will be overcome in the future. Other solid-state joining processes such as ultrasonic welding, friction welding and pressure welding are also used for joining microelectronic and highprecision components. A comparative assessment of the bonding parameters for various solid-state bonding processes are given in Table 9.1. The data given in Table 9.1 must only be used as a guideline since bonding parameters vary substantially depending on the material(s) and joint geometry.
9.3
Main bonding parameters and apparatus
Solid-state diffusion bonding is normally carried out at high temperatures, about 50%–70% of the absolute melting point of the parent material. Longer times than those used in conventional pressure welding processes, e.g. roll or forge welding, are used in order to allow creep processes to contribute to bonding and to lead to a reduction in the pressure required for intimate contact between the faying surfaces. Thus, in contrast to most solid-state welding processes, solid-state diffusion bonding is not normally associated with large amounts of deformation. Therefore, it is considered a suitable method for bonding high-precision components, e.g. electronic components such as microwave guides and filters. Solid-state diffusion bonding of most metals is conducted under a vacuum of 10–3–10–5 mbar or in an inert atmosphere (normally dry nitrogen, argon or helium) in order to reduce detrimental oxidation of the faying surfaces. Bonding of metals which have oxide films that are unstable at the bonding temperature (e.g. silver) may be achieved in air (see Section 9.6 for recent developments in solid-state diffusion bonding of materials with stable surface oxide). Figure 9.3 shows a schematic diagram of a diffusion bonder including specimen set-up. A laboratory diffusion bonding rig and its peripherals are shown in Fig. 9.4.
Table 9.1 Comparative assessment of bonding parameters in various solid-state bonding processes Bonding process
Time
Temperature
Pressure
Deformation
Ultrasonic welding Friction welding Diffusion bonding Cold welding Forge welding
Short Short Longest Short Short
Low High High Lowest High
Low/high High Medium Highest High
Low/medium High Lowest High High
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Loading system including electric motor and gear box
Loading rod
Insulator RF induction unit
Induction coil Specimen
Thermocouple
Insulator
Load cell Diffusion pump
Rotary pump
9.3 Schematic diagram of a diffusion bonder including specimen set-up.
9.4
Theoretical aspects and modelling of solid-state diffusion bonding
The aim, when diffusion bonding, is to bring the surfaces of the two pieces being joined sufficiently close that interdiffusion can result in bond formation. In practice, because of inevitable surface roughness on an atomic scale, it is not possible to bring the surfaces of two pieces within interatomic distances by simple contact. Even highly polished surfaces come into contact only at their asperities and the ratio of contacting area to faying area is very low. Thus the mechanism of solid-state diffusion bonding can be classified into two main stages. During the first stage, the asperities on each of the faying surfaces deform plastically as the pressure is applied. These asperities arise from the grinding or polishing marks that have been produced in the surface finishing stage. The microplastic deformation proceeds until the applied stress at the contact area becomes less than the yield strength of the material at the bonding temperature. In fact, initial contact occurs between the oxide layers that cover the faying surfaces. As the deformation of asperities proceeds, more metal-to-metal contact is established because of local disruption of the oxide film which generally fractures readily. At the end of the first stage, the
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9.4 A laboratory diffusion bonding rig and its peripherals.
bonded area is less than 10% and a large volume of voids and oxide remains between the bonded areas. In the second stage of bonding, the thermally activated mechanisms lead to void shrinkage and this increases further the bonded areas. There are several hypotheses to explain how a bond is formed in the solidstate.1 The ‘Film Hypothesis’ emphasises the effect of surface film characteristics on the joining process. According to this hypothesis, the observed differences in weldability of various metals are attributed to the different properties of their surface films and all metals are assumed to bond if thoroughly cleaned surfaces are brought together within the range of interatomic forces. A different theory, the ‘Recrystallisation Hypothesis’, suggests that the strain hardening of the faying surfaces during plastic deformation causes the atoms to move to another site on the interface at high temperature. Subsequently, new grains grow at the interface and the bond is established. The ‘Electron Hypothesis’ is based on the formation of a stable electron configuration as a result of metallic bond formation. In the ‘Dislocation Hypothesis’, exposure
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of dislocations to the free surface, as a result of plastic deformation, breaks up the oxide film and produces steps on an atomic scale which enhance the seizure of the joining parts. Finally, the ‘Diffusion Hypothesis’, the most commonly accepted hypothesis, considers the contribution of interatomic diffusion during bond formation. The difference in the energy level of surface atoms and of bulk atoms is the basis of this hypothesis. A considerable amount of work on modelling solid-state diffusion bonding has been carried out.4 However, all attempts at modelling diffusion bonding have two main aims: 1. Optimisation of the process variables, e.g. bonding temperature, pressure and time so that the proper bonding conditions for a particular material can be identified. 2. Obtaining a reasonable and profound understanding of the mechanisms involved and their relative contributions not only for different bonding conditions but also for different materials being joined. Various models for solid-state diffusion bonding have been proposed to date. A chronological list of the proposed models is summarised in Table 9.2. However, in none of the existing analytical or numerical models has the effect of oxide film on the bond formation been considered.4,5 Hence, while some of the models can accurately predict the extent of bonding for metals with soluble oxides, they are of more limited use for alloys with stable oxides, such as aluminium alloys and nickel base superalloys.
9.5
Various approaches to overcome surface oxide problem
Previous experimental work on the solid-state diffusion bonding of alloys with a stable oxide layer has shown that the presence of the oxide film is the main barrier to successful bonding.5 Hence, the development of a method to disrupt the oxide layer should lead to significant improvements in bond integrity. Within the last 40 years of investigation into the solid-state diffusion bonding of the metallic materials with stable surface oxides, e.g. aluminium alloys, various approaches have been produced to overcome the oxide problem, and the major approaches are reviewed below.6
9.5.1
Imposing substantial plastic deformation
The surface oxide films normally have a much lower ductility than the parent alloy and hence they rupture when parent alloys are subjected to a large amount of plastic deformation, as is shown schematically in Fig. 9.5. Metal-to-metal contact is thus promoted as a consequence of local disruption of the oxide film on both faying surfaces; various ways by which this is
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Table 9.2 Chronological summary of solid-state diffusion bonding models5 1966 A two-stage model assuming localised plastic deformation followed by the diffusion-controlled process, using a recrystallisation model. 1967 Similar to the 1966 model except that the migration of interface away from the voids was assumed to occur during the second stage and the remaining isolated voids are removed by volume diffusion in the third stage. 1973 The first quantitative model, based on creep dominant mechanisms. 1975 Similar to the 1973 model but including diffusion mechanisms and superimposing two different wavelengths (corresponding waviness and roughness of the joining surfaces). Removal of the voids was modelled by using the sintering equations. 1982 Based on an intensive use of the sintering equations assuming six different diffusion mechanisms. 1984 Modification of the previous model in order to reduce the existing discontinuity in bonding rate between the second and final stage. 1984 A diffusive creep cavitation model adopted for fine grain superplastic alloys, based on creep model. 1987 Using the same mechanisms as the 1984 model but assuming two concomitant modes of void closure. 1989 Assuming elliptic voids with successive incremental changes in shape to circular in order to overcome some of the limitations of previous models. 1992 An overview of the void shrinkage models including the models based on diffusional flow of atoms, assuming successive mechanisms when interface diffusion is slower than surface diffusion or vice versa. 1995 Finite element analysis of interfacial contact process due to power law creep; no diffusional mechanisms for void shrinkage are taken into account.
achieved, including how the oxide fragments into ‘single’ and ‘double’ blocks have been investigated.7 Early work on solid-state diffusion bonding of aluminium alloys in 1966 and more recently in 1995, has shown that a minimum amount of deformation (~ 40%) is required to produce bonds with reasonable strengths.8,9 Although high strength bonds can be achieved by applying significant plastic deformation during the bonding process, this method has limited
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Continuous oxide film on both faying surfaces
40% deformation
Rupture of oxide film and metallic bond formation between oxide particles
Single oxide block
Double oxide block
9.5 Bond formation during solid-state diffusion bonding of a material with stable oxide layer made by imposing a substantial amount of plastic deformation. Single and double block oxide fragmentation is shown (thickness of oxide layer is exaggerated).7
application due to the need for substantial deformation of the parent materials during the bonding process. Clearly, there are similarities and overlap between this approach and conventional pressure/forge welding.
9.5.2
Enhancing microplastic deformation of the surface asperities
An alternative approach to overcoming the oxide problem in solid-state diffusion bonding is to use a fairly rough surface finish, which may lead to higher bond strengths than those obtained using polished surfaces. It is suggested that the local plastic deformation in the initial stage of the bonding process leads to rupture of the oxide film as the asperities deform and metallic contact is achieved (see Fig. 9.6).10 The rougher the surface, the greater is the plastic deformation of the asperities; therefore more oxide fracture will occur and consequently metal-to-metal bonding is improved. In early work, based on electric resistance measurements of Al–Al bonds, it was suggested that faying surfaces treated by coarse emery papers produced many more metallic bonds than smooth faying surface.11 Subsequently, the effect of surface roughness on the shear strength of Al-8090 bonds was studied and, despite considerable scatter in the results, it was concluded that higher bond strengths were achieved when a rougher surface preparation was adopted.12 These results are consistent with investigations into the influence of surface preparation on the bond strengths of Al-7475 bonds.13 More recently, the effects of the size and shape of surface asperities on the interfacial contact process have been modelled and verified experimentally.14,15 In contrast, low temperature solid-state bonding of copper, using different surface preparations, showed a significant reduction in the bond strength as
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Continuous oxide film on both faying surfaces
Oxide layer on surface asperities just before initial contact
Metallic bond formation between ruptured oxide layers due to microplastic deformation on the asperities
9.6 Schematic diagram of solid-state diffusion bonding of materials with stable oxide layers, where metallic bond formation initially occurs between the disrupted oxide layers at contacting surface asperities.10
the surface roughness increased.16 This seems to be inconsistent with the above hypothesis, namely that increased surface roughness improves the bonding behaviour of materials. This inconsistency is probably due to the different properties of the surface oxides of copper and aluminium; the soluble copper oxide layer is not as detrimental as insoluble aluminium oxide to the bonding process.
9.5.3
Presence of active alloying elements
The use of active alloying elements, such as magnesium and lithium, when diffusion bonding aluminium-based alloys, has also been investigated.17,18 According to this work, these active elements chemically interact with and break up the continuous and amorphous aluminium oxide layer at an interface to form an array of discrete particles. The work refers to the transmission electron microscopy results which clearly show the bond line of an aluminium 8090 alloy without magnesium or lithium contains a continuous oxide film. In contrast, the bond line of aluminium 8090 alloy containing magnesium and lithium consisted of discrete oxide particles with metallic bonding between them. The nature of the aluminium oxide at the interface also changes during the bonding cycle, from amorphous to crystalline.10 A good correlation between bond strength and the extent of broken oxide was observed, leading to the conclusion that the greater the content of these elements, the greater the disruption of the oxide layer and consequently higher bond strengths are achieved.17,18 It was also concluded that magnesium is more effective than
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lithium in increasing bond strength. The effects of active elements in either shims or interlayers, placed at bond lines, have also been studied.17 The use of Al–Mg interlayers improved the bond shear strengths by up to 50% in comparison with bonds made using pure aluminium interlayers.
9.5.4
Non-conventional solid-state bonding methods
Diffusion bonding of an aluminium alloy at a temperature above the solidus line but below the liquidus line, e.g. at 580 °C, where solid and liquid phases coexist, should aide the disruption of the oxide layers and promoted intimate contact at the interface.19 This method was (surprisingly) referred to as solidstate diffusion bonding despite the formation of a liquid phase. A maximum shear strength of 270 MPa was achieved when the volume fraction of the liquid phase was 2~3% and this was increased to 400 MPa by a post-bond heat-treatment. However, if the volume of the liquid phase exceeds 3%, grain boundary cracking occurs and a large amount of porosity forms on the bond line. Because of the high temperature used, the detrimental effects of melting on the microstructure and on the shape of the base material are expected to be substantial. In addition, any minor fluctuation during the bonding process would destroy the part being heated up to its solidus temperature. Hence, in practical terms, the process does not seem viable. A different approach combines solid-state diffusion bonding and friction welding to join Al-6061 and Al-6061/SiC MMC.20 In this method, a torsional force is exerted while an axial force was acting on the parts to be joined. One of the parts had a conical end in order to expel the worn oxide film from the interface. The important advantages of this method over conventional solidstate diffusion bonding and friction welding are the short bonding time (~5 minutes) and less plastic deformation, respectively. However, in contrast to diffusion bonding, the method is applicable only to parts with certain shapes (preferably with round cross-section) and which also have to be machined to provide conical ends. Also, unlike friction welding, which is carried out in air, this method required a vacuum.
9.6
Non-chemical surface oxide removal using liquid gallium
The formation of a stable surface oxide film, as the main obstacle in solidstate diffusion bonding most alloys, is virtually instantaneous for many metallic alloys which contain elements with a high affinity for oxygen, e.g. Ni, Cr, Al, Co, Ti, W. Although surface oxide films can easily be removed by grinding the faying surfaces of an alloy, a new oxide layer forms virtually immediately due to the exposure of the ground metallic surface to ambient oxygen. The fabrication of high strength bonds is possible when using a sophisticated
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bonding rig which is capable of in-situ ion sputter cleaning of the faying surfaces prior to diffusion bonding.5 Obviously the use of such a costly and complicated approach is limited to laboratory tests, or specific applications where manufacturing cost is not an overriding parameter. A new surface preparation method has recently been developed which is based on the non-chemical oxide removal or modification of a surface oxide film prior to the solid-state bonding process.21 This new method is capable of removing or modifying aluminium and/or chromium oxides from the faying surfaces of a wide range of alloys from nickel-base superalloys to aluminium and ferrous alloys.3 This can be done by grinding the surface of the alloy using an emery paper impregnated with liquid gallium, and this process typically takes less than one minute. Using this approach, formation of excellent solid-state bonds in superalloys have been reported despite the presence of the original stable oxides, such as chromium oxide, on the faying surfaces of the materials being bonded. In some cases, the microstructures and compositions of the joint interfaces are identical to those of the bulk alloys. Figure 9.7 shows the microstructure of a superalloy which was bonded using the above method.3 The high temperature tests carried out on dissimilar joints between Inconel 600 and Inconel 718, showed that the lifetime of the bonded sample was as long as the parent Inconel 600 (unbonded). These results demonstrate not only that the new joining method is capable of producing high strength joints but also proves that the use of a small amount of gallium does not compromise the high temperature properties of Inconel 600.22
50 µm
9.7 Optical micrographs of cobalt base superalloys bonded using the oxide removal method. The arrows show the approximate location of the bond line.
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Similarly, after treating aluminium faying surfaces in the same manner, substantial improvements in bond strength have been achieved. The strengths and ductilities of the tested aluminium bonds were as high as those of the parent alloy, as is shown in Fig. 9.8.23
9.7
Summary
Diffusion bonding is a promising method for joining materials for which conventional welding methods have proved unsuccessful. More than three decades of research on diffusion bonding has resulted in a comprehensive understanding of bond formation mechanisms. Solid-state diffusion bonding may result in high strength bonds if plastic deformation is large enough to provide metal–metal contact as a result of oxide disruption on the faying surfaces. The presence of some alloying elements such as magnesium, in either the base material or interlayer, may enhance disruption of the oxide layer, leading to higher bond strengths. Achieving high strength bonds when using a complicated bonding process (such as ion beam cleaning to remove the oxide film before bonding) has very limited applications. Similarly, introducing a high amount of plastic deformation to improve the bond strength confines the use of such approaches
Tensile test of Al 6082 samples
Parent alloy
Bond 1
Bond 2
9.8 Diffusion bonds in aluminium 6082 alloy have strength and ductility as high as those of the parent alloy.23 Note that both bonded samples failed away from the centreline joint.
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to the applications where massive deformations can be tolerated. The approaches which rely on the effect of alloying elements, in order to improve the bond quality, cannot be employed for a vast range of the materials which lack those particular elements. A new diffusion bonding method, based on using a small amount of liquid gallium in order to remove or modify surface oxide prior to bonding, seems very promising. Using this new approach, bonds in nickel superalloys and aluminium alloys with strengths as high as those of the parent alloys can be produced. Ultrasonic welding and cold pressure welding are the most commonly used solid-state bonding processes for the manufacture of microelectronic components. Friction welding has vast applications in manufacturing automotive components but it is gaining momentum as an alternative microjoining process. Diffusion bonding requires much longer bonding times but is capable of producing bonds with minimal deformation in the components being joined.
9.8
References
1. Kazakov N.F., Diffusion Bonding of Materials, 1985 English version Pergamon Press. 2. Singer C., Holmyard E.J., Hall A.R. and Williams T.I., A History of Technology, 1958, Oxford University Press. 3. Shirzadi A.A. and Wallach E.R., Science and Technology of Welding and Joining, 2004, 9, (1), 37–40. 4. Wallach E.R., Trans. JWRI, 1988, 17 (1), 135. 5. Shirzadi A.A., PhD thesis, University of Cambridge, UK, 1998. 6. Shirzadi A.A., Assadi H. and Wallach E.R., Journal of Surface and Interface Analysis, 2001, 31, (7), 609–618. 7. Cantalejos N.A. and Cusminsky G., J. Inst. Met., 1972, 100, 20. 8. Cline C.L., Welding Research Supplement, Nov. 1966, 481s. 9. Urena A., Gomez de Salazar J. M. and Escalera M. D., Metal Matrix Composites; Key Engineering Materials, Trans. Tech. Publications Ltd., Switzerland, 1995, 104– 107, 523. 10. Dray A.E., PhD. thesis, University of Cambridge, UK, 1985. 11. Enjo T., Ikeuchi K. and Akikawa N., Trans. JWRI, 1978, 7 (2), 97. 12. Ricks R.A., Mahon G.J., Parson N.C., Heinrich T. and Winkler P.J., Proc. Conf. Diffusion Bonding 2, Cranfield Institute of Technology, UK, March 1990, ed. Stephenson D.J., Elsevier, Amsterdam, 69. 13. Tensi H.M. and Wittmann M., Proc. Conf. Diffusion Bonding 2, Cranfield Institute of Technology, UK, March 1990, ed. Stephenson D.J., Elsevier, Amsterdam, 101. 14. Takahashi Y. and Tanimoto M., Trans. ASME, 1995, 117, 330. 15. Takahashi Y. and Tanimoto M., Trans. ASME, 1995, 117, 336 16. Nicholas N.H., Nichting R.A., Edwards G.R. and Olson D.L., Proc. Conf. Recent Trends in Welding Science & Technology, 1990, ed. David S.A. and Vitek V.M., ASM International, 547.
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17. Maddrell E.R., PhD thesis, University of Cambridge, UK, 1989. 18. Maddrell E.R., Ricks R.A. and Wallach E.R., Proc. Conf. Aluminium-Lithium 5, Williamsburg, Virginia, USA, March 1989, Materials and Component Engineering Publications Ltd, 451. 19. Enjo T. and Ikeuchi K., Trans. JWRI, 1984, 13 (2), 63. 20. Yokota T., Otsuka M., Haseyama T., Ueki T. and Tokisue H., Mater. Sci. Forum, 1997, 242, 225. 21. ‘Surface Treatment of Oxidising Materials’ USA Patent 6,669,534 B2, 30 December 2003; and UK Patent 2380491, 25 July 2002. 22. Shirzadi A.A., Department of Materials Science and Metallurgy, University of Cambridge, Unpublished work. 23. Shirzadi A.A. and Wallach E.R., Proceedings of International Conference on Designing of Interfacial Structures in Advanced Materials and their Joints, Osaka, 2002, Japan, 25–28 November 2002, 655–658.
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10 Bonding using nanoparticles A H I R O S E , Osaka University, Japan and K F K O B A YA S H I , Fukui University of Technology, Japan
10.1
Introduction
When particles are reduced in size to less than 10nm, the characteristics are different from those in the bulk state [1–4]. The melting point and the sintering temperature are detectably lower than the bulk state as one of their characteristics. It has been also known that small particles have the natural propensity to sinter or Ostwald ripen for reducing the total free energy. Such a sintering phenomenon is based on the driving force originated from the large surface energy of the nanoparticles. If this activated nature of the nanoparticles effectively affects the surface atoms of the bulk metal, a metalto-metal bonding using the nanoparticles as a filler material can be achieved at a significantly lower bonding temperature than that in the fusion welding or the diffusion bonding. To this end, we have proposed the novel bonding process using nanoparticles as a new application of nanotechnologies [5– 14]. This bonding process can be alternative to the current microsoldering using a high-temperature solder, such as Pb-10Sn or Pb-5Sn solder. The use of these solders including lead has been prohibited in recent years on account of worldwide environmental preservation. The nanoparticle used in our researches is the composite type Ag nanoparticle, which consists of the Ag metallo-organic nanoparticles and the Ag2CO3 particles (hereafter described as the composite Ag nanoparticles). The second section in this chapter describes the structure and thermal characteristic of the composite Ag nanoparticle. In the third section, the strength and detailed interfacial microstructure of the joints of various metals such as Au, Ag, Cu, Ni, Ti and Al using the composite Ag nanoparticle are mentioned, and the mechanism of bonding using the Ag nanoparticles is also discussed. In the fourth section, we state the effects of bonding conditions, such as bonding pressure, bonding temperature and holding time, upon the strength of the Cu-to-Cu joints in more detail, and reveal that the bonding process using the nanoparticle can be alternative to microsoldering using the Pb-rich solders. Finally, in the fifth section, bonding of a Si chip using the 250 WPNL2204
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composite Ag nanoparticle is presented as an application of the bonding process to electronics assembly.
10.2
Structure and thermal characteristics of composite Ag Nanoparticle
10.2.1 Structure The composite Ag nanoparticles are made by reducing Ag2CO3 particles in Myristyl alcohol [15]. The composite Ag nanoparticles consist of the Ag metallo-organic nanoparticles and the Ag2CO3 particles that are unreacted ones. Figure 10.1 shows a scanning electron microscope (SEM) image of the particle of the composite sectioned by a focused ion beam (FIB). The crosssection of the particle reveals that the core of Ag2CO3 having several microns in size is covered with the Ag metallo-organic nanoparticles. Figure 10.2 shows the image of the Ag metallo-organic nanoparticles obtained with a transmission electron microscope (TEM). The average particle size of the Ag metallo-organic nanoparticles is around 9.2 nm.
10.2.2 Thermal characteristics Figure 10.3 shows the results of thermal analyses with DTA and TG of the composite Ag nanoparticles in air and Ar atmosphere. The composite nanoparticles have two exothermal peaks in air and an exothermal peak in Ar atmosphere. In air, a rapid exothermal reaction with a large amount of weight reduction occurs at around 463K, followed by a broad exothermal reaction with small weight reduction at temperature range from 463K to 508K, and finally a rapid exothermal reaction without any weight change occurs at around 508K. In contrast to this, in Ar atmosphere, the reaction around 463K
Ag metallo-organic nanoparticles
Ag2CO3 particle
500 nm
10.1 SEM image of cross-sectional view of composite Ag nanoparticle.
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10.2 TEM image of Ag metallo-organic nanoparticles.
105
90
2
85 423
473 Temperature (K) (a)
DTA TG
100
Exo.
1
373
80 573
523
Endo.
90
423
473 Temperature (K) (b)
523
TG (mass%)
373
100 TG (mass%)
Exo.
DTA TG
95
Endo.
DTA (µV)
1
DTA (µV)
252
80 573
10.3 DTA and TG traces of composite Ag nanoparticles with heating rate of 10K/min in (a) air and (b) Ar atmosphere.
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Intensity (arbitrary unit)
takes place in the same way; however, no second exothermal reaction appears and weight reduction continues slowly over 573K. Figure 10.4 shows the results of X-ray diffraction (XRD) analyses of the composite nanoparticles heated in air up to temperatures before and after the exothermal peaks, which correspond to 453, 473 and 513K, respectively. The Ag2CO3 peaks disappear at 473K after the first exothermal reaction. That is to say, the decomposition of Ag2CO3 occurs at around 473K. Here the decomposition of Ag2CO3 alone is known to occur over 673K [16, 17]. It has been reported that Ag2O (crystal) is reduced at a lower temperature (433K), by adding ethylene glycol, than the decomposition temperature of Ag2O (crystal) itself (703K) [18]. The organic shell covering the Ag nanoparticles, therefore, may play a role similar to that of ethylene glycol. In addition, as a result of the observation using FE-SEM, the metallo-organic nanoparticles grow larger at 473K in air and somewhat grow even in the Ar atmosphere as shown in Figs 10.5 and 10.6. Thus, it is considered that the first exothermal reaction at around 463K is the reduction of Ag2CO3 and the growth of Ag nanoparticles accompanying the partial decomposition of the organic shell (Figs 10.5(b) and (c)). In the next step after the first peak in air, the rapid growth of the composite nanoparticles is recognized at around 508K before and after the second exothermal reaction followed by the slow weight reduction (Figs 10.5(d) and (e)). In contract to this, in Ar atmosphere, nano-size particles are recognized at 523K that temperature is higher than that of the second peak in air (Fig. 10.6(c)), and the sintering is insufficient even at 773K (Fig. 10.6(f)). From these results, it is suggested that the second peak at around 508K is an exothermal reaction caused by the rapid sintering of Ag nanoparticles after the full decomposition of the organic matters that adhered to surface of Ag. Consequently, the sintering process in Ar atmosphere is delayed by a restraint of the decomposition of the organic matters owing to lower oxygen partial pressure.
30
Ag Ag2CO3
513K
473K
453K 32
34 36 2θ (degree, Cu-Ka)
38
40
10.4 X-ray diffraction patterns of composite Ag nanoparticles heated up to various temperatures with heating rate of 10K/min in air.
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10.5 SEM images of composite Ag nanoparticles heated to various temperatures with heating rate of 10K/min in air, (a) 453K, (b) 473K, (c) 493K, (d) 513K, and (e) 573K. (a)
(b)
(c)
(d)
(e)
(f)
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10.6 SEM images of composite Ag nanoparticles heated to various temperatures with heating rate of 10K/min in Ar atmosphere, (a) 453K, (b) 473K, (c) 523K, (d) 573K, (e) 673K, and (f) 773K.
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In conclusion, the composite Ag nanoparticles realize the rapid first step sintering at lower temperature, which is lined with the rapid decomposition of the organic elements owing to the interaction between the oxidation of the organic shell and the reduction of Ag2CO3 that works as a resource of oxygen. This characteristic of the composite Ag nanoparticle is beneficial for applying it to a bonding material. In addition, the final sintering requires the heating beyond the second exothermal peak under certain oxygen partial pressure. Thus, the bonding conditions using the composite Ag nanoparticles are found to need a bonding temperature above 508K, and a certain bonding time and oxygen partial pressure.
10.3
Bonding of various metals using composite Ag nanoparticle
10.3.1 Bonding strength and fracture morphology Bonding of Au, Ag, Cu, Ni, Ti and Al using the composite Ag nanoparticle has been performed. The shape of the joint specimen is shown in Fig. 10.7. Here Au and Ag specimens are prepared by spatter deposition of Au and Ag on the Cu disc, respectively. The upper disc and lower disc are bonded together using the composite Ag nanoparticle paste. Figure 10.8 presents the shear strength of the joint of each metal. General views of the fracture
Upper disc
Diameter: 5 mm Thickness: 2 mm Diameter: 10 mm
Lower disc
Thickness: 5 mm
10.7 Shape of disc specimens for shear test of various metal joints.
Shear strength (MPa)
50 40 30 20 10 0
Al
Ti
Ni
Cu
Cu/Ag
Cu/Au
10.8 Shear strength of various metal joints using composite Ag nanoparticle.
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surface of each metal joint are shown in Fig. 10.9. Higher magnification images of Fig. 10.9 are shown in Fig. 10.10. The Al-to-Al and Ti-to-Ti joints have significantly lower shear strength than the other metal joints. Both joints are entirely fractured at the interface between the sintered Ag layer and the substrate, as shown in Figs 10.9(a) and (b). It is assumed that the bonding of Ag to each metal cannot be achieved in these joints because plastic deformation is not recognized on these fractured surfaces, as shown in Figs 10.10(a) and (b). On the other hand, the Ni-to-Ni joint has a somewhat higher shear strength compared to those of the Al and Ti joints. According to the images of the fracture surfaces as shown in Fig. 10.9(c) and Fig. 10.10(c), the fracture occurs both in the sintered Ag layer and at the interface between the sintered Ag layer and the Ni substrate. Although the trace of plastic deformation of Ag or Ni is not recognized, the sintered Ag is elongated and deformed on the fracture surface inside the Ag sintered layer. In contrast to this, the Cu-to-Cu, Cu/Ag-to-Cu/Ag and Cu/Au-to-Cu/Au joints are entirely fractured in the sintered Ag layer, as shown in Figs 10.9(d), (e) and (f), and the fractured Ag layer is elongated and deformed even in the interfacial fracture surface. The images are shown in Figs 10.10(d)–(f). These joints have almost the same shear strength. The results suggest that the bonding in
Lower side Upper side (b) Ti-to-Ti joint
Lower side Upper side (c) Ni-to-Ni joint
Upper side Lower side (d) Cu-to-Cu joint
Lower side Upper side (e) Cu/Ag-to-Cu/Ag joint
Lower side Upper side (f) Cu/Au-to-Cu/Au joint 1 mm
10.9 SEM images showing general views of fracture surfaces of various metal joints using composite Ag nanoparticle.
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Lower side
Upper side
Lower side
Upper side
Lower side
Upper side
(d) Cu-to-Cu joint
(c) Ni-to-Ni joint
Lower side
Upper side
(b) Ti-to-Ti joint
(a) Al-to-Al joint
Lower side
257
Upper side
Lower side
Upper side
(f) Cu/Au-to-Cu/Au joint
(e) Cu/Ag-to-Cu/Ag joint
10µ
10.10 Higher magnification images of Fig. 10.9.
these joints is achieved through a metallurgical bonding between the sintered Ag layer and the substrate. This has been revealed by observations of the interfacial region with TEM. For instance, the TEM image and corresponding lattice image at the interface between the sintered Ag layer and Cu in a Cuto-Cu joint as shown in Fig. 10.11 clearly present the metallurgical bonding at the interface.
10.3.2 Bonding mechanism An oxidation reaction contributes to the decomposition of the organic shell as previously mentioned. The decomposition of the oxide film is needed to achieve metallurgical bonding between the sintered Ag layer and each metal. In order to verify the removal mechanism of the oxide film on the substrate, the standard free energy value of the oxide formation for each metal is noted. In the case of the bonding to Cu, the surface oxide film which existed before bonding disappears and the metallurgical bonding is achieved as shown in Fig. 10.10. Based on the shear strength of the joints, the order of the bondability to each metal is as follows: Au, Ag > Cu > Ni > Ti > Al. This order is the
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Sintered Ag layer Ag (111)
Ag
Cu
Cu (111) 100 nm
Cu (a)
5 nm (b)
10.11 TEM images of interface between Cu and sintered Ag layer in Cu-to-Cu joint using composite Ag nanoparticles, (a) general view, and (b) lattice image.
identical to the order of the standard free energy value of the oxide formation for each metal at 573K as shown in Fig. 10.12 [19]. The reduction reactions by C of formulas (1) and (2) are shown in the dotted line in Fig. 10.12. C + O2 → CO2
(1)
2C + O2 → 2CO
(2)
The values of ∆G0 for reactions of formulas (1) and (2) exist between the values of ∆G0 of the oxide formation for Ti and Ni, and between those of Ni and Cu, respectively at 573K. Considerable differences between the shear strengths of these metal joints are recognized. Moreover, it is known as mentioned in the previous section that the surface oxide films of Cu and Ag are reduced during the bonding process. These results suggest that the organic shell or the organic matter generated by the decomposition of the organic shell of the nanoparticles has a similar reducing power to C, and it plays a role of reducing the oxide film on the metal surface. In the case of Cu-to-Cu bonding, the reduction reaction can be as formulated in (3) or (4). 2Cu2O + C [Organic shell or Organic matter] → 4Cu + CO2 (3) 2Cu2O + 2C [Organic shell or Organic matter] → 4Cu + 2CO (4) On the other hand, it cannot be expected that the reduction reaction occurs in Al and Ti joints, because their oxides are more stable than both CO and CO2. For this reason, whereas the joints of Cu, Ag and Au – the oxides of which are less stable and can be reduced by the organic shell – show the same strength level, the joints of Al and Ti – the oxides of which are more stable than carbon oxides – have extremely lower strength. In conclusion, the Ag metallo-organic nanoparticles contained in the composite Ag
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259
O = 2Ag 2
Standard free energy values of the oxide formation (∆G0)
0 = 2Cu 2O 4Cu + O 2 2NiO O2 = 2Ni + C + O2 = CO2
–400
2C + O
2
Si + O 2
= SiO 2
= TiO 2 l O3 Ti + O 2 2/3A 2 O2 = + l 4/3A
–800
–1200 273
= 2CO
573
873 1173 Temperature (K)
1473
1773
10.12 Standard free energy values of oxide formation for various metals.
nanoparticles have the function to reduce an oxide layer of a specific metal through the decomposition of the organic shell. By this function, the bonding process can be applicable to the metals such as Au, Ag and Cu of which oxides are less stable than CO and CO2.
10.4
Effects of bonding conditions on bondability of Cu-to-Cu joints
10.4.1 Effects of bonding conditions on joint strength During the bonding process using the composite Ag nanoparticles, both the sintering of the Ag nanoparticles and the bonding of the Ag nanoparticles to the metal substrate are simultaneously attained. The removal of the organic shell from the Ag nanoparticles leads out the activated property of the nanoparticles. It is considered that the increase in the bonding temperature and the holding time urges the removal of the organic shell, and the increase in the bonding pressure predominantly promotes the sintering and the bonding of the Ag nanoparticles after the decomposition of the organic shell. Figure 10.13 shows the shear strength of the Cu-to-Cu disc joints at each bonding temperature with various holding times and bonding pressures. The joint strength increases only a little with holding time in this bonding condition. The increase in the holding time may promote the sintering of the Ag particles. The influence of the bonding temperature depends on the bonding
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60 300s
600s
150s Shear strength (MPa)
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150s 50 40 30 20 10 0
1
2.5 5 Pressure (MPa) (a)
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300s
50
1
2.5 5 Pressure (MPa) (b)
7.5
Shear strength (MPa)
60 150s
300s
600s
50 40 30 20 10 0
1
2.5 5 Pressure (MPa) (c)
7.5
10.13 Shear strength of Cu-to-Cu joints made using composite Ag nanoparticle with varying bonding temperature, holding time and bonding pressures, (a) bonding temperature of 533K, (b) bonding temperature of 553K, and (c) 573K.
pressure. Under a lower bonding pressure of 2.5MPa or less, the increase of the bonding temperature considerably raises the joint strength, while under a higher bonding pressure of 5MPa or more, the joint strength is scarcely affected by the bonding temperature. Finally, the increase in the bonding pressure brings a significant improvement of the joint strength at each bonding temperature. This is probably because the enhancement of the bonding pressure makes the sintered Ag layer denser and improves the interfacial bonding strength between Ag and Cu.
10.4.2 Microstructure and fracture morphology Reflecting the results in the previous section, the difference in the microstructure and fracture morphology depending on the holding time is not recognized so much. A significant influence of the bonding pressure on the microstructure of the joints is seen as shown in Figs 10.14 and 10.15. Under the low pressure conditions as shown in Fig. 10.14, submicron voids are observed in
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Sintered Ag layer
Void
Cu
Cu
(b)
(a) Sintered Ag layer
Cu 5 µm
(c)
10.14 BE images of Cu/Ag interface in joints bonded at various bonding temperatures of (a) 533K, (b) 553K, and (c) 573 with bonding time of 300s under bonding pressure of 1MPa. Sintered Ag layer
Sintered Ag layer
Cu
Cu (b)
(a) Sintered Ag layer
Cu (c)
5 µm
10.15 BE images of Cu/Ag interface in joints bonded at various bonding temperatures of (a) 533K, (b) 553K, and (c) 573K with bonding time of 300s under bonding pressure of 5MPa.
the sintered Ag layer at each bonding temperature. However, under the high pressure conditions as shown in Fig. 10.15, the voids are reduced or disappear, and the sintered Ag layer becomes denser. Moreover, the change in the fracture mode is recognized as shown in Figs 10.16 and 10.17. The fracture
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(c) At 300°C
Upper side
Loading direction
Lower side
(a) At 260°C
1 mm
10.16 SEM images of fracture surfaces of Cu-to-Cu joints bonded at various bonding temperatures of (a) 533K, (b) 553K, and (c) 573K with bonding time of 300s under bonding pressure of 1MPa.
(a) At 280°C
(a) At 300°C
Upper side
Loading direction
Lower side
(a) At 260°C
1 mm
10.17 SEM images of fracture surfaces of Cu-to-Cu joints bonded at various bonding temperatures of (a) 533K, (b) 553K, and (c) 573K with bonding time of 300s under bonding pressure of 5MPa.
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path changes from the Ag/Cu interface to the sintered Ag layer as bonding pressure increases. These results show that both strengths of the Ag layer and the Ag/Cu interface are enhanced, and the Ag/Cu interfacial strength exceeds that of the sintered Ag layer under high pressure conditions. The influence of the bonding temperature is not apparently recognized in the microstructure observation of the sintered Ag layer. However, the fracture morphology is affected by the bonding temperature under a pressure of 1 MPa. Figure 10.18 shows the higher magnification images of region A in Fig. 10.16(b) and the EDX spectra of points A and B, respectively. Figure 10.19 shows the higher magnification images of region B in Fig. 10.16(b) and the EDX spectra of points C and D respectively. In region A, Ag is detected on the Cu side and the fractured Ag layer is elongated and deformed. Therefore, the fracture is considered to propagate within the sintered Ag layer around the Ag/Cu interface. However, Ag is not detected on the Cu side and the plastic deformation is not recognized on the fractured surface in 2500
(a)
Counts
2000
Ag
1500 1000
Point A
500 10 µm
0 0
250 500 750 Energy (keV) (b)
1000
2500
(c)
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Cu
1500 Cu
1000
Point B
500 Ag
10 µm
0 0
250 500 750 Energy (keV) (d)
Cu 1000
10.18 Detailed fracture surfaces of region A in Fig. 10.16(b) and EDX spectra, (a) SEM image of lower side (sintered Ag side) and EDX spectrum of point A, and (c) SEM image of upper side (Cu side) and EDX spectrum of point B.
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(a)
Counts
2000
Ag
1500 1000 500 Cu
Point C
Cu
10 µm
0 0
2500
(c)
250 500 750 Energy (keV) (b)
1000
Cu
Counts
2000 1500 1000
Cu
500
Point D 10 µm
O 0 0
Cu 250 500 750 Energy (keV) (d)
1000
10.19 Detailed fracture surfaces of region B in Fig. 10.16(b) and EDX spectra, (a) SEM image of lower side (sintered Ag side) and EDX spectrum of point C, and (c) SEM image of upper side (Cu side) and EDX spectrum of point D.
region B (Fig. 10.19). It is considered from the results that the Ag/Cu interfacial strength in this region is not sufficient. In addition, the Cu detected from the sintered Ag side (Fig. 10.19(c)) may be caused by Cu oxide. This suggests that the Cu oxide film remains and the metallurgical bonding of Ag to Cu is not achieved in this region. The area corresponding to region A declines in the fracture surface of the joint bonded at a lower temperature of 260°C as shown in Fig. 10.16(a). On the contrary, it increases in the joint bonded at a higher temperature of 573K as shown in Fig. 10.16(c). These results suggest that the Ag/Cu interfacial strength is enhanced by the increase of the bonding temperature, and this effect is more prominent under the low bonding pressure conditions. The increase in bonding temperature is considered to more effectively promote the decomposition of the organic shell than that of the holding time does. As a result, both bonding pressure and bonding temperature are dominant parameters for improvement of the bondability to Cu using the composite Ag nanoparticles.
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10.4.3 Comparison with lead-rich solder joints Figure 10.20 shows the contour maps for the shear strength that are constructed from the results of Fig. 10.13. The two types of dotted line in Fig. 10.20 show the strength of the Cu-to-Cu joints using the Pb-5Sn solder and the Pb10Sn solder, respectively. The joints using the composite Ag nanoparticles have higher strength than those using the current Pb-rich solders under wide bonding conditions. The bonding conditions that satisfy the lower basis of the joint strength for alternative to the Pb-5Sn solder, 18MPa, are quite wide ranging including a low bonding temperature and a low bonding pressure. The rise of the bonding temperature and the increase in the bonding pressure can meet the higher basis of the joint strength for the alternative to the Pb10Sn solder, 30MPa. The results reveal that the bonding process using the composite Ag nanoparticles can be an alternative to microsoldering using the Pb-rich soldered with regard to joint strength.
10.5
Bonding of Si chip
An attractive application of the bonding process using the composite Ag nanoparticles is die-attachment of semiconductor chip, which Pb-rich hightemperature solders are currently applied to. Unfortunately Si, which is a
Holding time (s)
300
150 1
2.5 5 Pressure (MPa) (a)
Holding time (s)
600
300
150
1
2.5 5 Pressure (MPa) (c)
7.5
600
300
150 1
7.5
Shear strength (MPa)
Holding time (s)
600
48 42 36 30 24 18 12 6
2.5 5 Pressure (MPa) (b)
7.5
Strength of Cu-to-Cu joint using Pb-10Sn solder (30MPa) Strength of Cu-to-Cu joint using Pb-5Sn solder (18MPa)
10.20 Contour maps for shear strength of joints using composite Ag nanoparticles with varying pressures and holding times at bonding temperatures of 533K, 553K, and 573K.
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dominant semiconductor material, is assumed to have poor bondability to the Ag nanoparticles because of its stable oxide film as shown in Fig. 10.12. Therefore, Au coating has been applied to the bonding surface of a Si chip. The Au coated 5mm square Si chip has been bonded to a Au plated 5mm diameter x25mm height Cu cylinder in the same way as the disc specimens mentioned above with bonding pressures of 1.0MPa and 2.5MPa at a bonding temperature of 573K for a holding time of 5 min. The thermal stability of the joints has been evaluated using the thermal cycle test of the temperature range of 233K to 398K for 500 cycles. Figure 10.21 shows the results of the tensile test of the joints before and after the thermal cycle test. Figure 10.22 shows the cross-sectional microstructure of the joint bonded with 1.0MPa and 2.5MPa before the thermal cycle test. The tensile strength of the joints 40
Tensile strength (MPa)
35
Before thermal cycle test After thermal cycle test
30 25 20 15 10 5 0 1.0MPa
2.5MPa
10.21 Shear strength of joints bonded at 573K with bonding pressure of 1.0MPa and 2.5MPa using composite Ag nanoparticle before and after thermal cycle test.
(a)
(b) Si
Sintered Ag layer
Sintered Ag layer
Au
Au
Cu
Cu 10µm
10.22 BE images of cross-sectional microstructure of Si chip-to-Cu joints using composite Ag nanoparticle bonded at 573K with bonding pressure of (a) 1.0MPa, and (b) 2.5MPa before thermal cycle test.
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is around 25MPa in both conditions before the thermal cycle test. In addition, the sintered Ag layers of the joints are dense without any cracks or debonding as shown in Fig. 10.22. Here, the fracture of the Si chip-to-Cu/Au specimens entirely occur inside the Si chip. Therefore, the sintered Ag layer and the interface between the Si chip and the sintered Ag layer probably have a sufficient bonding strength more than 25MPa. However, after the thermal cycle test, the joint strength changes depending on the bonding pressure. Although the joint strength deteriorates in the case of a bonding pressure of 1.0MPa after the thermal cycle test, it is maintained in the case of a bonding pressure of 2.5MPa (Fig. 10.21). From this result, it is suggested that the joint with the bonding pressure of 1.0MPa is probably damaged by thermal stress. Though, the appropriate bonding pressure of 2.5MPa can realize the dense and reliable sintered Ag layer and sufficient bonding at the interface after the thermal cycle test. Moreover, we have reported that a Ag coated Si chip is successfully bonded to a Ag coated Al substrate using the composite Ag nanoparticles and the joint endures a thermal cycle test with temperature range from 213K to 573K [13]. Thereby, this bonding process using the composite Ag nanoparticles is applicable to the bonding of the Si chip.
10.6
Conclusion and future trends
This chapter has described the recent achievements on the bonding process using the nanoparticles that the authors’ research group has developed. This novel bonding process can realize metallurgical bonding at a low temperature around 573K without melting or long range diffusion. Since the bonding layer consisting of metallic Ag has a good heat-resistance, and excellent electrical and thermal conductivity, the bonding process has a great potential to be a reliable assembly process for electronics devices. However, there are some problems to be overcome in putting the bonding process to practical use. These include reducing bonding pressure and temperature to avoid the damage of devices. Moreover, there is a fear of the ionic migration of Ag that causes a short-circuit. Thus, the evaluations and countermeasures are necessary in the practical applications. Aside from this, development of nanoparticles other than Ag, such as a Cu metallo-organic nanoparticle, is also considered. These are our future research subjects.
10.7 1. 2. 3. 4. 5.
References
A. N. Goldstem, C. M. Esher and A. P. Alivisatos, Science, 256 (1992), p. 1425. P. Pawlow, Z. Phys. Chem., 11 (1909), p. 609. M. Takagi, J. Phys. Soc. Japan., 9 (1954), p. 359. N. Gladkich, R. Niedermayer and K. Spiegel, Phys. Status Solidi, 15 (1966), p. 181. E. Ide, S. Angata, A. Hirose, and K. F. Kobayashi, Acta Materialia, 53 (2005), p. 2385.
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6. E. Ide, S. Angata, A. Hirose, and K. F. Kobayashi, Mater. Sci. Forum, 512 (2005), p. 383. 7. E. Ide, A. Hirose and K. F. Kobayashi, J. Soc. Mater. Sci. Japan., 54 (2005), p. 999. 8. E. Ide, A. Hirose and K. F. Kobayashi, Mater. Trans., 47 (2006), p. 211. 9. E. Ide, S. Angata, A. Hirose and K. F. Kobayashi, Proc. Inter. Conf. New Frontiers of Process Science and Engineering in Advanced Materials Part 2, High Temperature Society Japan, (2004), p. 233. 10. S. Kobayashi, E. Ide, S. Angata, A. Hirose and K. F. Kobayashi, Proc. ASME InterPACK 2005, ASME, (2005), IPACK2005-73161 on No.1729CD. 11. S. Angata, E. Ide, S. Kobayashi, A. Hirose and K. F. Kobayashi, Proc. ASME InterPACK 2005, ASME (2005), IPACK2005-73164 on No.1729CD. 12. A. Hirose, E. Ide, S. Angata, S. Kobayashi and K. F. Kobayashi, Proc. Inter. Conf. Electronics Packaging 2006, IMAPS Japan, (2006), p. 59. 13. T. Yamaguchi, E. Ide, S. Kobayashi, H. Imaeda, A. Hirose, K. F. Kobayashi, M. Yamagiwa and Y. Murakami, Smart Processing Technology, Vol. 1, High Temperature Society Japan, (2006), p. 187. 14. T. Yamaguchi, E. Ide, S. Kobayashi, A. Hirose, K. F. Kobayashi, J. G. Lee and H. Mori, Proc. 12th Sympo. Microjoining and Assembly Technol. for Electronics, Vol. 12, Japan Weld. Society, (2006), p. 371. 15. Patent, WO 2005/075132 (2005). 16. P. Norby, R. Dinnebier and A. N. Fitch, Inorganic Chemistry, 41 (2002), p. 3628. 17. Y. Sawada and N. Watanabe, Thermochemica Acta, 138 (1989), p. 257. 18. A. Ono, N. Kondo, M. Kurosawa, M. Ohyama, S. Kodate, K. Okamoto and M. Ito, Fujikura Technology Review, 79 (2004), p. 107. 19. L. S. Darken and R. W. Gurry, Physical Chemistry of Metal, McGraw-Hill (1953), p. 342.
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11 Diffusion soldering and brazing M L K U N T Z and Y Z H O U , University of Waterloo, Canada
11.1
Introduction to diffusion soldering/brazing
Diffusion soldering/brazing is an attractive microjoining process for difficultto-join materials when high joint quality is required. This process is similar to conventional soldering and brazing in that a low melting point interlayer forms a liquid at the faying surface; however, it differs by employing longer process times to homogenize the joint. Diffusion soldering/brazing is a diffusion controlled process where the bond forms through a process known as isothermal solidification. This chapter will outline the mechanics of joint formation in the diffusion soldering/brazing process, typical applications, conventional modeling approaches for simulation of the process, and metallurgical or mechanical considerations.
11.1.1 Historical development Examples of similar processes to diffusion brazing can be dated as early as 2500 BC [1]. In modern usage, the process has been referred to by a great number of different names. Perhaps one of the most commonly used terms is the copyrighted ‘transient liquid phase (TLP) bonding’ [2]. Other common variations include ‘eutectic brazing’ [3], ‘eutectic bonding’ [4, 5], ‘activated diffusion bonding’ [6], ‘solid–liquid interdiffusion bonding’ [7], ‘liquid interface diffusion bonding’ [8], and ‘transient insert metal bonding’ [9]. To reduce the confusion surrounding the naming convention of the process, the American Welding Society (AWS) has attempted to resolve the issue by applying the name ‘diffusion brazing’ (DFB, or DB) to include all of the closely related variations [10]. In conventional brazing and soldering processes, the only distinction between the two is the melting point of the interlayer. The difference being that a braze has a melting point above 450°C, while a solder melts below 450°C. In diffusion brazing/soldering, the same convention generally applies, with diffusion brazing occurring at higher temperatures and diffusion soldering 269 WPNL2204
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occurring at lower temperatures. Other subtle variations differentiate the two processes. For example, in some diffusion brazing applications, the formation of intermetallic phases at the interface is avoided because it is considered detrimental to joint properties; however, in diffusion soldering the formation of intermetallic phases can be inevitable, and in some cases, is desirable because the melting temperature of these phases can be very high. Most microjoining applications are performed at lower temperatures and, therefore, diffusion soldering is regularly associated with small-scale applications while diffusion brazing is associated with large-scale applications. In practice, though, there are several high temperature diffusion brazing applications for microjoining. In this chapter, diffusion brazing will be used to refer to all processes that rely on isothermal solidification of a low melting point interlayer to form a homogenous joint. In cases where a low process temperature is used, diffusion soldering may be used interchangeably with diffusion brazing.
11.1.2 Applications The diffusion brazing process has seen widespread application in conventional ‘large-scale’ processes; however, in recent years development has shifted towards microjoining. Typical examples are found in the aerospace or land turbine industries, where diffusion brazing has been used in the manufacture and repair of turbine blades and components [11] with materials such as Nibased superalloys [12] and single crystal [13]. Specific examples of applying diffusion brazing to microjoining applications are relatively limited; however, there has been longstanding interest. In one case, Bernstein and Bartholomew [14] produced experimental bonds on electrical components in the ternary Ag-In-Sn system. Recently, however, there has been renewed significance in application of diffusion brazing technology to miniature components. For example, the need for high quality Pb-free joints in microelectronic interconnects has driven research in this area. Conventional Pb-free solders exhibit poor joint properties and employ high process temperatures which can cause component damage. The use of solders designed for diffusion soldering allows for lower processing temperatures and produces good joint properties. Additionally, diffusion brazing is a candidate process for joining biocompatible materials for medical implant devices where no braze/solder can remain at the interface. Several other applications can be found in the literature, including: metal matrix composites (MMC) [5, 15], and shape memory alloys (SMA) [16, 17], or intermetallic structures. In microelectronics applications, the use of Sn as an interlayer is common in diffusion soldering of interconnections. For example, Au-Sn [18,19], AgSn [20], and Cu-Sn [21–23] materials systems have been used in general applications, since Sn additions to solders decreases the melting point of Au, Ag, and Cu, which are commonly used for electrical connections. When
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materials are especially sensitive to temperature, additions of Bi or In have successfully been used to employ diffusion soldering at very low temperatures. In/Sn interlayers have been used to join Cu substrates [24, 25], and Au/In interlayers have been used to join Ti metallization layers on Si [26]. Lee et al. [27] used Sn/Bi coated Sn-3.5Ag interlayers with Cu base metal (Fig. 11.1). Furthermore, powder-based solder pastes containing a mixture of Bi or Sn/Bi powders with a non-melting Cu powder have been developed for diffusion soldering applications (Fig. 11.2) [28, 29]. For high temperature diffusion brazing applications, Cu coated Ti powders [30], Sn/Sb [31], and Cu/Ni [32] solder pastes have been used. The diffusion brazing process is well suited to joining materials that are difficult to join using conventional techniques. The process produces bonds that are nearly homogenous, with mechanical properties similar to the base metal. A unique characteristic of this process is the shifting melting point. The post-bond microstructure at the interface has a melting point that is similar to the base metal, and above the processing temperature. This enables multiple diffusion brazing operations to be completed subsequently without re-melting of the interlayer, or use of a lower process temperature. For this reason, the process is well suited to applications where the component is exposed to elevated service temperatures. In situations where conventional solders or brazes would re-melt, the homogeneous nature of the diffusion brazed joint prevents failure.
Ag3Sn
Solder
IMC UBM
11.1 Joint interface between a solder ball and UBM (Au/Ni/Cu), reflow temperature of 220 °C. A Sn-Bi coating melts at low temperatures and assists wetting of the solder ball on the metallization [27].
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(a)
100 µm
(b)
11.2 Diffusion soldering using a composite Sn-10%Sb solder paste with a soldering temperature of 245 °C and held for (a) 5 minutes, (b) 15 minutes. The melting point depressant Sb (light phase) is consumed through isothermal solidification with increasing hold time [31].
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Process description of diffusion soldering/ brazing
Simply, the process of diffusion brazing is the same as a traditional brazing operation with the addition of an extended hold time at the process temperature. The components are assembled at room temperature with a low-meltingpoint interlayer placed at the joint interface. The assembly is heated to the pre-determined process temperature, causing the interlayer to melt and wet the faying surfaces at the interface. The joint assembly is then held isothermally at the process temperature. During the hold, the width of the liquid is reduced through the mechanism of isothermal solidification. When the entire liquid width has been consumed, the assembly can be cooled to an intermediate temperature for a homogenizing heat treatment, if necessary. The assembly is then cooled to room temperature and the process is complete. As the name implies, diffusion brazing is a diffusion controlled process, and as a result, the well known fundamental equations for mass transfer derived by Fick can be used to describe it. The process can generally be considered a two-phase moving interface problem. After heating, the joint is a sandwich of two solid substrates with a liquid layer between the faying surfaces. In a non-steady state condition, where concentration gradients exist and resulting diffusion occurs, the direction and rate of interface motion is governed by a mass balance: ( C L – CS ) ⋅ d X ( t ) = DS ⋅ ∂ CS – DL ⋅ ∂ C L dt ∂x ∂x
(1)
where CL and CS are the liquid and solid concentrations, respectively, at the solid/liquid interface; and, DL and DS are the solute diffusivities in the liquid and solid, respectively. After melting of the interlayer, the direction of interface motion depends on the local concentration gradients in the liquid and solid phases. The mechanism of bond formation in diffusion brazing can be described by considering four discrete stages [33–36]. The generally accepted classification of these stages is: 1. 2. 3. 4.
heating dissolution and widening isothermal solidification homogenization.
During heating, the parts to be joined are heated to just below the melting temperature of the interlayer. The dissolution and widening stage begins at the onset of interlayer melting or eutectic reaction with the base metal. Additional heating results in dissolution of the base metal and a widening of the liquid phase. The isothermal solidification stage begins when the maximum
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liquid width has been reached. Solute diffusion across the solid/liquid interface results in boundary migration towards the joint centerline through isothermal solidification. The stage is complete when the liquid phase is entirely consumed. A solute peak remains at the joint centerline; this peak can be reduced through a homogenization heat treatment at an elevated temperature to avoid the formation of harmful phases upon cooling.
11.2.1 Process advantages and disadvantages The major advantages of diffusion brazing include: •
•
•
• •
•
Diffusion brazing often results in a homogeneous joint microstructure, and, in many cases the location of the joint is difficult to discern. This is of particular benefit when similar material properties across a joint are required. For example, there is no thermal expansion mismatch or galvanic corrosion potential due to variations in composition across the joint. Mechanical properties in the joint can approach those of the base metal. Hardness, tensile strength, and stress-rupture performance have been found to be similar to base metal properties. Near 100% joint efficiency can be attained in some applications. The melting temperature of the joint after the diffusion brazing operation is similar, or higher to the melting point of the base metal. Additional diffusion brazing processing can be carried out at the same temperature without the risk of re-melting the prior joints. Additionally, the component can be operated at high service temperatures without resulting in component failure. Components with very different geometries can be joined, especially joints with a high joint sectional area. Materials that cannot be joined with conventional fusion welding techniques can be joined without solidification cracking or formation of brittle intermetallic phases, and without distortion and residual stresses that can result from localized heat input. Excessive faying surface preparation is not necessarily required. Diffusion brazing is tolerant of a small amount of surface contamination such as oxides or impurities, especially if an appropriate fluxing treatment, or atmosphere are used.
The disadvantages of diffusion brazing are: •
Process times for diffusion brazing can be significantly longer than those required for conventional brazing operations. This can be a problem when short cycle times are required for process efficiency. Furthermore, in precipitation hardened materials, there may be a loss of base metal strength, or harmful equilibrium phases can precipitate and degrade mechanical properties.
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Relatively high process temperatures can be required for diffusion brazing operations. These temperatures typically affect the entire component, which can often result in damage of thermally sensitive materials. Close dimensional tolerance is required to ensure good part fit-up. Excessive variations in gap across the joint can cause localized lack of bonding where liquid has been squeezed out to areas with a larger gap.
11.2.2 Procedures for diffusion brazing The parts to be joined are prepared before bonding by assembling the components in the desired configuration. The faying surfaces of the substrates should be clean of oxides, scale, oil and contaminants before assembly. Chemical and mechanical cleaning methods can be used. Acceptable surface treatments include machining, grinding, sanding, and sandblasting. Common degreasing techniques may be required to clean the faying surface. Fluxes are also used to clean the interfaces in-situ during heating, and protect the joint from oxidation during bonding. The faying surfaces should be prepared such that a close tolerance is maintained across the interface. The surface condition must also be sufficiently smooth to avoid porosity and lack of bonding. The interlayer is placed at the interface where it will wet the faying surfaces upon melting, or be drawn into the joint through capillary action. In some cases, especially when the liquid spreads aggressively, a method of containing the liquid is required. If the liquid spreads away from the joint, there will be insufficient liquid to adequately fill the gap, resulting in an excessively porous joint. To prevent this from happening, a ceramic ‘stopoff’ is typically painted around the periphery of the joint. The interlayer can be placed at the interface by a number of methods. The interlayer contains a melting point depressant (MPD) solute and is usually added in the form of a thin foil. The MPD can also be added in a powder form, with or without a composite preform [37], as a coated powder [38] or as a coating (electroplating, sputter or thermal spray) [39]. If a simple eutectic system is considered, it can be shown that the bulk composition of the interlayer can be tailored to melt at the eutectic temperature, or can shift in-situ, through reaction with the base metal to form liquid.
11.2.3 Equipment for diffusion brazing Heat is usually applied in a radiation, convection, or induction furnace; a controlled atmosphere or vacuum is used in most applications. Depending on the geometry, pressure may be applied by fixturing, adding dead weight, or just through the weight of the components themselves. The addition of pressure has no effect on the rate of isothermal solidification, but can affect
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the process kinetics by squeezing liquid out of the joint, reducing the liquid width. If the pressure is too high, enough liquid could be squeezed out of the joint such that there is a lack of bonding at the interface. Fixturing should be used, when necessary, to maintain the required gap tolerance, typically within the range of a thousandth of an inch or less. An inert atmosphere or vacuum is usually used to protect the joint from oxidation at elevated temperatures.
11.2.4 Process parameters The process parameters for diffusion brazing have an impact on joint quality and process duration. The initial process conditions include the interlayer thickness, or width (Wo), and the interlayer composition (CF). The initial set up of a diffusion brazing joint is shown in Fig. 11.3. The processing variables include the pressure, heating rate, brazing temperature, isothermal hold time, and homogenizing heat treatment. •
•
Interlayer thickness: the initial thickness of the interlayer has a profound effect on the duration of the isothermal solidification stage. The time required for isothermal solidification increases exponentially with interlayer width. A sufficiently thick interlayer is required to wet the faying surfaces and fill any irregularities in the joint gap; however, the thinnest possible interlayer should be used to minimize the process time. For diffusion brazing operations, foil interlayers around 25 µm thick are typically employed. In diffusion soldering, much thinner interlayers are applied, usually less than 5 µm thick, using a deposition technique. Interlayer composition: like the thickness, an important consideration is the selection of the interlayer composition. The composition affects the time required for dissolution of the interlayer, as well as the maximum width of the liquid. For a binary system, the maximum liquid (Wmax) width after dissolution of the base material is given by:
Wmax =
C F ⋅ Wo C Lα
(2)
where CLα is the composition of the liquid at the brazing temperature. An interlayer with the eutectic composition will dissolve instantaneously at the eutectic temperature, and the extent of widening will depend on the brazing temperature. A pure interlayer will dissolve base material as part of the eutectic reaction, and additional widening with heating to the brazing temperature will occur, resulting in a significantly larger Wmax compared to a eutectic interlayer with the same initial thickness. This relationship between interlayer composition and liquid width will have a similar effect on process duration as Wo, increasing the isothermal solidification time with increased Wmax. The interlayer composition can
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X (t)
∆X 2
Stage 4: Homogenization
Stage 3: Isothermal solidification
CαL
X (0)
X (ts ) = 4
Joint schematic
Wmax 2
Composition profile
11.3 Diffusion brazing process mechanics, including initial conditions, stage 1: heating, stage 2: dissolution and widening, stage 3: isothermal solidification, and stage 4: homogenization. The thermal cycle is shown with a phase diagram schematic and corresponding joint schematic with composition profile [redrawn from. [57]
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•
•
•
•
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also affect the process kinetics of isothermal solidification in multicomponent systems with two or more diffusants. Pressure: there is no effect of pressure on the process kinetics; however, increased pressure can squeeze liquid out of the joint, decreasing the liquid width and shortening the overall process time. Excessive pressure can affect joint quality by squeezing too much liquid from the joint, resulting in porosity and lack of bonding. Fixturing can be used to maintain an adequate joint gap when the weight of the components results in excessive pressure at the interface; however, fixing the substrates at a set distance can result in Kirkendall porosity in some cases. Heating rate: the heating rate from room temperature to the melting point can affect joint quality. A sufficient heating rate is required to prevent excessive interdiffusion with the interlayer before melting occurs. This can be a problem for thin interlayers where loss of solute can reduce the liquid width below adequate levels. Brazing temperature: the isothermal hold temperature will affect the duration of the process. Increasing the process temperature will increase the solute diffusivity, and influence the kinetics of isothermal solidification. However, increasing the process temperature can also increase the maximum liquid width, resulting in an increase in the process duration. Optimization of the brazing temperature should be carried out with the consideration that some materials may be sensitive to higher temperatures, resulting in damage or microstructural changes. Isothermal hold time: the isothermal hold time is usually selected to ensure that isothermal solidification has been completed before the onset of cooling. This ensures a homogeneous joint microstructure without the solidification structure common to conventional brazing processes. Homogenizing heat treatment: in situations where a homogenization is required to avoid the formation of harmful phases in the joint region, the time required is a function of the solute peak remaining after isothermal solidification, the maximum solute level that can be tolerated, and the diffusivity of the solute, which is temperature dependent.
11.3
Diffusion soldering/brazing process mechanics
11.3.1 The four stages of diffusion brazing Stage 1: Heating The first stage in the diffusion brazing process is the period of heating from room temperature to just below the onset of interlayer melting, shown as point 1 in Fig. 11.3. As the assembly is heated, some solid state diffusion between the interlayer and base material occurs. The magnitude of the interaction depends on a number of factors including surface condition (i.e.
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roughness, cleanliness, etc.), and pressure exerted normal to the interface (higher pressure will help to flatten the asperities and improve surface contact). The exchange results in a typical diffusion couple compositional profile as shown in Fig. 11.3. The extent of diffusion is expected to be small in most cases; however, the amount of mass transferred during the heating stage will also depend upon the eutectic temperature and the heating rate as well as the diffusivities of the elements [34]. The heating rate during this stage is an extremely important parameter, especially when a very thin interlayer is used. A sufficiently high heating rate is required to limit interdiffusion between the interlayer and substrates. If the heating rate is too low, excessive diffusion of the low melting point constituent away from the interface can reduce the amount of liquid that forms to the point that there is not enough to fill the interface gap. This problem will be most severe with very thin interlayers and low solute concentrations. Li et al. [40] found that decreasing the heating rate from 5 to 1 K/s resulted in a need to increase the interlayer thickness from 0.6 to 2 µm to achieve acceptable bonds in alumina metal matrix composites using a copper interlayer. Clearly, it is important to control the heating rate during the heating stage of diffusion brazing. Dissolution and widening stage As the assembly is heated past the melting temperature, the interlayer will melt and wet the base metal at the faying surface. Depending on the material system and the composition of the interlayer, dissolution can occur upon the onset of the melting point, or in-situ through a eutectic reaction. Figure 11.3 shows the condition where the interlayer composition is eutectic, and Fig. 11.4 shows the case of a pure interlayer. With an increase in temperature above point 1, the equilibrium composition of both the liquid and the solid at the interface will track along the solidus and liquidus phase boundaries, respectively. To maintain equilibrium at the solid/liquid interface, the base metal is dissolved by the liquid resulting in a widening of the liquid zone due to conservation of mass. In the case of a pure interlayer, eutectic melting will initiate at the base metal/interlayer interface where, through diffusion, there will be a thin band that is at the eutectic composition [13]. The liquid will grow from each interface through dissolution of both the interlayer and the base metal. Since diffusion in the liquid phase is relatively fast (orders of magnitude higher than in the solid state), dissolution of the base metal occurs rapidly [41]. The liquid width is at a maximum (Wmax) at the peak, or brazing temperature where the liquid composition is at CLα as shown in Fig. 11.3. It is important to note that the time at which the maximum liquid width is reached does not necessarily coincide with the time at which the bonding temperature is reached. The kinetics of dissolution depend on a number of factors including solute diffusivity and heating rate such that widening of the
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2b 1
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Time
Liquid
CαL CLα
Interlayer
1 2a
Stage 2b: Widening
CαL CLα
Wmax
Liquid
2b Joint schematic
Composition profile
11.4 Variation of the dissolution and widening stage for a pure interlayer. In-situ melting occurs during heating through a eutectic reaction with the base metal. The interlayer is rapidly dissolved until a homogeneous liquid phase is reached. Isothermal solidification then proceeds in a similar manner to that shown in Fig. 10.3.
liquid will continue some time after the bonding temperature is reached until a steady state is achieved [42]. The time required to reach the maximum liquid width will depend on the initial width (Wo) and composition (CF) of the interlayer: a thick, pure interlayer will require the longest time; and conversely the dissolution of a eutectic interlayer will be instantaneous. Using a 80 µm thick pure Cu foil for diffusion brazing of pure Ag, Tuah-Poku et al. [35] found that widening of the liquid zone required times in the order of minutes. Clearly, under some circumstances the widening stage can be significant in terms of process time. Isothermal solidification stage After the dissolution and widening stage, an isothermal hold period ensues during which diffusion of the MPD solute across the solid/liquid interface
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into the base material occurs. This diffusion process is akin to a bi-phase diffusion couple. To maintain equilibrium on the phase diagram at a fixed temperature, the liquid composition at the interface is constant at CLα. Likewise, the composition of the solid at the interface is fixed at CαL. Furthermore, it is suggested in the literature that the composition of the liquid can be assumed uniform across the entire liquid width [41]. Since diffusion in the liquid phase is orders of magnitude faster than diffusion in the solid phase and the endpoints of the liquid phase are fixed and equal, this is considered an accurate assumption. The direction and rate of interfacial motion must satisfy a mass balance at each solid/liquid interface such that each of the two solid/ liquid interfaces move toward the joint centreline. This interfacial motion is shown schematically in Fig. 11.3. The mechanism of advancing interface motion is epitaxial growth of the solid phase into the liquid. This process has been coined ‘isothermal solidification’. The isothermal solidification stage is complete when the two interfaces meet at the centreline of the joint and there is no liquid remaining. The rate of interfacial motion will depend on the diffusivity of the solute in the base material, the miscibility gap between CαL and CLα, and the concentration gradient of the solute in the base material. The length of time required for completion of the isothermal solidification stage will also be greatly dependent upon the maximum width of the liquid. Grain boundary grooving and enhanced grain boundary diffusion has been cited by a number of researchers in order to explain faster than expected isothermal solidification rates [35, 42, 43]. Liquid penetration, or grain boundary grooving has been observed in essentially all studies of diffusion brazing with polycrystalline materials [43–45]. Tuah-Poku et al. [35] pointed out that liquid penetration at the grain boundaries and the accompanying departure from a planar solid/liquid interface made measurement of the liquid width very difficult, in some areas isothermal solidification was complete but there was still liquid remaining at pockets where grain boundary grooving had occurred. The extent of grooving as well as rate of isothermal solidification was found to be higher in fine grained samples versus coarse grained samples [44]; however, the difference in isothermal solidification rate between coarse grained samples and single crystal base metal were indiscernible [46]. Zhou and North [43] found that when the isothermal hold temperature approaches the melting point of the base metal (i.e. T ≥ 0.75Tm), the contribution of grain boundary diffusion on the total amount diffused becomes less significant, even more so with increasing grain size. The increase in interface kinetics with finer grain size has been attributed to enhanced grain boundary diffusion and grain boundary grooving. Homogenization stage When the isothermal solidification stage is complete, there will be a peak of solute at the solidus composition (CαL) remaining at the joint centerline,
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shown by the profile in Fig. 11.3. The assembly is held at an elevated temperature so the magnitude of the solute peak is gradually decreased over time through diffusion. This homogenization of the joint continues until an acceptable level of solute remains. What level is acceptable depends on the material and the application; however, it should be below the concentration at which harmful phases will precipitate in the solid state during cooling. Precipitates may degrade the mechanical properties of the joint. The length of time required for an adequate homogenization treatment depends on the magnitude of the solute peak, and the diffusivity of the solute. The temperature can be reduced during homogenization to shorten exposure at high temperature, or as part of an extended heat-treating procedure; however, this will also reduce the solute diffusivity. In the case where isothermal solidification of the liquid directly to an intermetallic phase occurs, as is common in many low temperature diffusion soldering applications, the homogenization stage may be necessary for favourable intermetallics to grow. It may even be possible to eliminate the intermetallic phase completely through growth of the primary phase, and thus achieve a homogeneous microstructure.
11.3.2 Critical stages in diffusion brazing The two most important stages in terms of joint quality are the isothermal solidification and the homogenization stages; coincidently it is also these stages that require the longest time for completion. In systems with a low solubility limit the isothermal solidification stage becomes longer and is more important. It is essential that the stage is not terminated by cooling before all of the liquid has solidified or a cast microstructure and solute rejection will result in segregated phases, which can degrade mechanical properties. Conversely, in some systems with a high solubility limit, the homogenization stage requires a longer time for completion and is considered more important. Premature termination of the homogenization stage can result in an inhomogeneous microstructure due to precipitation of additional phases that will adversely affect mechanical properties. Since the isothermal solidification and homogenization stages depend on diffusion of the solute in the solid base metal, they are orders of magnitude longer than the heating or dissolution stages. Thus, it is appropriate that most of the attention given to diffusion brazing parameters focuses on the time required for completion of these stages.
11.3.3 Isothermal solidification in multi-component systems Typical engineering materials are not pure, but contain a wide range of alloying additions. Diffusion brazing interlayers are often tailored to match
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the base metal chemistry, with the addition of one or more MPD constituents. However, dissimilar materials can also be used to tailor the melting range of the interlayer, for example the use of a Cu/In interlayer to diffusion solder Au [14]. The addition of a second solute to the interlayer can complicate the diffusion brazing process significantly. The problem can be formulated by considering a pure semi-infinite base metal substrate with a binary eutectic interlayer. Upon heating, the interlayer will melt and dissolution will occur. It is assumed that the dissolution will occur on a straight line from the initial interlayer composition to the pure base metal composition as shown in Fig. 11.5. Thus, the initial liquidus composition will lie at the intersection of this line and the liquid phase boundary on the Gibbs’ isotherm for the bonding temperature. As in the isothermal solidification of binary systems, the mechanism of the process in ternary systems is the diffusion of both solutes across the solid/liquid interface and into the base material. A mass balance can be written for each of the solutes in the system. From Equation 11.1, the mass balance for each solute can be written: ( C L1α – Cα1 L ) ⋅ d X ( t ) = D1 ⋅ ∂ C1( x , t ) x = X ( t ) dt ∂x
(3)
( C L2α – Cα2L ) ⋅ d X ( t ) = D2 ⋅ ∂ C2 ( x , t ) x = X ( t ) dt ∂x
(4)
Solute 1
L Final liquid to solidify Dissolution path
Initial CF interlayer composition
ts Base metal composition
α
Initial liquid
to
Solute 2
Co Solidus
Liquidus
11.5 Isothermal solidification mechanics for a multi-component system assuming a pure base metal and a binary interlayer. Dissolution is assumed to follow a straight line from the initial composition to the base metal composition. Isothermal solidification then proceeds via the shifting tie line mechanism until final solidification is reached.
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Inspection of the mass balances shows that under most circumstances, the interface velocity predicted by each will be different; however, there can only be one rate at which the interface migrates into the liquid. Thus, there must be some mechanism controlling the isothermal solidification behavior. This can be explained by Gibbs’ phase rule: f = n – p, where n is the number of components, p is the number of phases, and f is the degrees of freedom of the system. In the binary case there are zero degrees of freedom; however, in the ternary case there is an additional degree of freedom which allows isothermal solidification to proceed. This mechanism occurs by way of a shifting liquid composition. Interface motion is governed in such a way that the mass balance is maintained by allowing the composition of the liquid phase to continuously change as the process progresses. Thus, the composition of the last liquid to solidify could be significantly different than the initial liquid composition, and the consequence is that the rate of isothermal solidification will change continuously.
11.3.4 Wide-gap process variation In some situations, the largest joint gap that can be tolerated is so great that the time required for isothermal solidification exceeds that which is feasible. To shorten the process time, a composite interlayer can be used. The interlayer, usually in the form of a powder compact, contains the melting point depressant plus an additive that is designed not to melt at the bonding temperature. The higher melting point constituent does not melt entirely; however, it does participate in dissolution and is often similar in composition to the base material. The function of the additive is first, to reduce the amount of liquid that is necessary to fill the joint gap; and second, to increase the interfacial area between the liquid and solid phases and effectively increase the diffusion rate. This is shown schematically in Fig. 11.6. Gale and Butts [47] point out that a proper composite ratio is essential; if excessive liquid is formed, the high melting point additive will be completely dissolved. Conversely, if there is insufficient liquid to fill the gap, the joint will be excessively porous.
11.3.5 Temperature-gradient TLP bonding A novel diffusion brazing method for joining Al-based materials where a temperature gradient is imposed across the joint to increase the isothermal solidification kinetics was developed by Shirzadi and Wallach [48]. Since the equilibrium composition at the solid/liquid interfaces is temperature dependent, Shirzadi and Wallach propose that by simply imposing a temperature gradient across the liquid zone, a compositional gradient is induced in the liquid. Since diffusion in the liquid is relatively fast, solute transport across
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Base metal
Non-melting constituent
Liquid
Base metal
11.6 A joint schematic of the wide-gap variation of diffusion soldering/brazing. A composite interlayer containing a non-melting constituent is placed at the interface. The non-melting constituent often has a composition matching the base metal.
the liquid width occurs rapidly. The solute diffuses from the high concentration (low temperature) interface to the low concentration (high temperature) interface. This results in solidification at the low temperature interface since removal of the solute atoms locally increases the equilibrium solidification temperature. Thus, the liquid zone migrates in the direction of increasing temperature leaving behind a solute build up at the solidus concentration, as shown in Fig. 11.7. By means of a mass balance, the solute that is removed from the liquid by solidification will result in a shrinking liquid and eventual isothermal solidification when the two interfaces come together. The isothermal solidification time required in TG-TLP bonding is typically lower than what is expected in traditional diffusion brazing [49]; however, if a homogenization treatment is required the process time will generally increase.
11.4
Evaluating joint properties
The integrity of diffusion brazed joints can only be verified using destructive testing methods. Non-destructive inspection techniques are generally not viable for diffusion brazed joints. Thus, there is no way to ensure that the joint is defect free. Fortunately, diffusion brazing is a stable process, and with proper selection of process parameters, a high quality joint can be produced with high certainty. For procedure qualification and quality control, a combination of metallurgical examination and mechanical testing is required.
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Temperature
Th Tc
C
Hot interface Liquid Cold interface Solute diffusion
Solidified metal Joint schematic
Composition profile
11.7 Process schematic for the temperature-gradient TLP bonding process. A temperature gradient is imposed across the joint, causing the liquid zone to migrate towards the cold interface and shrink as it leaves behind a solute-rich solid.
11.4.1 Measuring diffusion brazing process kinetics Experimental determination of diffusion brazing process kinetics is typically accomplished through manual observation of the amount of liquid remaining after an isothermal hold period. The isothermal solidification stage is interrupted by premature cooling, resulting in athermal solidification of remaining liquid. The width of the solidification microstructure, typically eutectic, is then measured from a metallurgical cross-section of the joint. Difficulty in measuring the width of the solidified structure is usually encountered due to irregularity of the interface, as shown in Fig. 11.8 [11]. Grain boundary grooving and cellular protrusions result in a non-planar morphology, and the liquid width varies across the joint. Solutions to this are to take an average of measurements [46, 50], or alternatively, to measure the planimetric area of the solidified phase and divide by the observed length to determine the average width [9]. Additional problems with accurately measuring the liquid width using metallographic observation are caused by solidification of a primary phase during cooling from the brazing temperature to the eutectic temperature. The epitaxial nature of primary solidification results in a similar microstructure that is difficult to discern from the isothermally solidified material. Thus, when the width of eutectic is measured, it does not include a fraction of liquid that solidifies athermally, resulting in an underestimation of the actual
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W1
W3 W2
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Eutectic solidification structure
11.8 Quantifying isothermal solidification kinetics using visual inspection of a solidification structure after premature cooling. The typical non-planar interface makes an accurate measurement difficult. An average width can be obtained using several measurements, or by dividing the planimetric area over the length of the joint; however, it has been shown that this does not include a fraction of liquid that has solidified to a primary solid phase during cooling to the eutectic temperature.
width of liquid remaining [9, 42, 51]. The fraction of liquid that solidifies to the primary phase depends on the material system, the cooling rate, and the temperature difference between the brazing temperature and the eutectic. Some researchers have corrected the liquid width for this effect by applying the lever rule [52], and others have used a Scheil simulation [53]. Typically, a proportional correction is required to obtain the actual liquid width when measured using metallographic inspection. Recently, thermal analysis tools have been applied to measure the process kinetics of isothermal solidification. Differential scanning calorimetry (DSC) has been used to measure the amount of liquid remaining in-situ [54]. This method has been applied to both low temperature diffusion soldering applications using a powder based paste for an interlayer [55], as well as to higher temperature diffusion brazing applications with a foil preform [56]. Results have shown that DSC is a powerful method for accurately characterizing the solid/liquid interface motion during the isothermal solidification stage [57]. The benefits of this laboratory scale analytical technique are that the entire solid/liquid interface is characterized during the measurement, eliminating the need for costly metallurgical preparation of samples and the problems with measuring non-planar interfaces. Furthermore, DSC can be
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applied to composite interlayers or solder pastes with a non-melting constituent, where measurement of a liquid phase is difficult using manual inspection. There are modeling techniques available for simulation of the diffusion soldering/brazing process, and these methods are useful for lowering the costs of process parameter optimization, for example selection of isothermal hold temperature and time. There are several limitations, however, that reduce the accuracy of these models. While the predictions for isothermal solidification process kinetics are generally accurate for simple metallurgical systems [57], especially in higher temperature diffusion brazing; the accuracy decreases for low temperature applications and multi-component systems where several intermetallic solid phases can exist at the interface. In the latter case, a experimental approach is still the best approach. An example of experimentally observed process kinetics in diffusion brazed Ag with a Cu interlayer is shown in Fig. 11.9. This is a simple system since there are no intermetallics formed between Ag and Cu, as a result, the isothermal solidification rate can be accurately predicted using an analytical solution.
11.4.2 Metallography of diffusion brazing Metallurgical cross-sections of diffusion brazed joints can be examined for defects and microstructure. Conventional metallographic techniques are generally used, including optical and scanning electron microscopy. When further analysis, 1
W /Wmax
0.8
0.6
0.4
0.2
Measured liquid width Analytical prediction ξ = –7.56 µm/hr1/2
0 0
0.04 0.08 Normalized time (hr1/2/µm)
0.12
11.9 A comparison of measured and predicted isothermal solidification rates. The liquid width was accurately measured by a method using differential scanning calorimetry. The predicted interface rate constant of ξ = –7.56 µm/hr1/2 agrees very well with the measured isothermal solidification kinetics. The base metal was pure Ag, with a Cu interlayer, and a brazing temperature of 800 °C [57].
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such as phase identification is required, analytical techniques can be used, including energy dispersive X-ray spectroscopy and X-ray diffraction. The location of the interface after diffusion brazing may be difficult to determine when the joint quality is good. A homogeneous, defect-free interface is an indication of good mechanical properties. Defects that can be observed during visual inspection include porosity and lack of bonding. Observation of eutectic phase or solidification structure at the interface is an indication that the isothermal solidification stage was terminated before all of the liquid was consumed. Either a continuous layer of eutectic will form along the interface, or isolated pockets of eutectic can be observed. The eutectic solidification structure is an indication of incorrect process parameters and can usually be resolved by increasing the isothermal hold period. Figure 11.10 shows the progression of isothermal solidification in a diffusion brazed Ag joint using a eutectic Ag-Cu interlayer at 800°C. The joint microstructure was observed after 1, 2, 3, and 4 hours of isothermal hold time at the brazing temperature. The gradual decrease in width of the solidified eutectic phase can be easily observed; however, an exact measure of the liquid width over the entire joint is difficult using visual inspection alone. Additional phases and intermetallic compounds can be observed near the interface region. These phases may be brittle and have an adverse effect on mechanical properties. The formation of these phases is an indication of a high solute concentration at the interface during cooling. This can possibly be alleviated through the use of an appropriate homogenization heat treatment immediately after isothermal solidification to allow solute to diffuse away from the joint area. In some cases, the entire assembly will be solutionized and artificially aged after diffusion brazing and it may be possible to include this as part of the diffusion brazing thermal cycle.
11.4.3 Evaluating joint quality Selection of a mechanical testing protocol depends on the application. Generally, tensile testing is used to evaluate the tensile strength of the joint. In some cases, stress rupture testing may be used to evaluate high temperature creep performance. In diffusion soldering applications, shear testing is commonly used to evaluate the joint strength. The joint efficiency is a measure of the joint performance relative to the base metal properties. In diffusion brazing applications it is possible to achieve near base metal properties; however, failure will generally occur along the bond line.
11.5
Modeling of diffusion soldering/brazing
In the simplest of cases, the simple two component system (i.e. binary interlayer), the position of the solid/liquid interface during isothermal
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30 µm (a)
30 µm (b)
11.10 Progression of isothermal solidification during diffusion brazing of Ag using a 25 µm thick Ag-Cu eutectic interlayer at 800 °C (a) isothermal hold time of 1 hour, (b) 2 hours, (c) 3 hours and (d) 4 hours. The non-planar interface makes accurate measurement of the liquid width difficult. After 2 hours there is a discontinuous liquid phase, and after 4 hours isothermal solidification is complete. A cellular precipitation of Cu is visible in the base metal microstructure, necessitating further homogenization if a homogeneous microstructure is desired [57].
solidification can be shown to be proportional to the square root of the hold time [11]. To optimize the process parameters, such as the isothermal hold temperature and time, it is necessary to accurately measure the process kinetics of isothermal solidification, specifically the rate of interface motion. This can be accomplished through experimental means; however, it is often beneficial to use mathematical modeling tools to simulate the process to
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30 µm (c)
30 µm (d)
11.10 (Continued)
reduce the tedious and costly experimental procedure. Simulations are a useful tool for improving the process efficiency by minimizing the time required for isothermal solidification. Two approaches are available, simple analytical solutions and more complicated, but more accurate numerical techniques. In both cases, obtaining accurate results requires good knowledge of the input parameters. For example, accurate thermodynamic data, such as diffusivity and phase boundaries is required.
11.5.1 Analytical solution for isothermal solidification Simple analytical solutions have been derived for isothermal solidification in simple binary systems [11, 34, 35]. Several forms of these solutions have been used to calculate the instantaneous width of the liquid phase, and the
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time required for the entire liquid width to be consumed [58–60]. The solutions for isothermal solidification are formulated by considering a bi-phase diffusion couple, which is governed by the mass balance in Equation 11.5. If equilibrium is established at the solid/liquid interface, the compositions of the solid and liquid, at the interface, are fixed by a tie-line throughout isothermal solidification. Furthermore, diffusion in liquid is orders of magnitude faster than in solid, so it can be assumed that there is no concentration gradient in the liquid. A simplified form of the mass balance can then be rewritten as Equation 11.6. ( C L – CS ) ⋅ d X ( t ) = DS ⋅ ∂ CS – DL ⋅ ∂ C L dt ∂x ∂x
(5)
( C Lα – Cα L ) ⋅ d X ( t ) = D ⋅ ∂ C( x, t ) x = X ( t ) dt ∂x
(6)
In order to solve the mass balance at the solid/liquid interface, given by Equation 11.6, a substitution parameter, λ, can be applied, where: λ= x t
(7)
and, the initial condition of a pure base metal is given by C(∞) = 0. The solid/liquid interface (λi) is then defined as ξ, where the boundary condition of C(ξ) = (CαL) is applied. λi = ξ
(8)
Thus, the instantaneous position of the interface can be found by Equation 11.9, and the velocity is given by Equation 11.10. X(t ) = ξ ⋅
(9)
t
d X(t ) = 1 ξ dt 2 t
(10)
Fick’s second law for the rate of concentration change in a control volume, in terms of λ, can be written in one dimension as: 2 (11) – λ ⋅ dC = D ⋅ d C 2 dλ dλ 2 The solution to the differential is obtained by integration and application of the initial and boundary conditions. This gives the concentration profile in the solid as a function of position and time:
C ( λ ) = Cα L
erfc 2⋅
λ
D ⋅ ξ erfc 2 ⋅ D
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The mass balance, rewritten as a function of the substitution parameters is given by Equation 11.13, where k is the partition coefficient, CLα/CαL. Cα L ⋅ ( k – 1) ⋅
ξ = D ⋅ dC 2 dλ λ = ξ
(13)
Substitution of Equation 11.12 into the mass balance gives a solution for the solid/liquid interface, shown by Equation 11.14. No closed form solution can be found, so the equation must be solved using iterative techniques.
ξ = –2 ⋅ ( k – 1) –1 ⋅
2 exp – ξ 4 ⋅ D D ⋅ π ξ erfc 2 D
(14)
The isothermal solidification stage is complete when the solid/liquid interface reaches the joint centerline, X(t) = –Wmax/2. The time required for this to occur (ts) can be found by rearranging Equation 11.10: ts =
– Wmax 2⋅ξ
2
(15)
Inspection of Equation 11.14 shows that under isothermal conditions ξ is a constant. This is referred to as the isothermal solidification rate constant because it is an indication of the solid/liquid interface velocity. Increasing ξ results in faster solid/liquid interface motion, and shorter duration of the isothermal solidification stage. Furthermore, ξ is independent of the initial liquid width, thus it is useful to discuss process kinetics in terms of ξ rather than the time required for isothermal solidification when the temperature and initial liquid width are varied. The assumptions used in derivation of the analytical models can be summarized by [50]: 1. One-dimensional diffusion: solute diffusion in the base metal is assumed to be in one direction only, i.e. in a direction perpendicular to the base metal. Solute flux is a function of concentration gradient. 2. Quiescent liquid: there is no mixing in the liquid. Solute redistribution is a function of diffusional motion only. This assumption is generally considered valid unless induction heating is used for a heat source. 3. Constant diffusivities: the coefficient of diffusion is constant over the isothermal hold time and independent of solute concentration. In fact, diffusivity is related to the local chemical composition; however, the assumption is required to simplify the solution. 4. Semi-infinite base metal: the base metal can be considered semi-infinite so long as the diffusion distance of the solute (√Dt) is much less than the thickness.
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5. Equilibrium at the solid/liquid interface: the conditions at the interface are assumed to obey local equilibrium and follow the phase boundaries on the equilibrium phase diagram [61]. 6. Constant solid/liquid interface area: the interface is assumed to maintain a planar profile. It will be shown that grain boundary grooving has been shown to occur during widening and isothermal solidification and the effect on process kinetics will be discussed. A number of other authors have reported that experimentally measured solid/ liquid interface kinetics are found to be much faster than predicted using analytical solutions. In many cases, the discrepancy has been attributed to experimental setup. For example, if the liquid interlayer is squeezed or spreads away from the joint, there will be a reduction in the width of the liquid. There are other situations; however, where the assumptions used to derive the analytical solutions do not apply. For example, when the base metal substrate cannot be assumed semi-infinite compared to the interlayer thickness, as is common in many microjoining applications. In this situation it is possible that solute build up in the base metal will slow down the process kinetics and extend the time required for isothermal solidification. A numerical simulation approach would improve the accuracy in this case.
11.5.2 Numerical modeling techniques The finite difference method has been applied to model the kinetics of the diffusion brazing process. One advantage of the numerical approach is the ability to model the discrete dissolution, isothermal solidification and homogenization stages as one continuous process as shown by Zhou and North [62]. The one-dimensional model has evolved with additional development for more complex systems. Zhou and North [43] extended the model into two dimensions in order to account for grain boundary effects. Takahashi et al. [63] modeled the system with a saw-tooth profile for the solid/liquid interface in order to characterize the effects on a non-planar boundary on process kinetics. In comparison of one-dimensional numerical and analytical solutions for the isothermal solidification stage, there is no significant difference in the interface rate constant. The major difference in the results is in the dissolution and widening stage for which no complete analytical solution exists due to the complexity of variable temperature [64]. In this sense, if the prevailing conditions are such that the dissolution time and width are minimal, i.e. low interlayer solute content and small interlayer width, then the time required for the dissolution stage as a fraction of the total process time approaches zero and can be considered negligible. Zhou et al. [34] point out that the interface kinetics during the isothermal solidification stage can be estimated
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using an analytical solution so long as the effects of grain boundaries can be neglected. During dissolution, however, no single interface rate constant can be found with an analytical solution, and numerical methods are required to characterize solid/liquid interface motion. Modeling in multi-component systems For many simple applications, especially when the interlayer chemistry is tailored to match the base material with the addition of a MPD, the analytical solution for isothermal solidification will give a reasonable prediction of the process kinetics. In many engineering applications, however, the effects of a multi-component interlayer on the isothermal solidification rate must be considered. For a ternary system, there is one degree of freedom; hence, there must be a rule controlling the system that can be used to predict isothermal solidification behavior. The constant, ξ, is the rate constant that describes the interface movement. If the cross-diffusional effects are assumed negligible, then the rate constant can be found for each solute in the liquid through a mass balance [58]: Cα L ⋅ ( k i – 1) ⋅
ξi dC i = Diα ⋅ 2 dλ λ = ξ i
(16)
where ki is the ratio of liquidus to solidus equilibrium concentrations for the ith component, and Diα is the diffusion coefficient in the solid for the ith solute. The solution for the rate constant is given in Equation 11.17, which must be solved iteratively.
ξ i = –2( k i – 1) –1 ⋅
– ξ 2i exp 4 ⋅ Diα Diα ⋅ π ξi erfc α 2 Di
(17)
By examination, all of the terms in Equation 11.17 are constant for a given temperature. This means that unless all of the solutes have a certain combination of Diα and k i, ξ1 and ξ2 will be different [65]. There is no way that the interface can sustain a different rate of solidification for each solute. The only way to maintain a single velocity of the solid/liquid interface is for the solid and liquid compositions to track along the phase boundary lines. Numerical techniques have been applied to predict isothermal solidification process kinetics in ternary systems [46, 50], for example, finite difference equations for the mass balance of each solute diffusant can be written (i.e. Equation 11.18) then solved simultaneously to find the interface position using an iterative procedure. The shift in liquid composition is calculated
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using the phase boundaries, which can be approximated using thermodynamic predictions if available. i Dii ⋅ ∂C · ∆t = C Li (t + ∆t) · X(t + ∆t) ∂x
+ CSi (t + ∆t) · [X(t + ∆t) – X(t)] – C Li (t) · X(t) (18)
11.6
Summary and future trends
With the trend towards miniaturization of components and continually smaller devices, diffusion soldering/brazing will continue to be a candidate process that is well suited to these microjoining applications. Very thin joint thicknesses and small joint sections are possible with diffusion soldering/brazing processes, which is ideal for microjoining processes. Furthermore, the high joint quality that can be achieved in materials that are traditionally difficult to join makes diffusion soldering/brazing an attractive process. Diffusion soldering can be carried out at very low temperatures with the added benefit of a higher remelt temperature, which is ideal for temperature-sensitive materials and applications. Development of new interlayers, application techniques, and modeling approaches will increase the widespread application of diffusion soldering/brazing for emerging materials in applications for microelectronics, medical implant devices, shape memory alloys, and composites. Diffusion soldering/brazing can also be considered as a potential technique for nanojoining applications.
11.7 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14.
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47. W.F. Gale and D.A. Butts, Sci. Technol. Weld. Joi., 9, 4, 283–300 (2004). 48. A.A. Shirzadi and E.R. Wallach, US Patent No. 6,257,481. 49. A.A. Shirzadi and E.R. Wallach, Science and Technology of Welding and Joining, 2, 3, 89–94 (1997). 50. J.E. Ramirez and S. Liu, Welding Journal, 365s–375s (1992). 51. A. Sakamoto, C. Fujiwara, T. Hattori and S. Sakai, Weld. J., 68, 63–71 (1989). 52. Y. Nakao, K. Nishimoto, K. Shinozaki and C.Y. Kang, Joining of Advanced Materials (eds T.H. North), pp. 129–144, Chapman and Hall (1990). 53. K. Ohsasa, T. Shinmura and T. Narita, J. Phase Equilibria, 20/3, 199–206 (1999). 54. S.F. Corbin and P. Lucier, Metall. Mater. Trans. A, 32A, 4, 971–978 (2001). 55. S.F. Corbin and D.J. McIsaac, Mater. Sci. Eng. A (Switzerland), 346, 132–140 (2003). 56. M.L. Kuntz, S.F. Corbin and Y. Zhou, Acta Mater., 53, 10, 3071–3082 (2005). 57. M.L. Kuntz, Y. Zhou and S.F. Corbin, Metall. Mater. Trans. A, 37A, 8, 2493–2504 (2006). 58. C.W. Sinclair, J. Phase Equilibria, 20, 4, 361–369 (1999). 59. Y. Zhou, J. Mater. Sci. Lett., 20, 841–844 (2001). 60. S. R. Cain, J.R. Wilcox and R. Venkatraman, Acta Mater., 45, 2, 701–707 (1997). 61. J.S. Langer and R.F. Sekerka, Acta Metall., 23, 1225–1237 (1975). 62. Y. Zhou and T.H. North, Model. Simul. Mater. Sci. Eng., 1, 505–516 (1993). 63. Y. Takahashi, K. Morimoto and K. Inoue, Trans. JWRI, 30, 535–541 (2001). 64. Y. Zhou and T.H. North, Z. Metallkd., 85, 775–780 (1994). 65. C.W. Sinclair, G.R. Purdy and J.E. Morral, Metall. Trans. A., 31A, 1187–1192 (2001).
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12 Laser soldering Y H T I A N and C Q W A N G , Harbin Institute of Technology, P.R. China
12.1
Introduction
Laser soldering methods, which offer the possibility of localized heat and short laser pulses, have attracted more attention in the packaging and assembly of microelectronics, optoelectronics and micro-electro-mechanical systems (MEMS). In this chapter, the fundamentals of laser soldering including process control, thermal process, and formation mechanism of solder joint, as well as dissolution of substrate materials are discussed firstly. Then the development of a novel ultrasonic modulated laser fluxless soldering process is presented. Finally, reliability of laser soldered joints from the point of view of interfacial reaction and effect of intermetallic compounds (IMCs) is elaborated. In the last section, two case studies of laser soldering and reliability of solder joints in microelectronics packaging are presented.
12.2
Overview of laser soldering
In 1974, the CO 2 laser was first applied by CE Bohman in solder interconnections of microelectronic assemblies and the first commercially available laser soldering system appeared in 19761,2. Since then, research on laser soldering technology and equipment has developed rapidly. The ever decreasing size of electronic components, bringing a corresponding reduction in the size of interconnections, coupled with the availability of reliable laser units, has now made laser soldering technology economically viable and of particular value to the soldering of fine pitch devices. The laser soldering method has distinct advantages over conventional methods such as infra-red reflow, hot air reflow, vapour phase or hot-belt soldering, in which the entire assembly is passed through an oven. The key benefits of laser soldering are listed below. 1. Non-contact and local heating. The laser beam is precisely directed to the desired soldering location. This locally confined heat input will not 299 WPNL2204
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lead to thermal damage to the surrounding materials, especially to the heat-sensitive components. 2. Reliable solder joints. The relatively small rise in substrate temperature reduces mechanical stress and the rapid melting and cooling of the solder results in a fine microstructure which can improve the fatigue life of the solder joint. 3. Precise and controllable process parameters. The process parameters can be precisely controlled according to the different component types to attain uniform joint quality. Once programmed, the laser soldering system can provide repeatable results, run after run. 4. Flexible and easy to automate. The laser soldering process can be controlled in real time, and there is an opportunity to select various suitable wavelengths of lasers including Nd:YAG lasers and diode lasers, depending upon the materials. Combination of automation and laser results in high soldering speed and efficiency, especially the very high efficiency of diode laser systems (>40%), when compared to ovens or even other lasers, leading to greatly reduced energy usage. Earlier, large soldering areas were problematic for lasers. However, laser systems can now be equipped with scanner processing heads. With this equipment the soldering area can be scanned with a laser beam at high speed. By using split beams, several joints can be soldered simultaneously. One disadvantage of laser soldering is that it is not a mass-joining technique, as laser soldering is essentially a serial procedure, the beam being focused on each joint in turn. However, an advantage in soldering joints sequentially is that solder on adjacent solder pads is not molten simultaneously and hence the potential for solder bridging from one pad to the next is reduced. Owing to the above advantages, laser soldering methods show its potential benefits in the assembly of the highly reliable, heat-sensitive and electrostaticsensitive components, and has found more applications in the assembly of surface mounted devices (SMDs) such as resistors, capacitors, small outline packages (SOPs) and fine pitch quad flat packages (QFPs)3–5. Nowadays, laser soldering technology has become a more attractive solution in the solder bumping of plastic ball grid array (PBGA) packages and flip chip packages6–9, inner lead soldering of taped automated bonding (TAB) packages10, selective laser soldering of through-hole assemblies, lead-free soldering11, self-alignment of optical fibre in optoelectronic packages12–14, rework or desoldering on various package styles15–17, interconnection in microsensor manufacturing18, and localized laser assisted eutectic bonding in packaging of MEMS devices19,20. This chapter will focus on laser soldering for microelectronic packaging and assembly.
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Fundamentals of laser soldering
In the field of microelectronic packaging and assembly, laser soldering can be classified as laser reflow soldering, laser solder ball bonding and laser solder bumping according to its applications. The fundamentals of the abovementioned processes are the same, that is, a high-power-density laser beam is directed onto a small area of a potential joint for a short period. If the joint is irradiated for sufficient time and with sufficient power then both the solder and the local region of the components reach the solder melting temperature, allowing the solder to melt and flow. Once the laser beam is removed, the solder cools and solidifies to form the required solid metallurgical joint. The soldering material in laser soldering of microelectronics packaging and assembly includes solder paste, solder ball, solder wire, solder flux, thin film pre-coated/-plated on the substrate surface (pad) and on the component side (metallization layer, pins or leads). Critical factors for solder material in the laser soldering process are alloys used and flux composition. Typically, the following laser technologies are used in laser soldering. 1. CO2 laser 2. Nd:YAG laser 3. High-power diode lasers The CO2 laser is typically used in applications where high power density and high beam quality is required. However, in applications such as laser soldering, a very different power density regime is required and high beam quality does not bring any advantages to the application. The objective of laser soldering is to avoid any vaporization, simply to create melting of the solder and heating of the surrounding area to achieve satisfactory ‘wetting’ of the substrate21. Nd:YAG laser has advantages over the CO2 laser due to the fact that printed circuit board (PCB) substrates absorb less laser energy of the Nd:YAG laser at a wavelength of 1.06µm than that of the CO2 laser at 10.6µm, and the solder alloy can absorb more laser energy of the Nd:YAG laser that that of the CO2. Therefore, the Nd:YAG laser can provide high heating efficiency without heat damage to the PCB substrate Moreover, a Nd:YAG laser beam can be delivered by optical fibre, which makes it more flexible to the guidance of the beam to the solder joint than the CO2 laser, because the 10.6 µm wavelength of a CO2 laser beam needs to be delivered with a mirror system to the target instead of fibre. In recent years, semiconductor diode lasers at a wavelength of around 0.8 µm are becoming increasingly attractive options for selective soldering in microelectronics, due to the fact that the wavelength of the energy produced is highly absorbed by the metals used in solder but less readily absorbed by common PCB materials. These lasers also offer relatively high electrical conversion efficiency and efficient fibre optic energy delivery11,22,23.
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The laser soldering process is more time dependent than total energy dependent, in that time is required for the highly concentrated energy to spread throughout the solder mass and to heat the component parts to a high enough temperature for solder wetting to occur. The power required to produce good soldered joints is dependent upon the mass and thermal conductivity of the joint and can vary considerably with component style. The precise amount of power depends critically upon the physical shape of the pad and lead and whether solder paste or alloy is used. Since the heat energy is applied directly to the area to be heated, considerable heat loss is experienced through conduction to the cooler parts of the structure. The heat energy delivered by the laser beam must be able to supply this ‘lost’ heat. Since all the heat flows outwards through the structure from the target area of the laser beam, there is always the possibility of ‘punch through’ at this point unless the supply of heat energy is carefully controlled. It is therefore important to model the heat losses associated with the target area correctly.
12.3.1 Process control of laser soldering In order successfully to assemble a surface mounted component onto the PCB substrate, the laser soldering process must be controlled so as not to overheat the leads or solder joint by applying too much heat too quickly. The laser soldering process can be controlled in many ways. The simplest way is to monitor laser current feedback for laser power control. A real-time operating system enables accurate power vs. time profiles and their control. Systems can be equipped with a CCD camera and/or pyrometer. A CCD camera enables on-line visual monitoring of the process and a pyrometer enables online temperature monitoring; however, the following factors need to be taken into consideration when using pyrometers. Every material has different emissitivity. The angle of observation may affect the temperature reading and a different pin or pad size may affect the temperature value. A pyrometer needs to be calibrated for each specific case. The infra-red radiation method is a kind of non-contact process control method, which utilizes the change of infra-red radiation caused by temperature change during the laser soldering process to control both the laser power and heating time and to inspect the solder joint quality in real time. Using this method, an infra-red detector is embedded in the laser soldering equipment, and the temperature of the solder joint is measured in real time by using this infra-red detector. The formation of solder joints is inspected by measuring the temperature change during the laser soldering process. In this way, the formation and quality of solder joints can be controlled by changing laser power and heating time in real time. Once the temperature rises excessively, the laser input can be cut off instantly by a shutter to avoid burning the solder joint or lead of components.
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Figure 12.1 shows a schematic drawing of a Nd:YAG laser soldering system with an infra-red detector for process control24. The YAG laser at 1.06 µm is in the infra-red spectrum and is invisible to the human eye. A visible HeNe at 0.63 µm is coupled with the YAG laser to show the path during programming. The HeNe laser is positioned in the same optical path as the YAG laser so that the YAG will follow the exact path of the visible laser. The CCD camera is used for observation of alignment of the laser and the desired soldering location and also for the inspection of the laser soldering process. The laser input power is controlled by a programmed computer to provide precise laser energy input. The specially designed optical mirrors have multiple functions including (i) reflection of the Nd:YAG laser (wavelength 1.06 µm) totally and then focus onto the soldered location, (ii) focus of infrared radiation signal (wavelength 3~81 µm) caused by temperature rising of the solder joint onto the infrared detector, and (iii) blocking off the reflection of the Nd:YAG laser on the solder joint surface completely to avoid the interference of temperature inspection.
12.3.2 Thermal process of laser soldering Investigation of a thermal process can provide a fundamental and theoretical understanding of solder joint formation and its associated physical, chemical and metallurgical processes. The following two aspects make the formation of the solder joint during laser soldering unique:
Infra-red detector Computer Shutter
CCD
Infra-red radiation Laser Mirror optics
Laser beam
SMD Solder
PCB PCB X-Y work table
12.1 A Nd:YAG laser soldering system showing closed loop process using an infra-red detector for process control (provided by Microjoining Laboratory of Harbin Institute of Technology).
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1. Laser soldering offers a shorter and non-equilibrium heating process over the conventional reflow soldering methods such as infra-red reflow and hot air reflow which provide a long and uniform heating process. 2. The materials properties including abortion coefficients to the laser, thermal conductivity and heat capacity of pins and substrates of various electronic components are different. The small soldered area with complex geometry and relatively short heating duration which are characteristic of laser soldering for microelectronics makes the temperature measurement very difficult by experimental methods. Mathematical modelling can provide a better understanding of the laser soldering process and a valuable insight into which heat transfer mechanisms are important and how factors such as component dimensions, PCB characteristics and heating regime affect these. This in turn allows the laser parameters needed for successful soldering to be predicted25. The amount of laser beam power absorbed varies from material to material. For example, the Cu or kovar pins and solid solder alloy absorb 75% of the laser power whilst the PCB board absorbs only 25%. This information needs to be programmed into the model to make sure each material absorbs the correct amount of energy. In 1987, D.U. Chang proposed the first one-dimensional mathematic model to describe the thermal process of laser soldering of surface mounted devices26: r = a ln[2PAR(1 – exp (–0.975t/RC))]/πa2(Tm – T0)1/2 2
(1)
where r is standard thermal radius (radius of molten solder), a is radius of laser beam, P is laser power, A is absorption coefficient, R is the thermal resistance of solder to substrate, C is thermal capacitance, t is laser irradiation duration, Tm is the melting point of solder alloy, and T0 is the temperature of soldered component. A novel heat source used for modelling the laser soldering process of an optoelectronics butterfly package was proposed by D. Gwyer13. The temperature profile of a typical laser beam can be approximated by a Gaussian distribution. The Gaussian distribution laser heat source is defined as follows:
Q=
δ ⋅ I 0 –( r 2 / a 2 ) (– δ |z – z 0 |) e e 2 πr
(2)
where Q is laser heat source, I0 is peak laser intensity, and δ is absorption coefficient, r = ( x 0 – x ) 2 + y 2 is the distance from the centre of the laser beam, a is the radius of the laser beam, and x, y, z are spatial coordinates along the x-axis, y-axis, and z-axis respectively. Using this model, the laser energy will penetrate down the surface of each pin (in the z direction) and through to the part of the geometry where the
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solder is to be heated, and beyond through to the PCB underneath. When the penetration depth of laser energy down the surface of the pin is small compared to the thickness of the pin, the heat source can be treated as a surface heat source to simplify the model27,28. Figure 12.2 shows the simulated results of temperature distribution of a solder joint based on a surface heat source model during the laser soldering process28. The component is leaded quad flat package (QFP), and the width of pin is 0.4 mm and thickness is 0.14 mm. The PCB substrates are FR-4 resin and Al2O3 ceramic respectively. Since the thermal conductivity of FR4 is lower than that of the Al2O3 substrate, the temperature distribution on the surface of pins soldered on the FR-4 substrate is more uniform than that on the Al2O3 substrate. The temperature difference between the upper surface 350 Pin
Solder FR-4 resin substrate Laser powr 6 W
Temperature, °C
300 250
0.16 s
200 150
0.10 s
100
0.02 s
50 0 0.0
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0.6 0.8 Distance, mm (a)
1.0
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450 Pin
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400 Temperature, °C
350 300 250
0.325 s
200
0.250 s 0.025 s
150 100 50 0 0.0
0.2
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0.6 0.8 Distance, mm (b)
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12.2 Simulated results of temperature distribution of solder joint formed on (a) FR-4 resin substrate, and (b) Al2O3 ceramic substrate (from C. Q. Wang et al.28).
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of the pins and FR-4 substrate is smaller, resulting in a better shape of solder joint. On the other hand a large temperature gradient exists on the upper surface of the pins and Al2O3 substrate, and heat dissipates more steeply at the solder paste location. As a result of this, the formation of good solder joints in the case of Al2O3 substrate requires a higher level of laser power can than that of FR-4, and higher laser power can easily burn out the pins once inadequate contact between pins and solder paste occurs. Therefore, when assembling the electronic components on the ceramic substrate by laser soldering, preheating is suggested to lower the laser power and improve the quality of solder joint.
12.3.3 Formation of solder joint Figure 12.3 shows the schematic drawing of the laser soldering process for a surface mounted device. The dynamic process of a solder joint formation by laser soldering is shown in Fig. 12.4. 1. As shown in Fig. 12.4(a), temperature increased gradually after laser heating, and the solder at the centre of the laser spot begins to melt and agglomerate. The agglomerated solder droplet comes into contact with the pad on the substrate under surface tension, and then the solder around the solder droplet begins to melt and agglomerate through heat transfer, 2. The solder droplet starts to wet and spread on the pad when laser heating continues. The spreading direction is from component to the far end of the pad, as shown in Fig. 12.4(b). The solder droplet shows a strong adsorption effect during wetting, i.e., it adsorbs the surroundings of the previous spreading solder and forms a concave liquid-gas interface. 3. With the continuous radiation of the laser beam, the melted solder absorbs more heat energy, and the melted solder climbs on the vertical metallization of the component, adjusting the joint shape continually, as shown in Fig. 12.4(c). Laser
Metallization layer
SMD
Cu pad Solder paste
PCB
12.3 Schematic drawing of laser soldering process of a surface mounted device.
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4. The temperature of metallization on the component side increased under the irradiation of the laser. The solder on the far end of the pad begins to draw back and a good solder joint is achieved, as shown in Fig. 12.4(d). The shape of the solder joint is dependent on the location of the laser heating spot. The climb of the molten solder on the metallization layer of the component side is delayed by the drift of the laser spot location towards the far end of the pad, resulting in various joint shapes. The uneven temperature distribution of the solder joint during the laser soldering process plays an important role in the formation of the final solder joint shape. Considering the change of the free energy of the whole system: ∆G = ∆GV + ∆GS
(3)
where ∆GV is the volume item of free energy change and ∆GS is the surface item of free energy change. Assuming the component, substrate and solder volume are not dependent on the temperature after melting of the solder, ∆GV = 0, i.e., the system free energy change includes only surface item. When the temperature gradient exists on the solder surface, GS = A(t)γ (t) = ∑ Ai(t)γi(t)
(4)
where Ai(t) and γi(t) are surface area and surface energy per surface area at i position. Therefore, the solder joint shape formed by laser soldering depends on the temperature distribution of the solder surface. When the temperature
SMD
SMD
PCB
PCB
(a)
(b)
SMD
SMD
PCB
PCB
(c)
(d)
12.4 Illustrated diagram of dynamic formation of a SMD solder joint (a) the solder at the centre of laser spot begins to melt and agglomerate, (b) the solder droplet starts to wet and spread on the pad when laser heating continues, (c) the melted solder climbs on the metallization of the component, and (d) the formation of final shape of the solder joint.
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on the solder surface increases, γ(t) is reduced, and the joint shape in equilibrium tends to reduce the surface area of the solder. The relationship between surface temperature and joint shape is illustrated as Fig. 12.5. The formation of a SMD solder joint during laser soldering is a temperature forced convection mechanism. After laser heating, the solder paste surface reaches a high temperature and begins to melt because the solder paste, which is a suspension of solder powder in the organic flux, has a high absorption coefficient to the Nd:YAG laser and a low thermal conductivity. Therefore, the melted solder starts to agglomerate to reduce the surface energy. At the initial step of laser soldering, the component has a high thermal conductivity and behaves like a large cooling body. When the previously agglomerated solder droplet comes into contact with the PCB pad, the wetting process shows a strong direction characteristic, and the solder droplet wets towards the far end of the pad with high temperature. After the solder alloy melts, the physical properties of the solder change. Firstly, the interaction between the solder paste and laser is simplified, and it can be approximated as the surface absorption of pure solder alloy for the laser. At the same time, the thermal conductivity of the solder is increased significantly therefore the temperature increase of the solder surface is slowed down. In the meanwhile, the temperature difference between the solder surface and solder/pad interface decreases, and the temperature of the component side increases. At this time, the molten solder comes into contact with the metallization layer of the component, and the solder climbs towards the metallization to lower the system surface energy. As a result of this, the solder droplet at the far end of the pad is drawn back to form the solder joint finally, as shown in Fig. 12.5(a). Laser Laser
T2, γ2
T 1, γ 1 Solder (a) T1 > T2, γ1 < γ2
SMD
SMD
Pad
Pad (b) No climbing
Laser
T2, γ2
SMD
T 1, γ 1 (c) T1 < T2, γ1 > γ2
Pad
12.5 Illustrated relationship between the location of laser spot and shape of solder joint.
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If the laser heating spot deviates from the component significantly, the melted solder will agglomerate at the far end of the pad and cannot be in contact with the metallization because the surface tension is reduced when temperature is increased, and then the solder joint cannot be formed, as shown in Fig. 12.5(b). Therefore, to achieve a good solder joint shape after the solder wets the far end of the pad, the temperature of solder surface near the component must be higher than the temperature of solder surface at the far end of the pad, i.e., the laser heating position should be near the metallization of the component to make sure that the solder is drawn back under surface tension before solidification occurs, and the climb process of the solder on the component side goes on smoothly, as shown in Fig. 12.5(c).
12.3.4 Dissolution of the materials being soldered During the laser soldering process, the metallization layer at the component side and the pad at the PCB substrate side will dissolve into the molten solder. Sufficient dissolution of the materials being soldered is essential to clean the desired soldering surface and to enhance the strength of the solder joint. In the conventional soldering process, the dissolution thickness of the materials being soldered can be several microns to some ten microns. However, the thickness of the pad or metallization layer in the microelectronic assembly is very thin and it is necessary to control the dissolution during the soldering process or the pad or metallization layer will spall off and fail. For example, at a soldering temperature of 250 °C, Au film with 6 µm thickness will be completely dissolved within 1 s into the molten eutectic solder. The higher the temperature, the higher the dissolution rate will be. Another risk of complete dissolution of metallization is dewetting if the solder does not wet on the next layer. During laser soldering, the temperature of the whole system and the dissolution rate of the soldering materials change continually. At the same time, the solder alloy gradually melts and its contact area with the pad is increased. With the dissolution of the pad or metallization layer occurring, the concentration of solute atoms in the molten solder is increased continuously, which reduces the dissolution rate. All the factors mentioned above make the computation of dissolution of the pad or metallization layer during laser soldering more complicated. The dynamic dissolution of pad or metallization into the molten solder under the nonequilibrium temperature field of laser soldering condition can be achieved through solving the kinetics differential equation of solid-liquid mass transfer by numerical calculation29. The equation of dissolution reaction rate is defined as follows: dQ = aS( C1 – C ) dt
(5)
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where Q is the volume of dissolution of material into liquid solder, t is the contact time between the liquid and solid material, S is the contact area between liquid and solid materials, a is the dissolution rate coefficient, Cl is limited solubility, C is the concentration of solute element in the molten solder. Since Q can be described as Q = ρsVC
(6)
where V is volume of liquid solder, and ρs is the density of liquid solder. Equation 12.5 can be expressed as: d ( VC ) aS = ( Cl – C ) ρs dt
(7)
When the volume and temperature is constant, the solution of equation (12.5) is given by: S (8) Q = Vρs Cl 1 – exp – at V During laser soldering, the volume of liquid solder V, the contact area of liquid and solid material S and the dissolution rate coefficient changes with laser heating time and temperature therefore equation 12.8 is a variable coefficient differential equation for concentration of solute element. After discretization of equations 12.5 and 12.6: Q( n + 1) – Q( n ) 1 = [a(n) + a(n + 1)][S(n) + S(n + 1)][Cl(n) 8 ∆t n
+ Cl (n + 1) – C(n) – C(n + 1)]
C( n + 1) =
Q( n + 1) ρs V ( n + 1)
(9) (10)
where n is coordinate of time gridding, and t is time step. Then the complicated solution of the differential equation can be simplified as a solution of algebra equation. Assuming
∆t n [ a( n ) + a( n + 1)][ S( n ) + S( n + 1)] 8 Substitute equation 12.10 with 12.9, then, A( n ) =
Q( n + 1) =
Q( n ) + A[ Cl ( n ) + Cl ( n + 1) – C( n )] 1 + A /[ ρs V ( n + 1)]
(11)
(12)
By combining equations 12.11 and 12.12, the numerical computation of dissolution under unstable temperature field can be achieved.
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Figure 12.6 shows the dynamic curves of pad materials dissolution into the solder achieved by numerical calculation of the above-mentioned dissolution differential equations based on the temperature field computation results29. The pad material on ceramic substrate is Ag thick film conductor, and the pad on FR-4 resin substrate is Cu thick film conductor. Due to the solubility limit difference in Sn between Cu and Ag, there is a big difference in the dissolution dynamic process. The solubility limit of Cu in Sn is small, and it can be saturated during heating when the dissolution rate slows down. However, the solubility limit of Ag in Sn is bigger, the dissolution rate is increased
Dissolution rate, µm/s
Al2O3 ceramic substrate laser power: 20 W
0.8
0.75 0.6 Dissolution rate
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5 FR-4 resin substrate Laser power 12.5 W 4
4 3 3
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2 Dissolution thickness
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Dissolution thickness of Cu, µm
Dissolution velocity, µm/s
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0 1.2
(b)
12.6 Dynamic curves of pad materials dissolution into the eutectic SnPb solder achieved by numerical calculation (a) dissolution of Ag pad on ceramic substrate, and (b) dissolution of Cu pad on FR-4 substrate (from C.Q. Wang et al.29).
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constantly during heating, and then the amount of the dissolution is increased also. Therefore, the temperature needs to be accurately controlled when soldering is performed on Ag thick film conductor.
12.4
Fluxless laser soldering
Various fluxless soldering methods have been developed in recent years including the ion/atom sputtering method, plasma fluxless soldering and laser fluxless soldering. It has been shown that it is possible to perform laser fluxless reflow soldering of solder discs on pre-tinned Cu pads9. It was demonstrated that solder bumps could be formed on the pretinned Cu pads of a FR-4 substrate using a CO2 laser under an argon atmosphere, when solder discs instead of solder balls were employed. Pac Tech-Packaging Technologies GmbH of Germany and Pac Tech USA developed a new laser-based solder jetting technology, solder Ball Bumper Jet (SB2-Jet)14. This technology fulfils all the needs of fluxless soldering, local heating and reflow, no mechanical contact and stress during soldering, high solder alloy flexibility and capability of 3D-packaging. With a throughput of 10 balls/s, the SB2-Jet fulfils most of the requirements for today’s packaging of optoelectronics and MEMS devices in production. A further increase in speed to 20 and 30 balls/s is now in progress for the next generation. An additional feature of the SB2-Jet technology is the repair option and repair capability. This permits individual removal and replacement of solder balls and solder contacts and gives an increase in the yield and productivity of cost-intensive high end devices. M. Li and C.Wang developed a novel ultrasonic modulated laser fluxless soldering method30. Ultrasonic modulated laser fluxless soldering means that a continuous laser was modulated into a pulsed laser with a high frequency of 20 kHz. When this is done the pulsed laser has the function of an ultrasonic wave. The ultrasonic oscillation temperature was induced at the solder droplet surface heated by the ultrasonic modulated laser, which gave rise to the ultrasonic mechanical vibration at the solder droplet surface. The ultrasonic mechanical vibration propagated from the solder droplet surface to the soldering interface contributed to the solder wetting behaviour due to the ultrasonic cavitation effects. Figure 12.7 shows a schematic drawing of the ultrasonic modulated laser fluxless soldering method. The soldering process is performed in the low vacuum environment. The laser output is modulated from continuous Nd:YAG laser with 0.7mm focused laser beam diameter, which can provide ultrasonic modulated laser with 20 kHz frequency and 0.5 duty ratio. Figure 12.8 shows that the ultrasonic oscillating temperature with an amplitude of about 3 °C appeared at the surface of the solder droplet, and the temperature oscillation was synchronous with the modulated laser frequency30.
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Nd: YAG laser Lens Quartz glass 8 Pa vacuum environment
Solder Cu pad PCB substrate
12.7 Schematic drawing of ultrasonic modulated laser fluxless soldering method.
264 Vibration Temperature
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258
0.0
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–0.5 0
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100 150 Time, µs
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Vibration amplitude, µm
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254 250
12.8 Vibration of the solder droplet surface under the ultrasonic modulated laser after heated for 1.5 s. Laser power, 14W; duty ratio, 0.5; frequency, 20 kHz (from M.Y. Li et al.30).
The mechanical vibration with amplitude of 0.6 µm was also observed at the surface of the solder droplet, and the mechanical vibration frequency was about 20 kHz, which coincided with the modulated laser frequency and was also synchronous with the temperature oscillation frequency. According to thermodynamics theory, the ultrasonic mechanical vibration of the solder droplet is induced by the effect of the thermal expansion of the molten solder subject to the oscillating temperature field. According to wave-transfer theory, the wave pressure at the wetting interface transmitted from the surface of the solder droplet can be calculated by numerical modelling. As shown in Fig. 12.9, the pressure variation was induced at the solder wetting interface and negative pressure was observed. According to
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Pressure, × 104 Pa
1.5 1.0 0.5 0.0 –0.5 Negative pressure –1.0 0
50
100
150 200 250 Heating time, µs
300
350
400
12.9 Pressure variation in the solder/pad wetting interface under the ultrasonic modulated laser, indicating that the negative pressure occurred at the interface (from M.Y. Li et al.30).
ultrasonic cavitation theory, tearing force and vacuole can be generated at the wetting interface under negative pressure conditions, and then the vacuole will be broken in the ensuing compressive force, generating strong shock waves. Such a physical effect will be circulated at the soldered interface, and then the ultrasonic cavitation effect can remove the oxide film on the pad surface and result in a pronounced effect in improving the solder wetting behaviour. The ultrasonic modulated laser fluxless soldering method has many advantages. For example, it can be performed in a low vacuum environment, and the ultrasonic and heat energy is only added on the desired soldering area without damaging the component.
12.5 Reliability of solder joint Solder joint reliability has attracted tremendous attention because it not only transfers electrical signals and thermal energy from the devices, but also maintains mechanical integrity. One of the major influences on reliability was the solder joint’s microstructure. During the laser soldering process, the solder and pad will react with each other to form the intermetallic compounds (IMCs). The formation of intermetallic compounds between a solder ball and pad promotes wetting of liquid solder on the substrate surface, and it is believed to the develop a strong interconnection. However, the presence of excessive IMCs will weaken the long-term reliability of a solder joint, not only because most IMCs are brittle, but also because the coefficient of thermal expansion (CTE) mismatch exists between the IMCs and the solder
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joint. Therefore, investigation of the interfacial reactions is important in the optimization of laser soldering parameters, and is crucial to the understanding of solder joint reliability. In this section, two case studies on the reliability of solder joints were presented from a point of view of the interfacial reaction and IMCs formation during the laser soldering method.
12.5.1 Case 1: laser reflow soldering Laser reflow soldering can be used for solder bumping of flip chip and PBGA packages. The short laser irradiation duration provides finer microstructure within the solder bumps, and the finer microstructure results in better creep resistance compared to the solder bumps formed by the standard oven reflow methods such as infra-red reflow and hot air reflow method. Unlike the assembly process of QFP and other surface mounted components in which the solder pastes are subject to the reflow process only once, assembly of PBGA components will undergo the reflow process twice. Firstly, the solder balls were connected to the metallization layer of BT resin substrate to form solder bumps, and then the solder bumps were connected to the Cu pads of PCB substrate to form solder joints (Fig. 12.10). The IMCs formed at the interface of solder/Au/Ni/Cu/BT substrate during a first reflow bumping process will change when subjected to a second reflow soldering process. Evolution of IMCs at the solder/metallization interface during reflow bumping and reflow soldering processes are crucial for understanding the reliability of PBGA solder joints. Figure 12.11 shows an SEM cross-section of a solder bump formed by a Nd:YAG laser. The solder ball is eutectic SnPb with a diameter of 760 µm, and the metallization layer is Au/Ni/Cu with a 2 µm thickness of Au layer31. It was observed that lots of needle-like AuSn4 IMCs formed at the interfacial region of the solder bump, and grew sideways from the solder-metallization interface into the solder when the laser power is 15 W, and heating time is BT substrate
Solder bump
Solder joint
Au/Ni/Cu
Au/Ni/Cu
Cu
BT substrate
FR-4 PCB
(a)
(b)
12.10 Schematic drawing of PBGA solder bump and solder joint. (a) First reflow to form solder bump, and (b) the second reflow to form solder joint.
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200 ms, as shown in Fig. 12.11(a). The same phenomenon could also be found in Pb-free laser reflowed solder bumps when the same Au/Ni/Cu metallization layer is used31. The formation of these AuSn4 needle-like compounds was due to the rapid cooling rate of the solder bump and the large temperature gradient at the solder bump interface. The direction of the temperature gradient at the interface is the preferred growth direction of AuSn4 IMCs. During the solidification process of the solder bump, atom supply to form the IMCs depended on the heat flux and the diffusion of solute atoms. Therefore, the IMCs grew faster in the transfer direction of heat and mass. The needle-like AuSn4 IMCs disappeared at the interfacial region and distributed uniformly within the bulk solder as particles with the increase of laser irradiation duration, as shown in Fig. 12.11(b). When the heating time is long enough, all the Au in the pad dissolved into the whole solder ball. The dissolved Au reacted with Sn to form AuSn4 again during the solidification process. The detailed interfacial reaction kinetics during laser soldering can be found in the earlier research of Y. Tian8. Figure 12.12 shows SEM images of solder joints subjected to the first laser reflow and the second infra-red reflow. The interfacial microstructure of the solder joint formed by the second infra-red reflow was different, and strongly dependent on the first laser reflow parameters. When the laser heating time was small, the continuous AuSn4 layer dissolved partially, and the needle-like AuSn4 disconnected with the solder/metallization interface and spread inside the solder bulk as a rod shape, as shown in Fig. 12.12(a). The formation of the AuSn4 rods could be explained as follows: (i) needlelike AuSn4 broke off from the interface and fell into the solder since they were brittle and stressed by higher flowing rates of the molten solder; (ii) then they could further dissolve into solder because of the high solubility of Au in the molten solder; and (iii) because of the longer secondary reflow (b)
(a)
AuSn4
AuSn4 Au
12.11 SEM cross-section of a solder bump formed by a Nd:YAG laser of (a) 15 W heating for 200 ms, and (b) 20 W heating for 400 ms (from Y.H. Tian et al.31).
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(b) AuSn4
Pb-rich AuSn4
AuSn4
(c)
Pb-rich
12.12 SEM images of solder joints subjected to the first laser reflow and the second infra-red reflow at 207 °C for 53s (a) laser reflow 20 W, 100 ms, (b) laser reflow 25 W, 100 ms (c) laser reflow 20W, 200 ms (from Y.H. Tian et al.31).
time (53s) used, the needle-like AuSn4 would change into a rod-shape under a slow solidification rate. In addition to the AuSn4 rods near the interface, the segregation of the AuSn4 particles was also found at the eutectic cell boundary of Pb-Sn solder. At the same time, most of the Pb-rich islands were found near the segregation region of the AuSn4 particles. Segregation of AuSn4 particles indicated that re-precipitation of AuSn4 occurred during the solidification of the solder. Morphology and distribution of AuSn4 IMCs within the solder joint formed by laser reflow followed by the second infra-red reflow was affected by the first laser reflow parameters. When the laser heating time was small, a thick layer of AuSn4 IMCs was formed at the interface of the solder bump, and would not dissolve completely after the second infra-red reflow. The remaining AuSn4 particles at the interface were brittle, and would affect the integrity of the solder joint. On the other hand, when the laser input energy was large, after the second infra-red reflow, all the AuSn4 disappeared at the interface of the solder joint, but a large number of Pb-rich islands would appear, as shown in Fig. 12.12(c). These Pb-rich islands can also degrade the reliability of solder joints. Therefore, laser reflow processing parameters have an important
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effect on the reliability of solder joints. To achieve a PBGA solder joint with high reliability, modest laser reflow parameters should be adopted, as shown in Fig. 12.12(b).
12.5.2 Case 2: laser solder ball bonding The laser solder ball bonding method was applied to solder Sn3.5Ag0.75Cu Pb-free solder balls onto rectangular pads of a kind of microelectronics assembly32. The diameter of solder ball was 0.12 mm, and Au/Ni metallization on Cu pad with three Au thicknesses of 0.1, 0.9 and 4.0 µm was adopted on the component side, and the underlying Ni thickness was 0.2 µm. The processing parameters are as follows: laser power was 7 W, laser pulse was 5 ms, the temperature range of thermal cycle was between –40 °C and 125 °C, The ramp rate was about 16.5 °C/min, and the dwell time at both extreme temperatures was 20 minutes respectively. The interfacial microstructure of solder joints with various Au thicknesses shows great variation and results in significant effects on the reliability of solder joints. Sn3.5Ag0.75Cu solder joints with 0.1 µm Au metallization with and without thermal cycles are shown in Fig. 12.13. The thin Au metallization was dissolved quickly during the laser soldering process, allowing Ni to expose the solder bulk. Little Cu-Ni-Au-Sn IMC was formed at the interface, even after 1000 thermal cycles, and it was so thin that the exact atomic ratio cannot be determined by EDX. However, obvious coarsening of Ag3Sn particles in the solder bulk took place with increase in the number of thermal cycles. As for solder joints with 0.9 µm Au metallization, as shown in Fig. 12.14, the morphology of AuSn4 IMC changed gradually from needle-type to chunkytype after 1000 thermal cycles, and the thickness increased significantly. In
Cu-Ni-Au-Sn
Ag3Sn 10 µm
10 µm (b)
(a)
12.13 SEM images of Sn3.5Ag0.75Cu solder joints formed laser reflow with 0.1 µm Au metallization at different thermal cycles (a) 0 cycle, (b) 1000 cycles (from H.T. Chen et al.32).
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AuSn4 AuSn4 Cu-Ni-Au-Sn
10 µm
10 µm (a)
12.14 SEM images of Sn3.5Ag0.75Cu solder joints formed laser reflow with 0.9 µm Au metallization at different thermal cycles (a) 0 cycle, (b) 1000 cycles (from H.T. Chen et al. 32).
addition, a very thin Cu-Ni-Au-Sn IMC layer was formed below the AuSn4 IMC layer. The composition of Cu-Ni-Au-Sn cannot be accurately determined. For solder joints with 4 µm Au metallization, as shown in Fig. 12.15, AuSn, AuSn2 and AuSn4 IMCs with Cu concentration less than 0.8 wt% were observed at the interface. No Ni was detected in the Au-containing IMC, and it may be because the dissolution of the thick Au metallization and subsequent reaction with solder to form Au-containing IMCs need a longer time than the samples with 0.1 and 0.9 µm Au metallizations, and the thick IMCs at the interface present diffusion barriers for Ni to be involved in the interfacial reaction. AuSn4 IMC appears to be needle-type, while AuSn and AuSn2 IMCs formed a continuous planar layer after rapid wetting reaction. As the thermal cycles progressed, the interfacial morphology of AuSn4 IMC became planar. From the point of view of kinetics, needle-type morphology has a larger surface contact with solder bulk than a flat surface. In the wetting reaction, the rapid gain in IMC formation may compensate the surface energy spent in growing the needle-type IMC. However, in the following solid state reactions, the gain disappears, and the IMC changes to a layer-type with a flat surface because of the high interfacial energy between solid solder and IMC. At the same time, AuSn4 IMC grew at the expense of AuSn and AuSn2 IMCs, and the Cu and Ni content in AuSn and AuSn2 was less than 1 wt% during the phase change process. The thickness of AuSn4 IMC reached about 20 µm, and Cu and Ni content in AuSn4 was still less than 3 wt% after 1000 thermal cycles. As shown in Fig. 12.16(a), for the solder joints with 0.1 µm Au metallization, the cracks mainly occurred through solder near the interface of solder/IMC on the component side after 1000 thermal cycles. A detailed examination shows that the failure is characterized by intergranular cracking, which is
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AuSn4 AuSn4
AuSn2
AuSn2 AuSn
10 µm
10 µm
(a)
(b)
12.15 SEM images of Sn3.5Ag0.75Cu solder joints formed laser reflow with 4.0 µm Au metallization at different thermal cycles (a) 0 cycle, (b) 1000 cycles (from H.T. Chen et al.32).
mainly induced by grain boundary sliding of Sn-rich phases, as shown in Fig. 12.16(b). The likely reason is that there is not enough Au dissolved to form finely dispersed IMC particles to pin up the grain boundaries of Sn-rich phases. Therefore, the stress at the interface of solder and pad accelerates the movement of dislocations and vacancies at the grain boundaries, resulting in the grain boundary sliding of Sn-rich phases. The cracks of solder joints with 0.9 µm Au metallization were also observed at the same location as the solder joints with 0.1 µm Au metallization, as shown in Figs 12.16(c) and (d). However, the cracks were not so significant, and no obvious grain boundary sliding phenomenon was found. The finely dispersed IMC particles, which precipitated out due to supersaturated Au, may play an important role in pinning up the grain boundaries of Sn-rich phases. For 4.0 µm Au metallization, as shown in Figs 12.16(e) and (f) only micro-cracks were found on the AuSn4 IMC surface. An excessive AuSn4 IMC formation results in cracks in IMC because the thick IMC can easily crack at relatively light mechanical loads. A typical failure of solder joints is low cycle fatigue caused by the coefficient of thermal expansion (CTE) mismatch among different assembly materials while the system is subjected to environmental temperature change and power on/off. In thermal cycling experiments, the creep and fatigue microscopic failure mechanism can be induced due to the high homologous temperature condition. The time dependent creep is likely to occur together with fatigue. Creep failure is characterized by intergranular cracking because of grain sliding, while fatigue failure occurs due to formation and propagation of surface cracks. The creep-fatigue failure can be viewed as the interaction between fatigue cracks and grain boundary cracks. As shown in Fig. 12.17, the stress, which is caused by global CTE mismatch between the component and stainless steel and local CTE mismatch between
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IMC
Solder Grain boundary sliding Cracks
(a)
(b)
IMC Solder Cracks
(c)
(d)
IMC
Cracks Solder
(e)
(f)
12.16 Sn3.5Ag0.75Cu Solder joints after 1000 thermal cycles (a) 0.1 µm Au metallization, (b) magnified view of (a), (c) 0.9 µm Au metallization, (d) magnified view of (c), (e) 4.0 µm Au metallization, (f) magnified view of (e) (from H.T. Chen et al.32).
pad and solder, is mainly concentrated in the AuSn4 IMC layer at component side. However, the cracks do not necessarily occur through AuSn4 IMC since the strength of the IMC is generally greater than the strength of the solder. Furthermore, the thickness and morphology of IMC also play an important role in determining where cracks are initiated. A result of finite element modelling showed that the equivalent creep strain and plastic strain mainly
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12.17 A finite element model shows the maximum equivalent stress location.
concentrated at solder near the AuSn4 IMC, which agrees well with sites of cracks observed in the solder joints with 0.9 µm Au metallization subjected to the temperature-cycle experiment. The accumulated inelastic strain by finite element modelling is shown in Fig. 12.18, it is noted that the equivalent creep strain range is much greater than the equivalent plastic strain range. It indicates that the creep strain effect is dominant in contributing to the fatigue damage of solder joints during thermal cycling.
12.6
Summary and future trends
This chapter introduced the topic of laser soldering in microelectronics packaging and assembly, presenting the fundamentals of laser soldering, a novel laser fluxless soldering method and reliability of the laser soldered joint. The contents covers the laser reflow soldering of surface mounted devices, laser solder bumping of PBGA packages and laser solder ball bonding technology. Laser soldering technology is becoming increasingly important to electronics assembly as the trend to closer packing densities and finer line geometries continues. For example, fine-pitch flat-pack devices with lead spacings of
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0.03
0.02
0.01
0.00 0
2000
4000
6000 8000 10000 12000 14000 16000 Thermal shock time (s)
12.18 History of accumulated inelastic strain.
0.012 inch and ‘chip-on-tape’ packages having lead pitches of 0.010 inch had already been used for many applications. At these high pin densities, more commonplace methods for automatic soldering such as hot air and infra-red techniques are package dependent and not reliable. The future development of laser soldering will focus on the following aspects: 1. Diode laser soldering technology. It is believed that this technology will become more and more attractive for selective soldering and leadfree soldering11. The diode laser soldering technology was considered to be the next generation in non-contact soldering21. 2. Fluxless laser soldering in packaging of optoelectronics and MEMS devices14,19. The packaging of optoelectronics and MEMS devices are presenting challenges for interconnection and soldering technology. These requirements can no longer be met with standard flux-based processes which use a long temperature reflow profile and are implementing a lot of mechanical handling steps and processes. Basically, the packaging of these devices requires fluxless soldering, no thermal stress by localized heating, low or no mechanical contact or damage to sensitive membranes in MEMS or optical components. Laser fluxless soldering technology can fulfil the above specific needs in these applications. 3. Fundamentals study of laser soldering. Difficulties during the investigation of laser soldering are temperature measurement and solder joint shape prediction, as the solder joint is very small and the laser soldering is a non-equilibrium process. However, the prediction of solder joint shape is very important for the further prediction of solder joint reliability.
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Therefore, it is necessary to study the fundamentals of laser soldering to obtain the reliable solder joints. For students and researchers who would like to investigate laser soldering further, there are a variety of excellent books. A handy introduction and guide about laser soldering can be found from the ASM handbook33. In addition, chapters on laser soldering are included in the books by Klein and Lea34,35.
12.7
Acknowledgements
The author greatly acknowledges the Microjoining Laboratory of Harbin Institute of Technology (HIT) for providing the research results and some of the diagrams and photographs for this chapter. The author also appreciates the technical references of ViTechnology, Inc., Pac Tec, Inc., the Fraunhofer Institute, and research group related to laser soldering of the University of Arkansas, The University of Hull, University of Greenwich, Hongik University and Tampere University of Technology and so on.
12.8
References
1. Bohman C.F., ‘The laser and microsoldering’, Society of Manufacturing Engineers, Technical Paper AD 74810 1974 pp 19. 2. Kujawa T., ‘Laser soldering boosts productivity’, Lasers and Applications, September 1982 93–94. 3. Lea C., ‘Laser soldering of surface mounted assemblies’, Hybrid Circuits, No. 12, January 1987. 4. Ismail Fidan, ‘CAPP for electronics manufacturing Case Study: Fine Pitch SMT laser soldering’ Journal of Electronic Packaging, 2004 126(3) 173–176. 5. Whitehead D.G. and Polijanczuk A.V., ‘Reflow soldering by laser’ ASME Heat Transfer Division, 1990 143 47–56. 6. Changhai Wang and Andrew S. Holmes, ‘Laser-assisted bumping for flip-chip assembly’, IEEE Transactions on Electronics Packaging Manufacturing, 2001 24 (2) 109–114. 7. Xiaodong Zhang, Chunqing Wang and Yanhong Tian, ‘Design of laser scanning solder bumping system’, The 4th International Conference for Electronic Packaging and Technology, Shanghai China 2003, 141–144. 8. Yanhong Tian, Chunqing Wang, Xiaodong Zhang and Deming Liu ‘Interaction Kinetics between PBGA Solder Balls and Au/Ni/Cu Metallisation during Laser Reflow Bumping’, Soldering & Surface Mount Technology, 2003, 15(2) 17–21. 9. Jong-Hyun Lee, ‘Fluxless Laser Reflow Bumping of Sn-Pb Eutectic Solder’ Scripta Mater., 2000 42(8) 789–793. 10. Aschenbrenner, Fluxless Flip Bonding on Flexible Substrates: A comparison between adhesive bonding and soldering, Soldering & Surface Mount Technology, 1996 8(2) 5–11. 11. Chaminade C., Fogarassy E., Boisselier D., ‘Diode laser soldering using a lead-free
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13.
14. 15. 16.
17. 18.
19. 20.
21. 22. 23. 24.
25.
26. 27.
28.
29.
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filler material for electronic packaging structures’, Applied Surface Science. 2006 252 4406–4410. Yan Bohan, Wang Chunqing and Zhang Wei, ‘The Microstructure of eutectic Au-Sn and In-Sn solders on Au/Ti and Au/Ni metallizations during laser solder bonding process for optical fiber alignment’ The 7th International Conference for Electronic Packaging and Technology, Shanghai China 2006, 345–349. Gwyer D., Bailey C., Pericleous K., Philpott D. and Misselbrook P., ‘Mathematical Modelling: A laser soldering process for an optoelectronics butterfly package’, Inter Society Conference on Thermal Phenomena, 2002 121–127. Elke Zakel, Lars Titerle, Thomas Oppert, Ronald G. Blankenhorn, ‘Laser solder attach for optoelectronics packages’, www..pactech.de. Paul P.E. Wang, Steven Perng and Erick Russell ‘Laser rework technology’, http:// www.vitechnology.com. Hanreich G., Musiejovsky L., Wolter K.J., Fasching M., Nicoics J., ‘Optimisation of a laser soldering/desoldering process for Flip-Chips using a new thermal simulation tool’, www.iemw.tuwien.ac.at. Electronic Packaging, 68–69. Edgar Cerda, ‘Laser soldering applications for RF shield rework’, http:// www.vitechnology.com. Gernot Hanreich, Johann Nicolics, Martin Mundlein, Hans Hauser, Rupert Chabicovsky. ‘A new bonding technology for human skin humidity sensors’, Sensors and Actuators A 2001 92 364–369. Tan A.W.Y., Tay F.E.H., Zhang J., ‘Characterization of localized laser assisted eutectic bonds’, Sensors and Actuators A 2006 125 573–585. Yi Tao, Ajay P. Malshe, Willam D. Brown, ‘Selective bonding and encapsulation for wafer-level vacuum packaging of MEMS and related micro systems’, Microelectronics Reliability, 2004 44 251–258. Whitehead D.G. and Foster R.J., ‘Soldering with light’, Assembly Automation, 1995 15(2) 17–19. Greenstein M., ‘Optical absorbtion aspects of laser soldering for high density interconnects’, Applied Optics, 1989 28 (21) 1–5. Goldberg G., ‘Diode laser soldering technology: The next generation in Non-contact soldering’, Surface Mount Technology, 2004 (9) 262–266. Xuwen Liang, Chunqing Wang, Weiyan Jiang, Yiyu Qian ‘Solder Joint Qualification Based on the Wetting Recognition in SMT Laser Soldering’, Chinese Journal of Mechanical Engineering, 1998 34 (5): 63–69. Beckeet P.M., Fleming A.R., Gillbert J.M., Whitehead D.G., ‘Numerical modelling of scanned beam laser soldering of fine pitch packages’, Soldering & Surface Mount Technology, 2002 14 (1) 24–29. Chang D.U., ‘Experimental Investigation of Laser Beam Soldering’, Welding Journal, 1986 65 (10) 33–41. Chunqing Wang, ‘Experimental investigation and numerical simulation of the SMT laser micro-soldering thermal process’, Soldering & Surface Mount Technology, 1991 8 29–31. Chunqing Wang, Yiyu Qian and Yihong Jiang, ‘A numerical simulation model of SMT laser micro-soldering thermal process’, Chinese Journal of Lasers, 1993 20 (2) 126–131. Chunqing Wang, Yiyu Qian and Yihong Jiang, ‘Dynamic dissolution of the base metal into the molten solder in SMT laser micro-soldering’, Chinese Journal of Mechanical Engineering, 1993 29(2) 52–57.
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30. Mingyu Li, Chunqing Wang, Han Sur Bong, ‘Development of a flux-less soldering method by ultrasonic modulated laser’, Journal of Materials Processing Technology. 2005 168 303–307. 31. Yanhong Tian, Chunqing Wang, Deming Liu and Peter Liu ‘Intermetallic compounds formation at interface between PBGA solder ball and Au/Ni/Cu/BT PCB substrate after laser reflow processes’, Materials Science and Engineering B, 2002 95 (3) 254–262. 32. Chen H.T., Wang C.Q., Li M.Y., ‘Numerical and experimental analysis of the Sn3.5Ag0.75Cu solder joint reliability under thermal cycling’, Microelectronics Reliability, 2006 46 1348–1356. 33. Kelly Ferjutz, ASM Handbook: Welding, Brazing and Soldering, Vol. 6, 10th edition, ASM International, Ohio, 1993. 34. Klein Wassink R.J. and Verguld M.M.F., Manufacturing Techniques for Surface Mounted Assemblies, Electrochemical Publications, IOM, 1997. 35. Lea C., A Scientific Guide to Surface Mount Technology, Electrochemical Publications, Ayr, Scotland, 1988.
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13.1
Flux in soldering
In general the surface of a metal is covered with contaminants and an oxide layer (see Fig. 13.1). The contaminants are composed of absorbed gases, moisture and oil, and the oxide layer has a nanometer-scale thickness. The oxidation of a metal surface under ambient conditions is a natural phenomenon. About 1.5 nm of oxide forms within a few minutes and grows rapidly to about 2 nm within a week, and then much more slowly to a maximum thickness of 7 nm after as long as 20 years [1]. Since Sn2+ and Sn4+ have a similar value of free energy of formation, the surface oxide of Sn is likely to be a mixture of SnO and SnO2. The contaminants and oxide layer must be removed to obtain a reliable solder joint. Organic contaminants can be removed during the cleaning process by a cleaner such as IPA (isopropyl alcohol), methanol or ethanol, and the oxide layer is removed usually by flux. The roles of flux during soldering can be summarized as follows: 1. to remove the oxide from the metal surface 2. to prevent re-oxidation of the clean metal surface 3. to remove the oxide on the liquid solder to reduce the surface tension and increase fluidity, and 4. to assist heat transfer from heat source to the solder joint [2]. Flux can be supplied by a dispensing or spraying method, or by dipping into a flux bath. Abietic, neoabietic, palustric, dehydroabietic, pimaric resin acids are used for the flux resin acids. The oxides of CuO and SnO can be removed from the solder joint as follows: 2 R-COOH + CuO → (R-COO)2Cu + H2O 2 R-COOH + SnO → (R-COO)2Sn + H2O R: carboxyl residue (abietic acid , C19H29) 327 WPNL2204
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Contaminant layer Oxide layer Microstructurally and chemically inhomogeneous layer Base metal
13.1 Oxide layer on metal surface.
13.2
Demand for fluxless soldering
In conventional soldering processes, it is usual for liquid flux to be applied to the solder joint to improve the wettability of molten solder. After soldering, flux residues are produced around the solder joint. Since these residues are corrosive, they can cause long-term reliability problems [3]. Thus, the flux residue should be removed, and the halogenated or chlorofluorocarbon (CFC) solvents are generally used to remove the rosin-based flux residues. These hazardous solvents cause environmental concerns such as ozone depletion [4], and this forces us to develop new environment-friendly solvents, cleaning methods, no-clean fluxes, and fluxless soldering technologies. Meanwhile, flux residues which can contaminate optical surfaces have a negative effect on the optoelectronics like deflecting or attenuating of laser signal by flux residues [5]. The tendency of shrinking size and fine pitch of all electronic components is another reason to avoid flux application due to the difficulty of cleaning flux residues. For these reasons, the development of a fluxless soldering process has been pursued continuously.
13.3
Fluxless soldering process
In a fluxless soldering process, the oxide layer of SnO can be removed by several methods. In this chapter, reducing gas atmosphere such as hydrogen or carbon monoxide, formic acid vapor, plasma and ultrasonic methods will be introduced. Fluxless soldering by laser is described in Chapter 12.
13.3.1 Gaseous reduction Among reducing gases hydrogen (H2) is used occasionally to remove tin oxide from metal surfaces. Hydrogen reduces and dissolves solder oxides, and also reduces copper oxides in soldering temperature. Although hydrogen is used for high Pb-solder around 350 °C, at normal soldering temperatures about 220 °C~260 °C, this reaction is slow and inefficient. This causes
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difficulty in appling hydrogen to fluxless soldering. The initiation temperature to reduce solder oxide by hydrogen is given in Table 13.1 [6]. A Si-wafer coated with Sn/Au can be bonded to the Au-plated Si-substrate in N2+H2 atmosphere without flux [7]. An In/Ag-coated chip was also bonded to an Au-plated substrate in an H2 atmosphere for a fluxless process [8].
13.3.2 Acid vapor fluxless soldering In the activated acid vapor process, a dilute solution of a vaporized acid with an inert carrier gas is used for removing oxide layer instead of flux [3, 10]. The acid vapor carried on the metal surface reacts with the surface oxide, and removes the oxide. Formic [3] and acetic [9] acids which have low boiling temperatures are promising candidates for the acid vapor. The reaction between oxide and gaseous formic acid is as follows [3]: When T > 150 °C MO + 2HCOOH = M(COOH)2 + H2O (M : metal ) When T > 200 °C M(COOH)2 = M + CO2 + H2 H2 + MO = M + H2O Figure 13.2 illustrates an example of soldering equipment using formic acid vapor. [3] The formic acid vapor is supplied to the controlled atmosphere chamber with dry nitrogen gas, and fluxless soldering occurs.
13.3.3 Plasma assisted dry soldering (PADS) The PADS process converts surface oxides of solder to new compounds and oxyfluorides, which leads to fluxless soldering [10-12]. In this process fluorinecontaining gas such as CF4 or SF6 is dissociated in a RF-generated plasma or Table 13.1 Initiation temperature for hydrogen to reduce solder oxide [6] Solder
Dominant type of oxide
Initiation temperature (°C)
63Sn/37Pb 96.5Sn/3.5Ag 99.3Sn/0.7Cu 95Sn/5Sb 48Sn/52In 91Sn/9Zn
Tin oxides Tin oxides Tin oxides Tin oxides Indium oxides Zinc oxides
~430 ~430 ~430 ~430 ~470 ~510
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N2 Sample
Flowmeter HCOOH Heater
Computer Sensor
N2 tank Hotplate Formic acid
Reflow chamber
13.2 Schematic of formic acid vapor soldering equipment. [3]
microwave to form an active gas of fluorine atoms [10]. The produced fluorine atoms react with solder oxides as below: SnOx + yF = SnOxFy
(1)
The PADS process can be carried out under an inert gas such as N2 or even in air. The solder converted by the PADS process can be stored for a week in an air or for two weeks in an inert gas. The converted film breaks during soldering to expose melted fresh solder, and joining is performed. Figure 13.3 shows an example of the PADS equipment and schematics of chamber structure [13].
13.3.4 Fluxless soldering under argon or hydrogen plasma Plasma treatment using the energetic particle from a glow discharge is one of the alternative methods to the application of flux [14]. Plasma can remove the oxides or organic contamination through physical and chemical action (see Fig. 13.4). As a dry cleaning process, plasma treatment produces little or no waste and is therefore far more attractive than solvent or acid-based techniques where gallons of waste may be generated. Economically, the plasma process is not as costly as one would presume, considering that it is expected to eliminate environmental concerns associated with solvent cleaning, cleaning apparatus, flux and flux application systems [15]. Conventional plasma treatment includes argon or hydrogen plasma [16– 19]. The active hydrogen ion is generated in a microwave plasma system with the samples downstream from the plasma. Reflow must be in situ because oxide would re-grow even at room temperature if any metal surface is exposed to air. The addition of hydrogen to argon plasma gas is helpful to WPNL2204
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13.3 PADS equipment and schematic of chamber structure [13].
H C
Ar
O
H H
Contaminant
Oxide layer Solder
13.4 Illustration of hydrogen and argon plasma reaction on solder surface.
lower the oxidation rate but cannot ultimately stop it. If Ar gas is ionized in the plasma, Ar+ can enhance the formation of monatomic hydrogen (see following chemical equation) which in turn improves the reduction of oxides. Ar+ + H2 = ArH + H0
13.3.5 Plasma cleaning and soldering using an Ar and H2 gas mixture Ar and H2 gas mixture is more environmentally-friendly but less harmful to the passivation layer of the Si die. The plasma gas typically consists of H2 gas within 10 vol.% and Ar. Figure 13.5 shows a typical RF plasma etching system used to etch the oxide of a solder surface. Though it is the vacuum chamber type, the atmospheric plasma system is also available today.
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RF generator
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13.5 RF Plasma system (a) schematic of RF plasma system (b) an example of a RF plasma etcher.
Figure 13.6 shows the surface of an Sn-3.5%Ag bump after different treatment. In Fig. 13.6(a), the wrinkled oxide layer covering a solder surface is shown. The bump was reflowed without flux to have the oxide layer grow artificially. Figure 13.6(b) is the surface of the bump reflowed with flux. The surface became smooth after the oxide layer was removed chemically by flux. Figure 13.6(c) is the surface of the bump treated by Ar + H2 plasma. The surface became rough due to the partial etching of the solder surface as well as oxide layer by physical and chemical reaction with the plasma. Auger depth profile analysis can be used to estimate the effect of plasma treatment on surface oxide etching characteristics. Figure 13.7 shows the AES results before and after the Ar+10%H2 plasma treatment to the Sn
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13.6 Sn-3.5wt%Ag Solder bump surface after different treatment (a) after reflow on hot plate at 230 °C for 2 min without flux (b) after reflow with flux on hotplate at 250 °C for 10s (c) after Ar + 10%H2 plasma treatment at 50W for 60s.
solder disk [20]. The figure shows the sputtering time at which oxygen is not detected. The sputtering time for oxide removal is 3.1 min before plasma treatment. After Ar+10%H2 plasma treatment, this time decreased to 1 min. The sputtered oxide thickness is proportional to the sputtering time. Therefore, it is fair to say that the oxide layer became around 3.1 times thinner for Sn by the plasma treatment. Figure 13.8 is the AES result of the Sn solder after Ar+H2 plasma treatment under different H2 content in operating gas from 0 to 20% [20]. H2-added plasma treatment was more effective for oxide layer etching than Ar-plasma alone. But the etching characteristics of the Ar + H2 plasma were similar in 10~20%H2 content. Therefore, about 10%H2 addition to Ar gas seems to be sufficient to remove the surface oxide effectively. Masahiko et al. have reported the effectiveness of oxide etching of H2-added Ar-plasma in their application for BGA and CSP package [17]. Mannos et al. reported the effect of the chamber pressure on the plasma etching characteristic [21]. Figure 13.9 shows the effect of pressure on ion bombardment energy. At very low pressure (below 10 mtorr), high ion energy bombardment (physical sputtering) is dominant. As the pressure increases, the neutral particle density increases with decrease in ion energy. Ion-assisted etching which combines the effects of both physical and chemical etching, takes place. At higher pressures, above 100 mtorr, chemical etching by neutral particles plays an important role. A thin Sn oxide layer can be destroyed by mechanical hydrostatic force when the molten solder exists under the oxide layer [22]. This is called ‘ice floe’ theory, which can explain the reason why soldering is easy with the presence of some metallic Sn under the oxide layer [23]. The thin oxide of a solder surface is easy to break by volume expansion when melted in an inert gas atmosphere of an ordinary belt furnace [18]. However, the plasma treatment should be performed just before the flip chip bonding to minimize the growth of the natural oxide on solder bump surface.
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C O Sn(O) Sn
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13.7 AES depth profile of Sn solder after plasma before and after plasma treatment (the etch rate is 7 nm/min for SiO2) (a) before plasma treatment (b) after Ar + 10%H2 plasma treatment at 50W for 60s [20].
When the plasma treated solder surface is exposed to air, the natural oxide starts to grow on the solder surface. Therefore, it is important to know how long the Ar + H2 plasma treatment is effective in air. Figure 13.10 shows the AES depth profile results of plasma-treated Sn solder according to exposure
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70 No treatment Ar Ar + 5% H2 Ar + 10% H2 Ar + 15% H2 Ar + 20% H2
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13.8 Oxygen concentration of Sn surface etched by Ar + H2 plasma with different H2 content (power: 50W, time: 30s, pressure: 270 mtorr) [20].
Ion energy
Physical sputtering
Ion-assisted etching
Chemical etching
0.001
0.01
0.1 1.0 Pressure (torr)
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13.9 Effect of pressure on ion bombardment energy [21].
time [20]. The oxide layer grows just after the plasma-etched solder is exposed to ambient conditions. The thickness of the oxide layer from a 3-day specimen is similar to that from a 1-day specimen. The oxide layer of a 1-day specimen is definitely thicker than that of a 2-hour specimen. Therefore, the minimization of the
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Oxygen concentration (atomic%)
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3.5
13.10 Oxygen concentration of Sn surface according to time in ambient after Ar + 10%H2 plasma treatment at 50W for 30s. (Pressure: 270 mtorr) [20].
time span between plasma treatment and bonding will lead to a sound solder joint. Currently, research on atmospheric plasma is very active [24, 25]. The application of atmosphere plasma is expected to play a major role in the transition from the batch plasma treatment process to the more productive continuous process for fluxless soldering.
13.3.6 Fluxless solder ball bumping by plasma reflow In a fine pitch flip chip package, the residual flux may exist in the region where the inspection and removal of the residue is almost impossible. Therefore, fluxless flip chip soldering is becoming an active research area. The shape of Sn-3.5Ag solder bumps reflowed in different conditions is compared in Figure 13.11. In laser fluxless soldering (Fig. 13.11(a)) the abrupt local laser radiation made the ball bonded partially with the UBM somewhat mis-aligned from the center. The joint between solder ball and UBM had a sharp notch at which stress may be concentrated under shear loading. On top of the bump, there was a crater originating from rapid cooling. The crater may cause a void at a flip chip solder joint if gas is trapped in it. The crater will not appear if the laser current or pulse width increases. However, at high current or long pulse width, the frequency of capillary clogging also increases, so that the life of the capillary may be reduced. Therefore, after initial bonding at low laser current, reflow of the ball is required to reshape the bump and to increase adhesion strength with UBM.
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13.11 Sn-3.5wt%Ag solder bump shape reflowed in different condition (a) as-bonded solder ball by laser, (b) reflowed with flux, (c) reflowed in Ar + 10%H2 plasma.
The shape of the bump reflowed on a hot plate with flux to reshape it is shown in Fig. 13.11(b). As was expected, the bump was reshaped so that the crater from laser soldering disappeared. The solder wetted well on the UBM, and the notch at the joint also disappeared. However, there were flux residues around the solder bump, so adequate flux cleaning was required. The residual flux may affect the reliability of the flip chip package due to the negative effect on underfill materials [26]. In the case of an optoelectrotic device package, the laser signal may be deflected or attenuated by flux residues and impact the performance of the components [27]. Solder bumping by plasma is also illustrated in Fig. 13.11(c). The Sn3.5Ag ball was reflowed in Ar+10%H2 plasma at 100 W for 60 s. The figure shows that the crater from laser bonding disappeared by bump reshaping and solder wetting on the UBM. Therefore, Ar + H2 plasma is quite effective for a fluxless reflow process. During Ar + H2 plasma reflow, self-alignment occurs as shown in Fig. 13.12 [28]. The off-centered solder ball about 30 µm from UBM began to reshape at 40 s (Fig. 13.12(b)) and completely selfaligned to the UBM pad at 60 s (Fig. 13.12(c)). However, excessive reflow time can cause spattering around the bump (Fig. 13.12(d)). The shear strength of a bump is affected by plasma reflow time and plasma power (see Fig. 13.13 [28]). At a plasma power of 100 W, the shear strength reached a maximum of 85 gf after 100 s and decreases thereafter. At a power of 200 W, the strength reached a maximum at 20 s, and decreased thereafter. The shear strength of the bump reflowed on a hot plate at 250 °C with flux is also plotted in Fig. 13.13. The strength was kept 60~80 gf which is a similar value to plasma reflow. The mechanism of Ar + H2 plasma reflow can be explained as follows (see Fig. 13.14). An As-bumped solder ball is locally bonded by laser with UBM and surrounded by an oxide layer with a crater on top as in Fig. 13.14(a). When the ball is reflowed in Ar + H2 plasma, the oxide layer on solder and UBM surface is broken by Ar+ ion bombardment. The momentum energy of Ar+ ion is transferred to heat energy by collision to the solder
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(a)
(b)
(c)
(d)
13.12 Top view of Sn-3.5wt%Ag solder bump reflowed in Ar + 10%H2 plasma at 100W (a) 20s (b) 40s (c) 60s (d) 120s [28].
0
10
Hot plate reflow time (s) 20 30 40
50
60
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Ar + 10% H2 100W Ar + 10% H2 200W Reflow with flux at 250°C
0 –10 0
20
40 60 80 Plasma reflow time (s)
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13.13 Sn-3.5wt%Ag solder bump shear strength on Cu/Ni UBM with different reflow condition [28].
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UBM Si-wafer (b)
(b)
IMC
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(d)
13.14 Mechanism of Ar + H2 plasma reflow (a) As-bonded solder ball by laser ball bonding (b) oxide layer is removed by Ar + ion bombardment (c) reshaping and self-alignment begins by liquid surface tension (d) molten solder reacts with UBM and IMC formed [28].
bump and the solder is heated to melting point (Fig. 13.14(b)). The reducing atmosphere of H2 plasma prevents molten solder from re-oxidation. The reshaping and self-alignment begins by surface tension of liquid solder (Fig. 13.14(c)). The molten solder reacted with the UBM and IMC formed (Fig. 13.14(d)). The optimum plasma reflow condition will depend upon plasma systems due to plasma complexity. Therefore, more systematic research should be done to enlarge the process window of plasma reflow as a flux-free process. It is said that an electronic device can be damaged by plasma, and thus methods to avoid this should be prepared. The plasma reflow can be applied to optoelectronics and MEMS packages where the use of flux is restricted.
13.3.7 Fluxless soldering using ultrasonic energy In flow soldering, ultrasonic wave can be a method of fluxless soldering. The ultrasonic energy in the solder bath can remove oxides and other contaminants from metallic surfaces [29]. Aluminum and stainless steel, which require very active and corrosive fluxes, can be joined by ultrasonic soldering. Mechanical scrubbing has been used to break up oxides and assist flip chip
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bonding. Some success has been achieved using ultrasonic in traditional thermosonic bonding mode [30-36]. Figure 13.15 shows a typical ultrasonic flip chip bonding system which consists of ultrasonic generator, transducer and horn. Substrate temperature, bonding load, and ultrasonic power and time are major parameters to optimize the ultrasonic flip chip bonding process. The principle of ultrasonic bonding is not completely understood but can be explained as follows. The chip undergoes strong vibratory motion on the substrates which are contacted by bonding load. The surface oxides and contaminants at the bonding interface are destroyed and removed by the vibratory motion. This motion reveals fresh solder and UBM (Under Bump Metallurgy) metal surface, and two fresh surfaces are bonded [37]. The excessive bonding load will suppress the amplitude of ultrasonic vibration and the relative movement of two workpieces. Similar results are reported for the thermosonic bonding of Au bumps [32, 36]. The extended bonding time causes bonded joints to fail from fatigue. The shear strength of Sn-3.5wt%Ag solder bump flip chip is affected by bonding parameters (see Fig. 13.16 [37]). The shear strength increased proportional to bonding temperature. The die shear strength of the chip bonded by ultrasonic wave below the melting temperature of Sn-3.5Ag solder (221 °C) is quite high. Figure 13.16(b) shows the die shear strength with the bonding load at 25 W ultrasonic power. The die shear strength increased as high as 50 gf under the bonding load of 0.8 N/bump but decreased above 1.0 N/bump. Figure 13.16(c) displays the die shear strength variation with ultrasonic power, which means that strong friction improved the bond strength. The die shear strength increased with ultrasonic power. As one of the fluxless soldering methods, ultrasonic bonding has an advantage to bond a chip at low temperature. Thus, ultrasonic bonding is useful for a flip chip device which is sensitive to high operation temperatures. However, during ultrasonic bonding the bonding area remains in the solid state [38], and self-alignment does not occur in flip chip bonding. Thus, in the electronics industry more precise placement of workpieces such as chip and substrate are required to achieve a good bond. Transducer Load Load Bumped die Ultrasonic generator Horn
Substrate Hot plate
13.15 Schematic diagram of ultrasonic bonding system.
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13.16 Effect of ultrasonic bonding parameters on shear strength per bump from die shear test of Sn-3.5wt%Ag bump flip chip (a) bonding temperature (bonding load; 0.8N/bump, ultrasonic power; 25W), (b) ultrasonic power (bonding load; 0.8N/bump), (c) bonding load (ultrasonic power; 25W) [37].
13.3.8 Other processes for fluxless soldering Vacuum atmosphere was applied for fluxless soldering in electronics packaging [39]. The vacuum soldering reduces trapped gas volumes, and prevents oxidation of solder and parts to be joined through low residual partial gas pressures. However, in vacuum soldering, combination with reducing gas is required to overcome obstacles to achieve ideal wetting. Fluxless soldering using Sn-Au solder which could be achieved by Snlayer plated with Au was suggested [40, 41]. The Au-layer prevents the oxidation of Sn-layer, and reacts with Sn to make a eutectic reaction at
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217 °C. The sample to be bonded is heated under air, H2 or H2+N2 atmospheres. Other various solder systems for fluxless soldering were suggested by this research group. This can be a promising process for fluxless soldering, but metal coating is required as a solder. Fluxless soldering with rolled solder in N2 atmosphere was applied successfully to bond glass and Si-wafer [16]. In some case, plasma cleaning is employed after the first reflow for this process. A back-light-unit (BLU) for LCD (liquid crystal display) was dip soldered without flux using ultrasonic equipment. A metal cap of BLU was fluxlessly soldered on a glass tube [42]. Ultrasonic soldering at ambient temperature for bonding of Si-wafer to PCBpad was performed [43]. It is emportant to control bonding load and time to get a sound joint for this process.
13.4
Summary
For fluxless soldering formic acid vapor or H2 were used before or in the 1980s. However, this cannot promise great success to achieve fluxless soldering. In the 1990s plasma assisted dry soldering (PADS) was introduced to remove tin oxides, and this process works well for a Pb-Sn solder. However, fluorine can etch SiN or SiO2, and RF power may damage IC devices. Eutectic reactions such as in Sn-Au, Ag-In, Au-In, etc., systems were developed for fluxless soldering in the 1990s and have continued to be studied. This can be a promising process for fluxless soldering although metal coating is required as a solder. Plasma fluxless soldering under Ar + H2 was suggested in the 2000s, and this seems to be successful. However, as mentioned above concerns remain about RF power. The rolled solder which was applied to bond glass and Siwafer without flux in N2 atmosphere seems also to be successful. However, in some cases plasma cleaning is needed, and the limitations of preformed solder can be a drawback. Ultrasonic dip soldering has been applied for fluxless soldering. In this case excessive dissolution from the part to be joined is to be avoided. Most fluxless soldering processes mentioned above still have some limitations for application but these problems will gradually be solved in future.
13.5
References
1 S. C. Britton and K. Bright, Metallurgia, Vol. 56, p. 163 (1957). 2 J. L. Jellison, J. Golden, D. R. Frear, F. M. Hosking, D. M. Keicher and F. G. Yost, ‘Advanced Soldering Process’ (in The Mechanics of Solder Alloy Wetting and Spreading ed. by F. G. Yost, F. M. Hosking and D. R. Frear, Van Nostrand Reinhold, New York, p. 230 (1993). 3 W. Lin and Y. C. Lee, ‘Study of Fluxless Soldering Using Formic Acid vapor’, IEEE-Transactions on Advanced Packaging, Vol. 22, No. 4, pp. 592–601 (1999).
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4 M. J. Molina and F. S. Rowland. Nature, Vol. 249 p. 810 (1974). 5 J. H. Lau ed., Flip Chip Technologies, McGraw-Hill, New York, p. 83, (1995). 6 C. C. Dong, A. Schwarz and D. V. Roth, Feasibility of fluxless reflow of lead-free solders in hydrogen and forming gas, NEPCON 97, Malaysia, (1997). 7 C. C. Lee and R. W. Chuang, IEEE Trans. CPT, Vol. 26, No. 2, pp. 416–422 (2003). 8 Y. Chen, W. So and C. Lee, IEEE Trans. CPMT, Vol. 20, No. 1, pp. 46–51 (1997). 9 H. Matsuki, H. Matsui, and E. Watanabe, ‘Fluxless Bump Reflow Using Carboxylic Acid’, Proceeding of 2001 Int’l Symposium on Advanced Packaging Materials, pp. 135–139 (2001). 10 N. Koopman, S. Bobbio, S. Nangalia, J. Bousaba, and B. Piekarski, ‘Fluxless Soldering in Air and Nitrogen’, 1993 IEEE-ECTC Conference Proceedings, pp. 595–605, (1993). 11 S. Nangalia, N. Koopman, V. Rogers, M. W. Beranek, H. E. Hager, E. A. Ledbury, V. A. Loebs, E. C. Miao, C. H. Tang, C. A. Pico, E. J. Swenson, D. Hatzis, P. Li and C. Luck, ‘Fluxless, No Clean Assembly of Optoelectronic Device with PADS’, 1997 IEEE-ECTC Conference Proceedings, pp. 755–762 (1997). 12 S. Nangalia, P. Deane, S. Bonafed, A. Huffman, C. Statler, and C. L. Rinne, ‘Issues with Fine Pitch Bumping and Assembly’, 2000 Int’l Symposium on Advanced Packaging Materials Proceedings, pp. 118–123 (2000). 13 www.coe.uncc.edu/~smbobbio/fluxless/fluxless_intro.html 14 G. Takyi, N. N. Ekere and J. D. Philpott, ‘Solderability Testing in Nitrogen Atmosphere of Plasma Treated HASL Finish PCBs For Fluxless Soldering’, 1998 Electronic Components and Technology Conference Proceeding, pp. 172–179 (1998). 15 G. Takyi, Doctorial Thesis, Department of Aeronautical, Mechanical and Manufacturing Engineering, University of Salford, p. 3 (1998). 16 C. B. Park, S. M. Hong, J. P. Jung, C. S. Kang and Y. E. Shin, ‘A Study on the Fluxless Soldering of Si-wafer/Glass Substrate Using Sn-3.5mass%Ag and Sn37mass%Pb Solder’, Materials Transactions, Vol. 42, No. 5 pp. 820–824 (2001). 17 F. Masahiko, M. Tsugunori, D. Kazuhide and N. Hiroshi, 1999 IEEE-ECTC Conference Proceedings, pp. 408–414 (1999). 18 T. Nishikawa, M. Ijuin, and R. Satoh, ‘Fluxless Soldering Process Technology’, 1994 IEEE-ECTC Conference Proceeding, pp. 286–292 (1994). 19 K. Tango, J. Shibata, N. Hosada and T. Suga, ‘Bonding at Room Temperature and Fluxless Reflow of Copper and Lead-free Solder’, Advances in Electronic Packaging, EEP-Vol. 19–1, pp. 1053–1057 (1997). 20 S. M. Hong, C. S. Kang and J. P. Jung, Jour. of Elec. Mater., Vol. 31, No. 10, (2002). 21 D. M. Mannos and D. L. Flamm, Plasma Etching An Introduction, Academic Press Inc., London, 1988. 22 R. J. Wassink, Soldering in Electronics, Electrochemical Publication, Ayr, Scotland, p. 226, (1989). 23 P. E. Davis, M. E. Warwick, P. J. Kay and S. J. Muckett, ‘Intermellacic Compound Growth and Solderability Part 2: Reflowed coatings’, Plating and Surface Finishing, Vol. 69, pp. 72–76 (1982). 24 A.-A. H. Mohamed, R. Block, K. H. Schoenbach, Proc. of 2001 Pulsed Power Plasma Science, IEEE, p. 393 (2001). 25 E. S. Paller, J. E. Scharer, R. Cao and K. E. Kelly, Proc. of 2001 Pulsed Power Plasma Science, IEEE, p. 393 (2001). 26 C. Beddingfield and L. M. Higgins, ‘The Effects of Flux Materials on the Moisture Sensitivity and Reliability of Flip-Chip-on-Board Assemblies’, IEEE Transactions on CPMT Part C, Vol. 21, pp. 189–195 (1998).
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27 J. H. Lau ed., ibid, pp. 83–121. 28 S. M. Hong, C. S. Kang and J. P. Jung, IEEE Trans. on Adv. Pack., Vol. 27, No. 1, pp. 90–96 (2004). 29 R. W. Woodgate, The Handbook of Machine Soldering-SMT and TH, 3rd edn. John Wiley & Sons, New York, p. 306, (1996). 30 H. R. Faridi, J. H. Devletian and H. P. Le, ‘A New Look at Flux-Free Ultrasonic Soldering’, Welding Journal September, pp. 41–45 (2000). 31 S. Y. Kang, P. M. Williams, T. S. McLaren and Y. C. Lee, ‘Studies of thermosonic bonding for flip-chip assembly’, Materials Chemistry and Physics, Vol. 42, pp. 31– 37 (1995). 32 S. Y. Kang, P. M. Williams and Y. C. Lee, ‘Modeling and Experimental Studies on Thermosonic Flip-Chip Bonding’, IEEE Transactions on CPMT, Part B, Vol. 18, No. 4, pp. 728–733 (1995). 33 Q. Tan, W. Zhang, B. Schaible and L. J. Bond, ‘Thermosonic Flip-Chip Bonding Using Longitudinal Ultrasonic Vibration’, IEEE Transactions on CPMT, Part B, Vol. 21, No. 1, pp. 53–58 (1998). 34 T. Tomioka, I. Mori and K. Atsumi, ‘Bondability of Thermosonic Flip Chip Bonding’, Proceedings of 6th Symposium on Microjoining and Assembly Technology in Electronics, pp. 163–168 (2000). 35 M. Hashimoto, T. Yonezawo, and K. Higashi, ‘Development of Flip-Chip Bonding Technology Using Ultrasonic Vibration’, Proceedings of 6th Symposium on Microjoining and Assembly Technology in Electronics, pp. 175–178 (2000). 36 R. Kajiwara, M. Koizumi and T. Morita, ‘Ultrasonic Flip Chip Bonding Technology for LSI Chip with High Pin Counts’, Proceedings of 7th Symposium on Microjoining and Assembly Technology in Electronics, pp. 161–166 (2001). 37 S. M. Hong, C. S. Kang and J. P. Jung, Mater. Trans., Vol. 43, No. 6, pp. 1336–1340 (2002). 38 J. E. Krzanowski, ‘A Transmission Electro Microscopy Study of Ultrasonic Wire Bonding’, IEEE Transactions on CPMT, Vol. 13, No. 1, pp. 176–181 (1990). 39 G. Hagen, K. Wolter and A. Wagner, Vacuum soldering in electronics packaging, SMTA International Conference, Sept. (2002). 40 C. C. Lee, C. Y. Wang and G. S. Matijasevic, ‘A new bonding technology using gold and tin multilayer composite structures’, IEEE Trans. Compo., Hyb., and Manufac. Tech., Vol. 14, Issue 2, pp. 407–412 (1991). 41 R. W. Chuang, D. Kim, J. Park and C. C. Lee, ‘A fluxless process of producing tinrich gold tin joints in air’, IEEE Trans. Compo. And Pack. Tech., 27, pp. 177–181, (2003). 42 J. S. Cheon and J. P. Jung, Danyang Soltec Report, (not published). 43 J. M. Kim, J. P. Jung, Y. N. Zhou, Ultrasonic soldering of Si-wafer for flip chip with Sn-3.5Ag, SMTA International Conf. on Soldering and Reliability, April, (2007), Toronto, Canada (accepted for publication in J. Elec. Mater., 2008).
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14 Laser microwelding I M I Y A M O T O , Osaka University, Japan and G A K N O R O V S K Y, Sandia National Laboratories, USA
14.1
Introduction
The physical phenomenon behind light amplification by stimulated emission of radiation was originally described by A. Einstein.1 Progress from the equations to actual beams of coherent photons took ~40 years,2 and though the laser was briefly characterized as “a solution looking for a problem,” it has certainly found many problems to solve in almost 50 years since then. Examples of laser technologies which have become commonplace include (but are not limited to) materials processing, metrology, chemical analysis, chemical catalysis, communications, entertainment, data storage and recovery, holography, guidance systems, various medical procedures and product recognition. The subject covered in this chapter is laser welding, and in particular, laser microwelding. The history of laser microwelding has its roots in the history of laser welding itself. In the early days of laser technology, the beam power available was quite limited. In order to create fusion, good focusing optics and pulsed waveforms were needed in order to achieve adequate laser intensity. Needless to say, with only a relatively small amount of power, the applications which could be attempted were effectively micro-applications. Some application examples listed in an early text3 include: spot welding of small thermocouples, and the hermetic closure of a small detonator. Other early references4–6 refer to similar activities involving joining to thin films, joining ribbons and small diameter wires, and joining electrical interconnects to an early transistor near glass feed-throughs. These pioneering applications exhibited features which still cause industry to turn to lasers: the ability to accomplish a fusion joint near a heat sensitive material and the ability to join very small features, without needing physical contact. As increased powers became available, the size range of workpieces which could be joined grew, and while laser microjoining never disappeared, it became less significant to overall industry developments. A continual trend towards miniaturization over the years has nevertheless always been a source 345 WPNL2204
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of new applications for this precisely-controllable, precisely-directed, intense heat source. A more contemporary application is to “field program,” configurable microcircuits by linking (or cutting) interconnections. This can take place in vertical (i.e. different layers) or lateral (same layer) geometries.7 This capability, combined with laser cutting, makes in-situ programming a much more flexible process, opening up a broader range of circuit customization options. As noted in an earlier chapter, the term microwelding is somewhat elusive. Many references to microwelding really imply meso- or milli-scale welds. Adhering to the definition of Chapter 3, laser microwelding concerns the joining of parts having at least one dimension <100 µm. Accordingly, the focused laser spot size needs to be approximately this size or smaller. Achieving such a small focus diameter has become much easier with the introduction of near-diffraction limited industrial lasers in recent years. In the past, extensive intra-cavity aperturing and multi-element lenses of very short focal length would have been needed. Remaining limits to achieving small diameter focused spot size include lens aberrations and ultimately the diffraction limit set by the laser’s wavelength. However, it is now possible to readily achieve a spot diameter on the order of 20 µm. Joint geometries possible via laser microwelding are only limited by access of the laser beam and by fixturing considerations. Thus, butt, lap, corner, edge, fillet, and T-joints and the cross-wire geometry for wires are all possible. However, <100 µm thick material is definitely quite “floppy” and hard to maintain in consistent thermal contact. Unless specialized techniques are employed (such as “laser spike” lap welding, see Section 14.8.2), joint gaps must be kept consistently small relative to the beam size and material thickness. Typically this is accomplished by fixtures in contact with the parts (which somewhat negates the non-contact nature of the laser heat source). Even after a successful joint is produced, the assembly will still be fragile and must be handled carefully. Laser technologies of interest to microwelding, important characteristics of laser beams, the beam-workpiece interaction (including material and thermal flow), characterization and control of the process and the welds resulting will be introduced in this chapter. Defects encountered in practice will also be noted. Further, suggestions for fixturing and part manipulation will be given, along with recommendations for laser power, wavelength, pulse duration and frequency, culminating in recommendations for what is needed in a laser microwelding system. Illustrative examples will also be provided, though some of them will violate our <100 µm criterion for microwelding where the principal at issue is well-illustrated and no better example is available.
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Fundamentals of laser microwelding
14.2.1 Laser beam propagation and focusing The usefulness of a laser beam as an intense source of heat derives from its coherence, the ability of the photons which make up the beam to align such that their electromagnetic fields sum to a large amplitude, and do so consistently over an appreciable length of the beam path, rather than partially cancelling because of random misalignments in time and space, as is the case with noncoherent light sources. This coherence is related to the extreme monochromaticity (narrow bandwidth) of the light. Coherence may be expressed as either a length or a time (with the two related by the speed of light). A second feature of laser light which makes it precise enough for microwelding is that it originates as a collimated (nearly parallel) beam with minimal divergence θ, the angle between the envelope of the beam and its axis (some references define divergence as twice this value). The radius w of a Gaussian beam (defined later in this section) propagating in space along the z-axis is described by the following equation (where λ is the laser wavelength): w = w0[1 + {(z – z0)λ/(π w 02 )}2]1/2
(1)
The minimum radius w0 is located at z0 which is inside the laser cavity. w0 is often referred to as the minimum waist. The magnitude of z – z0 relative to π w 02 /λ will define if one is in the near field ({z – z0} << π w 02 /λ) or far field ({z – z0} >> π w 02 /λ) of the beam. In the far field, which is the most useful region for laser welding: θ ~ Kλ/d
(2)
where K is ~1 (it ranges from 2/π for a Gaussian beam to 1.22 for a uniform beam) and d is the diameter of the limiting aperture from which the laser beam emerges, closely approximated by 2w0. The θw0 product of the laser beam is essentially unchanged by passage through an optical element. Thus if the beam is expanded to a larger value of w0, its divergence must decrease. The importance of low divergence lies in the ability of a lens of focal length F to focus a Gaussian beam to a minimum spot radius rs (more about this shortly) as given by the relation: rs = Fθ
(3)
Combining equations 14.2 and 14.3 gives for a focusing lens, assuming that the beam fills the entire lens (where d in equation 14.2 is the lens diameter D): rs = Fλ /D
(4)
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F/D is the definition of the F number of a lens. Practical lenses approach an F number moderately greater than 1 at best, hence the rule of thumb that the focusing limit of a beam is on the order of its wavelength (the so-called “diffraction-limited” condition). In reality, spherical aberration of the lens will add substantially to rs as the F number decreases. A final factor which is important to laser processing is the depth of focus, the distance along the propagation direction where the intensity of the focused beam remains essentially constant and equal to that at the focal point. The distance along the beam path δz at which the intensity decreases by a factor of 1/ξ2 is derived from equation 14.1 as:8 δz = ± (π w 02 /λ)(ξ2 – 1)1/2
(5)
If the location where ξ = 1.05 is chosen (corresponding to ~10% intensity decrease), δz = ± w 02 /λ
(6)
Two important derived parameters which combine the laser’s power with the beam size or divergence are the beam radiance (or brightness) and the beam intensity. The radiance is the power per unit area per steradian, while the intensity is the power per unit area. The radiance is more often applied to measuring a laser’s output and is of great interest to the laser designer. On the other hand, the intensity of the laser energy delivered to the surface of the parts being welded controls the melt behavior, so the latter parameter is the more useful of the two for the welding engineer. It is conventional to describe the radius: rs of a Gaussian beam by the radial location at which the intensity has decreased to 1/e2 of the peak intensity I0. This corresponds to a condition in which 86% of the beam’s total power P is located within the circle of radius rs. An axisymmetric Gaussian beam has an intensity distribution given by: I(r) = I0 exp[–2(r/rs)2]
(7)
Integrating this intensity distribution over space to determine P and then rearranging gives: I0 = 2P/π rs2
(8)
From equation 14.8 it is clear that the peak intensity is a strong function of the focused spot size. Real beams from high power lasers typically have multiple resonance modes and thus are not Gaussian. They may be elliptical, rectangular or even irregular in shape; nevertheless, to describe their “size”, the diameter of a circle of area equal to that contained within the measured intensity contour line corresponding to 1/e2 of the beam’s peak intensity is often reported. Non-Gaussian beams containing high order modes propagate according to the equation:
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349
(9) 2
which would be equivalent to equation 14.1 if M = 1. The factor M (sometimes referred to as the “times diffraction limit” factor9) has become a common quantitative measure of a laser’s beam quality. One can estimate the value of M2 needed to achieve a small enough spot size for microwelding (~10 µm) as a function of wavelength and focusing lens focal length from equation 14.4. Hence a non-diffraction-limited beam of a given M2 value will give a spot size of M2 times that predicted from equation 14.4, neglecting aberration. For a near-infrared (IR) laser of ~1 µm wavelength, with a typical focal length of 100 mm and lens diameter of 25 mm, in order to achieve a 10 µm radius spot, an M2 value of <2.5 is required. Alternatively, the required divergence from equation 14.3 can be stated: <0.1 mrad. Introduction of a fiber optic beam delivery system allows convenient delivery of laser energy to the workpiece. The fiber’s effect on the beam depends upon whether the fiber is single mode or multi mode, and whether the index of refraction of the fiber is stepped (different at the outer surface) or graded (variable throughout the fiber, typically in a parabolic manner). Fibers can be used to “homogenize” the intensity distribution of multi-mode beams, and for the case of laser microwelding, if small enough, can select for the Gaussian (TEM00) mode, which as we have seen, aids focusing to a small diameter. Since the power demands for laser microwelding are modest, small diameter fibers can be used without running into fiber damage due to excessive thermal heating or exceeding the fiber’s peak intensity damage threshold. Where fiber technology really aids laser microwelding is in the production of low-divergence laser beams of small waist, hence low θw0 product. The beam waist in a single mode fiber laser is essentially limited to the fiber diameter, which can be quite small (~15 µm) for lasers appropriate to microwelding. Much of the above discussion was condensed from Chapter 2 of Reference 3, to which the reader is directed for a more thorough description.
14.2.2 Laser-matter interaction Since the first laser light was produced by a doped ruby solid state laser in 1960,2 a variety of lasers have been developed employing a variety of lasing media (gas, liquid, solid and free electrons) and lasing at wavelengths ranging from X-rays to far-IR. Pulse durations range from femto-seconds (10–15s) to continuous wave (CW). Industrially-available lasers suitable for practical microwelding include: the flashlamp- or diode pumped solid state Nd:YAG laser (1.06 µm), various laser diodes (0.8–1.1 µm), the fiber laser (1.04–1.5 µm), the thin disk laser (1.03 µm) and the CO2 laser (10.6 µm). The laser-
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matter interaction is affected by the intensity (also known as power density), wavelength and dwell time of the laser beam. Unlike electrons (as in E-beam welding), an appreciable fraction of light impinging upon a surface is reflected. Various models exist which explain the dependence of the reflectivity (the fraction reflected) on the irradiated material’s properties.10–12 For metals, it is generally the case that at a given wavelength, the reflectivity will increase with the electrical conductivity. However, the reflectivity of a metal surface depends greatly upon the wavelength of the light incident, the angle of incidence, the surface roughness, the presence of surface films or layers, the composition of the material and its temperature. Figure 14.1 shows the absorptivity of various metals to a range of wavelengths (absorptivity ~[1 − reflectivity] for opaque materials like metals) 13. However, because of the multitude of effects which influence the reflectivity, the data shown in Fig. 14.1 can only be considered semiquantitatively indicative of the actual behavior which may be encountered in practice. For the intensity regime of interest to laser microwelding of metals, the mechanism of energy absorption involves damping of the incoming photons’ electromagnetic radiation field by the electron cloud associated with the lattice, which eventually converts the photons’ energy to increased lattice vibration indicative of localized heating.14 For a metallic target, the absorption occurs within a fraction of a wavelength distance from the surface (~100 nm) and is described by an exponential decrease in the electric field strength of the laser radiation with penetration distance below the surface.15 The effective depth of penetration is given by an attenuation coefficient
Absorptivity
CO2
0.8
Nd: YAG
SHGNd: YAG LD direct
1.0
0.6
0.4 Carbon steel 0.2
Al Polished Ag
0.0 0.2
0.4
Ni
Cu 0.6
1 2 4 Wavelength (µm)
6
8 10
20
14.1 Absorptivity vs wavelength for different metals. Wavelength of selected lasers are also indicated.
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whose value is wavelength dependent.16 For non-metals the laser penetration distance is typically at least 10x greater (or even >>10x for transparent materials). The effect of this localized surface heating varies from gentle heating to explosive ejection of ionized atoms. Figure 14.217 shows the range of physical behavior expected to result as a function of power density and pulse duration available to commonly-available lasers. Non-equilibrium effects where the electron and lattice temperatures differ appreciably,18 seen at greater intensities than those plotted on Fig. 14.2 (associated with ultrashort pulse durations) are not discussed here, as they are not relevant to welding. However, for the case of transparent materials, high laser intensities (associated with fs~ps pulse durations) can promote absorption as will be discussed in Section 14.8.5.
14.2.3 Laser-plume interaction Before the laser beam can reach and interact with the target material, it must pass through a laser-induced plume created above the surface. While the plume becomes more noticeable as the power reaching the surface increases, and is especially dramatic in keyhole welding, it is still present to a small extent in conduction welding. In the plume, phenomena such as absorption, scattering and refraction, all of which depend on laser wavelength, the welding velocity, the power density of the focused laser beam in free space (that
108
Plas
ma
prod
Power density (W/cm2)
ucti
107
Material removal No
10
on
va
po
6
riz
Hole drilling
ati
on
Cutting S vap urface oriz atio n
Small melted depth
105
No
We ldi
Heat treating 104 10–7
10–6
10–5 10–4 Pulse duration (seconds)
ng
me
10–3
ltin
g
Large heat affected area 10–2
14.2 Laser-material interaction phenomena expected as a function of pulse duration and beam power density (does not include ultrashort pulse durations).
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which would obtain with no specimen present) and the identity and density of the gaseous species present in the plume can affect the energy actually reaching the workpiece surface, as seen in Fig. 14.3. Beam absorption is due to the inverse Bremsstrahlung effect,19 in which electrons present in any plasma are accelerated by the electric field of the laser beam. The absorption coefficient of the laser beam in a plasma is proportional to the electron density and the square of the wavelength of the laser beam.20 The effects of absorption are important in CO2 laser welding where the temperature of the plume can reach as high as 10,000 K, since the plume is heated by absorption when it is ionized.20 The laser power impinging upon the workpiece is significantly reduced by the shielding effect of the absorbing plasma. Hence, in CO2 laser welding it is common to direct an assist gas such as He, Ar or N2 to the laser-plasma interaction point to reduce the electron density by cooling the laser-induced plasma. In Nd:YAG laser welding, on the other hand, the absorption by inverse Bremsstrahlung is negligible, since the absorption is approximately 100 times smaller than for the CO2 laser. While the plume temperature still does increase with increasing laser power, it is more typically <4000K. In contrast to absorption, scattering is more important at shorter wavelengths (the reason the sky is blue), since Rayleigh scattering is proportional to the inverse fourth power of the wavelength of the laser beam.21 In Nd:YAG laser Laser beam
Scattering
Absorption Refraction
Keyhole
14.3 Schematic illustration showing interaction between laser beam and laser-induced plume above keyhole.
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welding, because metal vapor condenses into particles (which act as scattering sites), significant laser power is lost by scattering. In welding at Nd:YAG and diode laser wavelengths, scattering becomes especially significant when the welding is done in a glove box to protect the workpiece from oxidation, which increases the amount and size of particles in the vapor present in the plume. Refraction can usually be neglected in laser microwelding, since the height of the plume is very small, thus the increase in the beam spot due to defocusing (sometimes called thermal “blooming”) is small relative to the focused beam spot, except for the tightly focused laser beam obtained from a single-mode fiber laser.22 Table 14.1 summarizes the importance of absorption, scattering and refraction for different lasers proposed for microwelding.
14.2.4 Classification of welding mode Laser welding of metals is generally classified into thermal conduction and keyhole (or deep penetration) welding modes.23,24 The mode present depends on the power density of the focused laser beam at the work surface, and the temperature distribution developed, as shown in Fig. 14.4. In keyhole welding, the metal surface is heated by the focused laser spot to a temperature where surface evaporation creates a recoil force sufficient to depress the molten surface into a cavity (the keyhole). While the term “keyhole” welding was originally coined to refer to a welding procedure where a through hole completely penetrating the workpiece was created by the heat source, (which then was translated along the seam, guaranteeing a full penetration weld joint), it is now used to refer to any process that produces a deep cavity, whether or not it completely penetrates. Conduction welding Opaque materials In laser welding of a material at low power densities the interior is heated by Table 14.1 Importance of absorption, scattering and refraction in laser welding
Semiconductor laser (λ = 0.8–1.1 µm) Nd:YAG laser (λ = 1.06 µm) Fiber laser (λ = 1.04 µm) CO2 laser (λ = 10.6 µm)
Absorption
Scattering
Refraction
0
2
0
0 0 2
1 1 0
0 1 1
0:small, 1:intermediate, 2:large
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3-D conduction
2-D conduction
Laser beam
Shallow hole
Keyhole Single reflection T < TBoil
A few reflections T ~ TBoil
Multiple reflections T ~ TBoil
(a)
(b)
(c)
14.4 Laser beam reflection and surface temperature of laser absorbing face in laser welding of (a) conduction mode, (b) intermediate mode, and (c) keyhole mode.
thermal conduction from the work surface as shown in Fig. 14.4(a). Analysis of the temperature distribution resulting is dealt with in more detail in Section 14.3. However, the minimum laser power needed to melt the material’s surface can be evaluated by a simple calculation. When a material is irradiated by a stationary Gaussian heat source, the surface temperature at the center of the heat source T(t) is given by:25 AW tan –1 4 T ( t ) = 23/2 d π Kd
α t
(10)
where d is diameter of laser spot evaluated at the 1/e-power point, W is the laser power, A is the absorptivity of the surface, K is the thermal conductivity and α is the thermal diffusivity. Thus the minimum laser power for melting metal of melting temperature Tm is given by Wm =
π KdTm A
(11)
The absorptivity is strongly dependent on the material being irradiated and the laser wavelength. Figure 14.1 shows the absorptivity of selected metals plotted vs. wavelength,13 with the wavelengths of commonly available welding lasers indicated. The laser absorptivity has a tendency to decrease with increasing wavelengths; indeed, at IR wavelengths the absorptivity is inversely proportional to the square root of the d.c. electrical conductivity. As may be
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readily noted, silver and copper have very low absorptivity for Nd:YAG and CO2 lasers. The absorptivity of copper is approximately 10% at 1.06 µm (Nd:YAG laser), whereas it increases to 50% at 532 nm, the second harmonic generation (SHG) wavelength for Nd:YAG. The absorptivity is also influenced by the surface conditions. When a metal surface is contaminated by oil or covered by an oxide film, the absorptivity is increased relative to that obtained for the clean metal. Figure 14.5 shows the minimum laser power needed to melt selected metals with fundamental and SHG Nd:YAG laser wavelengths at a focused spot diameter of 50 µm calculated with equation 14.11. It is seen that silver is the most difficult metal for conduction welding, requiring approximately 900 W at 1.06 µm. Copper requires nearly 300 W at 1.06 µm. The easiest materials are nickel and steel, with minimum laser powers for melting between 20 and 30W, which are 30 to 45 times smaller than for silver. Conduction welding of bulk metal is generally not efficient at infrared laser wavelengths, as a large fraction of the laser energy is lost by reflection (especially in silver and copper), and the laser energy actually absorbed by the work is dissipated 3-dimensionally. However, conduction welding has the advantages of being very reproducible and providing high surface quality, since the welding process is very stable. Copper can be welded at considerably lower power using a SHG Nd:YAG laser, since the absorptivity of copper increases to 50% at 532 nm. This is even higher than for aluminum as seen in Fig. 14.1. Figure 14.6 shows an example of copper welding with a SHG Nd:YAG laser.
400 1064 µm (fundamental Nd: YAG wavelength) 532 µm (second harmonic wavelength)
Laser power for melting (W)
350 300 250 200 150 100 50 0
Ag
Cu
Al
Ni
Steel
14.5 Minimum laser power for melting different metals with fundamental (1.06 µm) and SHG (0.532 µm) Nd:YAG laser (1/e-diameter: 50 µm).
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Bottom surface
Top surface
14.6 Butt welding of copper plates of thickness 80 µm by SHG Nd:YAG laser (pulse energy: 0.66 J at 2 kHz, welding velocity: 0.2 m/min) (Courtesy of Miyachi Technos).
Transparent materials In contrast to metals, laser beams deeply penetrate transparent materials such as polymers and glass, with very little laser energy being reflected or absorbed. There are two mechanisms for a transparent material to absorb laser energy for welding. The first mechanism employs an absorption process described mathematically by the same equation as for the metallic case, but with a smaller attenuation coefficient. In transmission welding a transmitting polymer overlays an absorbing polymer. The under layer material can be doped with a colorant to promote absorption of the laser beam at or near the interface. The upper layer is then heated by thermal conduction from physical contact with the laser-heated under layer to provide a lap weld at the interface, as shown in Fig. 14.7(a).26 Alternatively, a CO2 laser, whose longer wavelength radiation is absorbed strongly, can be used for lap welding a thin upper sheet to a thicker lower sheet. In this case, the laser is absorbed at the surface of the upper sheet and the lower sheet is heated by thermal conduction. The upper sheet is limited in thickness by the requirement that heat conduct through it to the lower layer. A second method of transparent material welding depends upon nonlinear absorption processes. When an ultrashort pulse duration laser (hence high power density) is sharply focused into a bulk glass, laser energy is absorbed in the beam waist region by multiphoton and avalanche ionization.27,28 In this case, the interior of the glass is melted without needing colorant, as shown in Fig. 14.7(b).
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Upper polymer plate
357
Laser beam
Upper glass plate
Lower glass plate Lower polymer plate with colorant (a)
(b)
14.7 Laser welding of transparent material by (a) linear absorption of colorant and (b) nonlinear absorption process of bulk material.
An advantage to welding of transparent material where melting is entirely confined at a subsurface interface is the lack of exterior surface contamination due to spattering or vapor redeposition from an exposed fusion zone. (This clearly does not apply to the case where CO2 radiation is used.) Conduction to keyhole transition welding When the power density of the focused beam spot is increased sufficiently to heat the work surface to temperatures where appreciable evaporation of the material begins to occur, the surface of the molten zone is depressed by the vapor recoil pressure. This depression results in a shallow cavity. Since the surface of this cavity is below the original surface of the material, the “surface heat source” laser can now be said to heat the metal interior, as is shown in Fig. 14.4(b). The resultant bead is no longer strictly conduction type, since its shape is no longer semicircular and the aspect ratio (depth/width) of the bead cross-section exceeds 0.5, which is the theoretically-expected aspect ratio for conduction-only heat transfer. A large fraction of the incident laser beam is still lost to reflection, since the beam path exhibits only a small number of reflections (barely more than once) in such a shallow hole. Figure 14.8 shows the transition in weld bead shape obtained by changing the laser power of a diode laser, from the transition type (a, b) to the keyhole type (c, d).29 The transition from conduction to keyhole mode welding is also observed by changing the focal position with respect to the work surface at constant laser power. It is interesting to contrast the conduction to keyhole mode transition behavior for laser welding vs. electron beam welding. A very sharp transition occurs in laser welding using a sharply-focused beam spot from a singlemode fiber laser, since the laser beam total energy absorption increases
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(a)
(b)
(c)
(d)
14.8 Cross-sections of weld bead and aspect ratio of diode laser welding in stainless steel at power densities (× 105 W/cm2) of (a) 0.6, (b) 0.72, (c) 1.08 and (d) 1.56.
dramatically because of the sharply increased number of reflections (each reflection adds another increment of energy) when a deep keyhole cavity is produced, as shown in Fig. 14.9(a).30 In contrast, for electron beam welding, the bead changes gradually, as shown in Fig. 14.9(b),30 since no change in energy absorption occurs by the formation of the keyhole. Laser welding near the bead transition region should be avoided, since the penetration depth is sensitively affected by minor changes in the focus position with respect to the work surface, surface reflectivity, and fluctuations in the beam power caused by plume effects. The laser power reflected from the keyhole depends on the aspect ratio of the keyhole and laser wavelength, since the reflected power depends on the number of reflections24,31,32 and the reflectivity of the molten metal. This indicates that the reflection from the keyhole provides useful information concerning the penetration depth, which may be used in a feedback control scheme as will be described in Section 14.5.2. Keyhole welding A further increase in the power density creates a deep keyhole with approximately the same diameter as the focused beam as is seen in Fig. 14.4(c). With such a deep keyhole cavity, most of the laser energy is absorbed by multiple reflections in the keyhole. In this manner, the laser, despite its “surface heat source” nature, can directly heat locations deeply below the original surface of the workpiece. In CW laser welding, the keyhole is maintained by a dynamic balance33 between the surface tension pressure acting to close the (assumed essentially cylindrical) keyhole: Pγ = γ/r
(12)
and the evaporation pressure acting to expand it: Pvap = mnu2
(13)
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df
Bead width (µm)
60
25W (λ = 1.06 µm) 450 mm/s 40 µm in thickness
359
df
db
40 (df + db)/2 20
db 0
–100
0
100
200
Focal position ∆f (µm) (a)
Penetration depth (mm)
7 6 5
Electron beam welding (150kV–70mA v = 500 mm/min
4 3 2 1 0
–100
–50 0 Focus position (mm)
50
(b)
14.9 Bead transition in laser welding and electron beam welding (a) fiber laser welding; (b) electron beam welding of stainless steel made at different focal positions. (Note differing horizontal scales!).
where γ is the surface tension of the molten material, r is the radius of the keyhole, m is the mass of material evaporated, n is the density of the evaporated mass, and u is the velocity of the evaporating material. In Fig. 14.10 the Pγ and Pvap are plotted vs keyhole radius, and there are two points, N and S, where Pγ = Pvap is satisfied.33 The keyhole is unstable at N, because the keyhole radius is driven by the pressure difference |Pγ – Pabl| so as to move away from N when the keyhole radius is slightly perturbed due to changes in metal flow, assist gas pressure or absorbed energy. As the result, the keyhole either disappears or jumps to S. Conversely, the keyhole is maintained stably at S; the pressure difference |Pγ – Pabl| works as a restoring force to move back to S, should the keyhole radius be perturbed to larger or smaller values
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Pγ
Pressure
Pvap
N S
Radius of keyhole
a
14.10 Ablation pressure Pvap and surface tension Pγ plotted vs radius of keyhole.
by the same types of perturbations noted earlier. Thus the keyhole oscillates around S at a resonant frequency, assuming that the perturbation is not very large. The resonant frequency of the keyhole oscillation decreases with increasing keyhole size. (Detection of this oscillation frequency may provide useful in-process monitoring information.) In the above development it is assumed that the surface tension force will always act to close the keyhole. However, this may not be the case for all geometries. Assuming that a cylindrical keyhole of radius r is formed in a cylindrical molten pool with radius R in foil with thickness h, the surface area of the molten region S (including the top, bottom and cylindrical surface of the keyhole) is given by: S = 2π(R2 – r2 + rh)
(14)
Then the increase in the surface area due to an infinitesimal increase in the keyhole radius, dS/dr, is given by34 dS/dr = 2π(h – 2r)
(15)
Equation (14.15) indicates dS/dr > 0 when h > 2r, which describes the relatively deep, small cavity of the previous section. The positive sign shows that a decrease of the keyhole radius will result in a decrease in the surface energy. Thus a keyhole radius decrease is favored, until and unless the surface tension is balanced by the evaporation pressure. On the contrary, in the case of h < 2r, a relatively wide, shallow cavity, dS/dr < 0. This indicates that an increase in keyhole radius leads to a decrease of the surface energy, suggesting that an increase in keyhole radius will not be resisted, resulting in keyhole that will not close, resulting in a weld with a lack of fusion or bridging. Laser beam energy is absorbed in the keyhole by multiple reflections.
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Therefore the absorptance in the keyhole is dependent on the reflectivity of the molten metal32 and the aspect ratio of the keyhole, which determines the number of reflections in the keyhole. Figure 14.11 shows the calculated ratio of the reflected laser power plotted against the aspect ratio of the keyhole. It is seen the reflection loss becomes less than 10% at aspect ratios larger than 5 in Nd:YAG laser welding and 9 in CO2 laser welding.31 In keyhole welding, the surface of the keyhole is kept at the evaporation temperature of the material by the heat balance between laser heating and cooling due to the latent heat of evaporation, conduction and convection. Although evaporation plays an important role in producing and maintaining the keyhole, the latent heat of evaporation is very small, less than 0.5% of incident laser power, since only a small fraction of evaporated mass is ejected from the keyhole; most is re-deposited elsewhere along the keyhole wall.35 In pulsed laser welding, the keyhole formation is transient and dynamic as shown in Fig. 14.12. During the ramp-up period of the laser pulse, the molten metal flows out of the keyhole and forms raised areas on the top and bottom surfaces. When the ramp-up of the laser pulse is too fast, the molten metal splashes out to the surrounding areas as spatter, so that a concave bead is produced due to the lost molten metal. In the ramp-down period of the laser pulse, the molten metal flows back into the keyhole space to fill it up (assuming it hasn’t been spattered away). When the ramp-down time is too short for the cavity to be smoothly refilled by the molten metal, pores are left in the molten metal. Pores can be driven to the surface by buoyancy forces,
Degree of reflected laser power in (%)
70 Wavelength 10.6 µm Wavelength 1.06 µm
60 50
Keyhole depth t: 0.5 to 12.0 mm Keyhole spread diameter d: 0.47 mm Inclination angle front: 88° Inclination angle side: 88.5° Inclination angle rear: 88° Keyhole shift: 10% of depth
40 30 20 10 0 0
5
10 15 Aspect ratio(s)
20
14.11 Calculated reflected laser power vs keyhole aspect ratio (s = depth/width) and laser wavelength.
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Laser beam
Molten metal
Molten metal Keyhole
Melt flow Recoil pressure Ramp up period
Ramp down period
14.12 Keyhole behavior and molten metal flow in ramp up and ramp down periods in pulsed laser welding.
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when the ramp down time is long enough to allow them to reach the surface before solidification freezes them in place.
14.3
Thermal modeling of laser microwelding
14.3.1 Levels of complexity As noted in the previous section, laser energy is introduced into opaque material essentially at the surface, or at least within a surface layer only a few atomic layers thick.15 Thermal conduction into the material from this surface heating is well understood and easily modeled both analytically and via finite difference and finite element methods. Radiative and convective losses from the surface are somewhat more complex, but are still capable of solution. However, this simple behavior becomes more complex when the energy transport occurs at a high enough rate to cause melting and evaporation. Creation of a plume or a plasma above the workpiece-beam impingement point can cause changes in the energy distribution of the laser beam reaching the surface. The three mechanisms noted in Section 14.2.3: thermal defocusing due to refraction or “blooming,”22 scattering by small condensed vapor particles,21 and at long wavelengths, absorption by ionization18–20 can cause significant redistribution of the beam energy before it even reaches the surface. Since these effects tend to be transient and somewhat chaotic, modeling the energy input to the surface becomes complex.36 At the other side of the beam-workpiece impingement surface, as the solid material melts, convective transport within the fusion zone becomes important, even predominant, with forces due to buoyancy, aerodynamic drag, surface tension (including a contribution from the Marangoni force due to the temperature dependence of the liquid–vapor surface energy) and eventually, with the onset of evaporation, a vapor recoil force coming into play (see Chapter 3 for more details). As noted earlier, the vapor recoil force has a key role in explaining the phenomenon of “keyholing”, i.e. the creation of a deep, tunnel-like cavity in the fusion zone which provides exceptional depth/width ratio welds. With a change in the impingement surface from a plane into a depression and then a dynamically-varying keyhole, effects of multiple reflection from the surface then need to be added to the already complex story. The history of modeling of laser welding37–41 has traced the deepening understanding of the complexity of the beam-material interaction. In microwelding, keyholing may be unnecessary, as the thickness of the parts of interest is quite small. Avoiding keyholing removes the necessity of modeling fusion zone surface displacement, redistribution of energy due to reflection, and significant changes in energy deposition with depth, relative to the simple assumption of surface energy deposition only. Nevertheless, keyholing is of extreme practical interest as, due to the keyhole’s higher
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energy absorption, it simultaneously provides maximal weld depth/width ratios with minimal heat input (and hence minimal distortion) at high travel speeds. Furthermore, the smaller, narrower fusion zone is less prone to liquid instabilities which can result in weld defects as was seen on pages 359–360. The reduced distortion is especially advantageous in lap geometries, where less comprehensive tooling is needed to keep parts in close contact. (For spot welds a keyhole technique known as laser “spike” welding42 described in Section 14.8.2 is useful for dealing with variable gaps.) When making butt welds, keyhole welds do require exceptional beam aim point accuracy to locate the weld at the faying surface. This is a critical feature of all butt geometry laser welds, but is especially critical in microwelds. Thus, while conduction mode laser microwelds are certainly possible, they pay a heavy penalty in productivity and in fixture complexity to combat distortion. A further simplification is that essentially all laser microwelding is autogenous, i.e. there will be no filler metal introduced, such as occurs in GMAW or other wire-fed processes. The only exceptions to this rule known to the authors are droplet welding, discussed in Section 14.8.1, laser-assisted ultrasonic wire bonding,43 which will not be discussed in this chapter, as it is not a fusion process, and laser-assisted chemical vapor deposition.44 The latter is not a fusion welding process either, but rather a decomposition or chemical reaction, and likewise will not be treated here. The effect of part dimensionality is another concern which can lead to either simplification or complication. For example, the relative thinness of most microsize sheet materials, combined with normal metallic values of thermal conductivity, makes a 2-D assumption for simplicity almost a given. The other side of the coin is that details of heat sinking, surface-related heat transfer (the surface to volume ratio of microparts is typically much larger than for macroparts) and the effect of nearby edges may seriously degrade the accuracy of analytic 2-D solutions. Another factor is that since 2-D solutions typically involve much lower thermal gradients than 3-D, the regions where radiation and convection are applicable (i.e. high temperature regions) are expanded. Nevertheless, it is almost always useful for engineering purposes to try the simplest mathematical solution possible. Table 14.2 summarizes some of the analytical tools which may be found useful, before resorting to numerical models.
14.3.2 Conduction model predictions for microwelds Line and point source calculations were made using an available software package (SOAR)45 to approximate weld sizes and thermal cooling gradients for a range of weld power and travel speed. These are collected in Table 14.3. In addition to the temperatures, the software calculates the melting efficiency. The process efficiency (related to the reflectivity of the material
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Table 14.2 Selected analytical solutions of increasing complexity Model
Equation (see Section 14.11)
Applicability
Comments
Stationary point source
1
Spot welds
Semi-infinite solid
2
Keyhole spot welds
2D, full penetration
3
Conduction weld
Semi-infinite solid
Extended surface sources Gaussian Top-hat (square)
Conduction weld 4 5
Semi-infinite solid, better approximation for HAZ temperatures
Moving line source
6
Keyhole
2D, full penetration
Extended line sources Cylindrical Parallelepiped
Keyhole 7 8
2D, full penetration, better approximation for HAZ temperatures
Partial penetration
Semi-infinite solid, partial
keyhole or transition conduction/keyhole
penetration, better HAZ temperature approximation
Extended volume Double ellipsoidal
9
Laser microwelding
Stationary line source Moving point source
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Heat input (W)
Travel speed (mm/s)
dT/dx* (°C/mm)
100 (70)22 25 10 10 10 10 1 1 1 1
500 500 (310†)22 500 100 100 50 50 5 1 100
1.7 2.5 1.6 3.2 4× 1.6 4× 3.6 3× 4×
× 103 × 104 × 105 × 104 104 × 104 105 × 105 105 105
–dT/dt** [–dT/dx+] (°C/s) 1.9 3.0 3.2 1.2 2.4 3.0 1.2 1.0 2.4 2.4
× × × × × × × × × ×
106 107 108 107 106 106 107 106 105 107
[8.5 × 105] [1.2 × 107] [8.0 × 107] [3.2 × 106] [4.0 × 106] [8.0 × 105] [2.0 × 107] [1.8 × 106] [3 × 105] [4 × 107]
Weld width (mm)
Weld length (mm)
Melt efficiency
0.106 0.018 0.0024 0.0117 Depth:0.03 0.023 Depth:0.003 Depth: 0.003 Depth: 0.003 Depth: 0.003
0.381 0.024 0.0023 0.012 0.06 0.024 0.006 0.007 0.006 0.006
0.47 0.32 0.1 0.1 PS: 0.1 0.1 PS:0.007 PS:0.0008 PS:0.0002 PS: 0.013
*Behind pool, estimated from calculated centerline locations for 1454 and 1354 °C contours **Behind pool, estimated from travel speed × Tmp/weld length49 + Behind pool, estimated from travel speed × dT/dx * †0.08 mm thickness PS: point source in finite thickness solution, not full penetration. 1W as a line source was insufficient to induce melting at any travel speed.
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Table 14.3 Line and point source calculations (data) for various welding parameters, 304 SS, 0.1 mm thick
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being irradiated) is not included in these calculations. Since one can either use a point or line source model for a given power and travel speed, it is interesting to compare the results. If insufficient energy is given to support through-thickness melting, the line source is inappropriate. When both models give a non-zero melt volume, one must decide which is more correct. This conflict is resolved by calculating an approximate power density value at the surface (a spot size estimate is needed). If the calculated power density >5 kW/mm2, then a keyhole-like 2-D line source is more appropriate.46 For example, see Table 14.3, rows 4 and 5: 10 W @ 100 mm/s. If one assumes a 10 µm spot diameter, the power density is 130 kW/mm2, a keyhole scenario. If instead the spot size is 100 µm in diameter, the power density drops 100fold, and it becomes a conduction mode weld. (A few actual data values where full penetration was achieved in 100 µm (or comparable) thickness of 304SS are included in the table for comparison.) The final column of Table 14.3 shows the calculated melting efficiency (defined as the ratio of the enthalpy needed to melt (only) the fusion zone to the amount of enthalpy entering the part). As is typical of all fusion welding processes,47,48 the melting efficiency varies from near zero to about ~0.5. It is noticeable that at any given heat input, the efficiency tends to increase with increasing travel speed. This is not to say that the weld size will increase with increasing travel speed. Changing from the 2-D line source, keyhole mode to the 3-D point source, conduction mode also tends to reduce the melting efficiency at equivalent power and travel speed.
14.3.3 Metallurgical consequences A simple estimate of the cooling rate at the back of the weld fusion zone is given by:49 dT/dt = travel speed × Tmp/weld length
(16)
where we have replaced in the original expression the length from the source to the back of the pool by the total length of the pool, and the temperature rise from ambient to melting with the melting temperature (liquidus) itself. Alternatively, one can multiply columns 2 and 3 in Table 14.3, which gives agreement within a factor of ~3. The cooling rate can then be used to estimate the size of dendritic or cellular microstructural features resulting from solidification from the following:50 λ = a (ε)–n
(17)
where a, n are constants equal to 80, 0.33 for primary arms and 25, 0.28 for secondary arms, respectively, in grade 310 austenitic stainless steel. Measured dendrite arm/cell sizes for austenitic stainless alloys at similar estimates of cooling rates range from 0.4–1.1 µm.51
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While the cooling rate is maximal at the location chosen, and hence estimates the finest feature size, it has been noted that this will underestimate the expected feature sizes by only about two times given the fractional power law dependence on cooling rate. Because of the extremely rapid cooling rates encountered in microwelding, one would expect that solidification mode changes from slow rate macroscopic welds would be expected in stainless steels. In particular, it has been noted that single phase solidification structures result at high cooling rates, rather than mixed δ ferrite (a body-centered cubic crystal structure) and γ austenite (a face-centered cubic crystal structure) structures that would be expected at macro size weld cooling rates.52 Thus, the range of compositions exhibiting completely austenitic mode solidification will expand with rapid solidification. Unfortunately, solidification (hot) cracking tends to accompany this mode of solidification. Adequate weldability materials will need to be specified, either by ensuring extremely low tramp element content, or by adjusting the allowable composition range to favor ferritic solidification.53
14.4
Weld defects
Weld defects can be broadly classified into two categories: those associated with material problems, and those associated with continuity of the joint.
14.4.1 Materials defects The first category of defect is related to the changes which are undergone in microstructure when a solid metal is melted and re-solidified. For example, LIGA* material is extremely high strength by virtue of the extremely fine microstructure left by the electroplating process which produces it. When welded, a completely annealed, grain coarsened structure is left behind, creating a metallurgical notch effect due to the now undermatching weld strength. Other aspects are due to the extreme surface to volume ratio of micro fusion zones, which lead to sensitivity to contamination, either pre-existing or introduced by less-than-perfect shielding. Particularly for reactive materials such as Ti, Nb, etc., the high levels of O, C, N that may be nearly unavoidable at the microscale may lead to extremely brittle materials. In other cases, such as with LIGA-based materials, residual hydrogen may lead to porosity, or *LIGA: LIthographie, Galvanoformung, Abformung, acronym for high aspect ratio (depth to width), micron-level resolution, metallic part production process using synchrotron radiation to expose a masked photoresist material, followed by dissolution of the exposed photoresist, electroplating the developed cavities, dissolution of the mold and then chemically-etched release from the substrate.
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elements added to the plating bath may cause problems. Figure 14.13 shows an example of a LIGA-produced material which was laser spot welded and exhibited practically every type of metallurgical weld defect known to metallurgy: porosity, hot cracking and cold cracking, simultaneously. The problem was traced to the addition of saccharine (a stress-reliever) to the plating bath.54 Dissimilar materials can introduce a number of defects brought about by their lack of matching physical properties. One of the most serious involves the formation of intermetallic compounds in the fusion zone, which tend to be brittle.
14.4.2 Development process defects Gross discontinuity type defects such as lack of fusion (Fig. 14.14) or decohesion, excessive spatter and drilling (Fig. 14.15), burn-through (Fig. 14.16), excessive oxidation (Fig. 14.17) and cracks (Fig. 14.13) are generally resolved early during process development. Once these issues are solved by fixture and process parameter adjustments, what are left are more subtle issues, such as pores due to unstable welding conditions (created by dynamic detachment of portions of the keyhole, aka spiking, or root porosity), geometrical notches such as underfill and humping and thermal distortion. These will be seen in the next section. They tend to be more difficult to resolve, sometimes needing process monitoring and feedback control, which will be treated in Sections 14.5.1 and 14.5.2.
100 µm
14.13 Example of laser weld in Ni-Fe LIGA material showing porosity, hot cracking and cold cracking.
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0.5 mm
14.14 Example of LIGA pure Ni, showing lack of fusion in laser weld.
14.15 Example of laser micro spot fusion zones exhibiting spatter and drilling.
14.4.3 Manufacturing process defects and their avoidance In laser welding, the keyhole formed in the molten pool is not hydrodynamically stable. This instability can lead to weld defects when uncontrolled perturbations due to material and process variations occur. In this section, the formation and suppression of weld defects in keyhole welding are described. Underfilled bead There are two ways to produce underfilled beads. The first is encountered in pulsed laser welding when molten metal is splashed out of the keyhole
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14.16 Example of laser micro weld in lap geometry showing burn through of top layer.
14.17 Example of laser weld in pure Ni (LIGA) showing excessive oxidation.
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because the pulse shape is too steep during the ramp-up period, as is schematically described in Fig. 14.12. The underfilled bead is simply produced by a lack of molten metal in the weld bead. An example is given in Fig. 14.18(a).55 This defect can be prevented by prolonging the ramp-up time or by a dual beam mode, which combines a pulsed laser with a CW laser, whose result is shown in Fig. 14.18(b). The other type of underfilled bead defect is observed when the gap is too large in lap or butt welding, even in CW welding, as shown in Fig. 14.18(c). This defect is prevented by reducing the gap. Porosity Figure 14.19 shows the penetration depth plotted against defocus distance (distance from sharp focus) for various laser pulse lengths in spot welding of 316 stainless steel. The bead and porosity shapes and locations are also schematically noted vs parameter space. It is noticeable that deeply penetrating partial penetration welds are most prone to this problem, and within the scope of the study, porosity-free welds beyond about 0.75mm were not achievable with the pulse shape used.56 As noted earlier, an excessive pulse ramp-up slope can lead to melt ejection; however, when the ramp-down period is too steep, smooth metal flow needed to refill the keyhole cavity is prevented, the keyhole sides collapse, and bubbles are retained in the molten pool. With the rapid cooling usually present with laser welding, pores are left in the weld when the rapidly solidifying pool traps pores before they can float to the surface. However, it is possible to tailor the pulse shape to allow the pores sufficient time to reach the surface, as shown in Fig. 14.20.56 In pulsed laser welding, it is clear that defects (underfill due to spatter and pores) result when the pulse conditions are inappropriate. Porosity is much more likely to be produced when the bead does not penetrate completely through the plate. When the pool fully penetrates the
(a)
(b)
(c)
14.18 Underfilled bead in pulse laser welding (a) underfill caused by excess energy, (b) underfill is overcome by adding CW power, and (c) underfill in lap joint due to appreciable gap.
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Power, P (kW)
Laser microwelding 6
τ = 1 ms: Eo = 4.9 J/p τ = 2 ms: Eo = 10.1 J/p τ = 3 ms: Eo = 15.3 J/p τ = 5 ms: Eo = 24.2 J/p τ = 10 ms: Eo = 45.2 J/p
4 2 0
373
0
5
10 Time, t (ms)
15
20
1.5
Penetration depth, dp (mm)
SUS 316
: τ = 10 ms : τ = 5 ms : τ = 3 ms : τ = 2 ms : τ = 1 ms
1.0
Porosity free area 0.5
0
0
5
10 15 Defocused distance, fd (mm)
20
Power, P (kW)
14.19 Weld bead and porosity produced with different defocus distances and pulse widths in 316 stainless steel.
6 4 2 0
(a) (b) (c) (c) 0
10
τ = 10 ms: Eo = 38.7 J/p τ = 12 ms: Eo = 43.0 J/p τ = 14 ms: Eo = 50.3 J/p τ = 16 ms: Eo = 57.2 J/p 20
30
Time, t (ms) a
b
c
d
14.20 Controlled pulse shape and cross-sections of resultant weld bead. Porosity floats towards to the surface due to buoyancy forces, as the pulse tail is prolonged.
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back surface, any pressure at the bottom of the pool is relieved, and surface tension forces also tend to fill the cavity. Pores are also produced at the root of the keyhole due to spiking (rapid fluctuations to high intensity) as in electron beam welding, where the cooling time is too short for the bubbles to float to the surface. Humped bead When sufficient laser power is available, the maximum welding velocity in CW laser welding is limited by the onset of humping, shown in Fig. 14.21(b). Figure 14.21(a) schematically illustrates the top view of the keyhole surrounded by the molten pool, with the molten metal ahead flowing around the periphery of the translating keyhole. The molten metal ahead of the keyhole is transferred around the perimeter of the keyhole as it moves forward. The metal flow Keyhole Molten pool
d
D v
Vortex Flow speed V (b)
(a)
0.70 0.60
Velocity ratio V/v
0.50 0.40 0.30 0.20 0.10 0.00 0
2
4 6 8 Keyhole diameter d (mm) (c)
10
12
14.21 (a) Metal flow around the keyhole provides pressure drop at the keyhole and vortex behind the keyhole (b) appearance of humped bead (c) calculated flow velocity at keyhole perimeter.
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produces a pressure drop at the keyhole perimeter and a vortex flow behind the keyhole. Both phenomena combine to make the molten metal flow unstable,57 resulting in a bead which exhibits regularly arranged humps after solidification, as shown in Fig. 14.21b. The flow velocity V of the molten metal at the perimeter of the keyhole is approximated by:22 V=
dv D–d
(18)
where v is the welding velocity, d is the keyhole diameter (or focused spot size) and D is the width of the molten pool at the narrowest position of the melt flow. Figure 14.21(c) shows the flow velocity at the perimeter calculated for different keyhole sizes at a welding velocity of 1 m/s in stainless steel foil of 40 µm thickness. It should be noted that the size of the molten region is essentially constant with decreasing keyhole size, since the temperature field away from the heat source depends little on the heat source size. Equation 14.18 indicates that the flow velocity decreases with decreasing keyhole diameter d, while it is proportional to the welding velocity; thus humping can be suppressed by decreasing the keyhole diameter. Upper sheet
Lower sheet
100 µm
240 µm (b) Cross section
(a) Surface of upper sheet
Upper sheet
240 µm
Lower sheet
100 µm
(d) Cross section
(c) Surface of upper sheet
14.22 Surface appearance and cross-section of pulsed Nd:YAG laser welding with (a)–(b) narrower gap (unbridged) (c)–(d) wider gap (bridged).
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Non-bridged bead in lap welding Welding of thin material is a difficult task, since a non-bridged bead often results. This is because in principle the keyhole grows unstably when the material thickness is smaller than the keyhole diameter, as described on pages 359–360. It was recognized long ago that fixturing was critical for successful Gas Tungsten Arc (GTA) lap welding of thin materials58 where the upper sheet’s h < 2r, and this concept is equally valid for laser microwelding. Two general rules have been defined for laser lap welding of thin foil:59 (1) gaps in joint fitup should be less than 10% of the foil thickness, and (2) tooling spacing should be smaller than two times the thickness. Two attempts at lap welding of thin material with a pulsed Nd:YAG laser are now described, one initially unsuccessful, and the other successful. The first attempt involved spot lap welding of aluminum alloy A3003 of thicknesses 0.1 mm (top) and 1 mm (bottom) with a focused spot size of 0.15 mm.60 Figure 14.22 shows the surface appearance and cross-sections of bridged and non-bridged beads. It was found that bridging could not be reliably obtained, so a diagnostic study was performed to understand the problem and examine mitigative alternatives. Figure 14.23 shows measured values of the reflected laser beam intensity and thermal emission detected coaxially with the laser beam through dichroic mirrors for two welds, one successful and the other unsuccessful. The thermal emission increases until reflection from the lower sheet is detected (at ~2.5 ms), showing the upper sheet is perforated for the non-bridged bead. In the case of the bridged bead, the thermal emission increases continuously with little increase in the laser reflection, showing a common keyhole is produced for the upper and lower sheets. Note that in Fig. 14.22 the bridged weld has an apparently larger initial gap; however this weld was made employing an adaptive procedure (described in a later section) that increased the weld energy until bridging was obtained. The second attempt involved seam lap welding of extremely thin foil (down to 2.5 µm in thickness) to a much thicker washer, using a pulsed multimode Nd:YAG laser. A sophisticated laser welding technique was developed, where the surface of the washer was precision lapped before welding, and the laser beam was irradiated through a quartz window pressing against the foil, guaranteeing intimate contact between the foil and the washer, as shown in Fig. 14.24(a).59 Figure 14.24(b) shows the appearance of successful lap welds of 2.5 µm thick nickel foil to the thick stainless steel washer. (This is the minimum thickness material successfully welded reported in the open literature).
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Laser pulse duration: 4.8 ms 3
Power [µW]
2.5 2 1.5 1 0.5 0 0
1
2
3 Time [ms]
4
5
(a) Laser pulse duration: 4.8 ms 3
Power [µW]
2.5 2 1.5 1 0.5 0 0
1
2
3 Time [ms]
4
5
(b)
14.23 (a) Thermal emission and (b) reflected laser intensity from the lower sheet in pulsed Nd:YAG laser welding, which corresponds to Fig. 14.22 (䊉 = partially penetrated laser weld; 䊊 = non-bonded laser weld with hole).
14.5
Laser microwelding technologies
14.5.1 In-process monitoring Sound weld beads can nearly always be obtained in the laboratory by carefully optimizing welding conditions. In production, however, it is not easy to maintain optimal parameters of output power and mode structure of the laser beam, since lasers and peripheral optics can deteriorate in long-term operation.
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14.24 (a) Set up for lap welding thin foil to heavy section washer. Fully fixtured, ready to laser weld (b) SEM photograph of lap-welded sample of 2.5 µm-thick nickel foil to stainless washer at the top of the picture. The second weld pass located at the bottom of the picture was used to trim the foil to size.
Besides, the weld quality is also affected by non-laser conditions such as variations in input materials, gaps in butt or lap welding, contamination by oil and dust, oxidation of the work surface, and manufacturing concerns with throughput that might prevent the use of meticulous fixturing, even if the laser parameters are kept optimized. Thus it is desirable to evaluate the weld quality during the welding process by in-process monitoring. An early form of in-process monitoring of laser welding was demonstrated in 1984.61 Light emission detected by a photo sensor was used to monitor CO2 laser welding as shown in Fig. 14.25. When a laser-absorbing plasma was produced, due to inappropriate gas pressure or an incorrect direction of the assist gas nozzle, the weld bead became wider and shallower with accompanying strong light emission. Thus the assist-gas condition could be monitored by detecting excessive light emission from the welding zone. Since then, a variety of in-process monitoring techniques have been developed, including those which: detect light emission from the plume, detect laser beam reflected energy, detect thermal emission, perform video image analysis, measure sound pressure, etc. Figure 14.26 shows the relative usage of monitoring techniques reported (mainly) at ICALEO between 1999 and 2004. Detection of light emission from the plasma/plume is the most widely adopted because the phenomenon is well understood, and light emission can be acquired without significant noise in comparison with, for example, sound intensity signals. Light emissions can be acquired coaxially through a pinhole in the mirror in CO2 laser welding or through a dichroic mirror in Nd:YAG laser welding as is shown in Fig. 14.27.62 In-process monitoring techniques acquiring multiple signals are also being developed since more information is available. Figure 14.28 shows an example of a camera-based monitoring system to estimate the weld bead quality from the dimensions of the weld bead.63 In
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15
Light intensity from side (a.u.)
A
C
B
10
A Plasma bead
B
C
Sound bead
Humping bead
5
0
0
50 100 Assist gas pressure (cm of H2O)
150
14.25 Assist gas pressure vs light intensity measured from side direction in CO2 laser welding stainless steel (CW power of 1 kW).
Others
Sound
Plume radiation
Reflection
Image analysis
Thermal radiation
14.26 Relative frequency of monitoring techniques reported at ICALEO meetings, 1999–2004.
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Laser beam
CCD camera
CCD camera
Imaging optic
Imaging optic Pinhole mirror focusing optic
Workpiece vs (a)
CCD camera
Laser beam
Imaging optic
Focusing optic dichroitic mirror
Workpiece vs
Focusing optic dichroitic mirror
Workpiece vs
(b)
14.27 Schematic drawing of different sensor setups for CO2 and Nd:YAG laser systems.
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Laser beam
(c)
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Nd: YAG laser optics
CO2 laser optics
Laser microwelding
CCD camera Laser welding torch
381
Semiconductor laser Collimate lens Width Cylindrical lens Height Slit beam Step Welding bead
Angle θ Under cut
14.28 In-process monitoring system employing projected semiconductor laser stripe for inspection of weld bead and the detected image.
this system, the weld bead is illuminated by an inclined line-shaped semiconductor laser beam, and dimensions such as bead width, undercut, step (workpiece vertical misalignment) and the height of the bead reinforcement are evaluated by image analysis of the stripe produced by the semiconductor laser as recorded by a CCD camera mounted normal to the work surface. This system has been used in tailored blank welding at the Toyota Motor Co. Figure 14.29 shows another example of a camera-based monitoring system,62 where video images, as well as 3-D perspective views and centerline plots of the intensity distribution are shown. From these data records and analysis, one can recognize full or partial penetration. Frequency analysis of the monitored signal can provide useful information to evaluate the weld quality. Figure 14.30 shows an example evaluating underfilled vs correct bead contour from the frequency of the light intensity signal. The surface contour of the lap weld bead varies due to the gap between the sheets, becoming unacceptably underfilled at large gaps. Figure 14.30 shows the a.c. signal content of the frequency region between 4 and 6 kHz increases dramatically when the gap becomes larger than 0.3 mm. When the gap is <0.3 mm, a common keyhole is produced in the two plates, while if the gap is larger, two smaller keyholes which independently oscillate at higher frequencies are present.64
14.5.2 Adaptive control The next step to guaranteeing sound laser spot weld beads involves incorporation of real-time control of the laser parameters based on the signal detected by an in-process monitor. In the previous section, an example of multiple-signal monitoring of spot weld bridging of Al A3003 was given.60 At the time, a bridged weld was shown which had a large gap, though it was noted that an adaptive process
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Partial penetration
0
150 100 50 0 –1 –0.5
Vs
–0.5 –1 PL = 2 kW –1
–0.5 0 X (mm)
0.5
Signal (a.u.)
Y (mm)
0.5
0 X (mm) 0.5 –1
0.5 0 –0.5 Y (mm)
Intensity profile 250 200 150 100 50 0 –1 –0.5 0 Y (mm)
0.5
Full penetration
0
150 100 50 0 –1 –0.5
Vs
–0.5 –1 PL = 3 kW –1
–0.5 0 X (mm)
Signal (a.u.)
Y (mm)
0.5
0 –0.5 X (mm) 0.5 –1 Y (mm) 0
0.5
0.5
Intensity profile 250 200 150 100 50 0 –1 –0.5 0 Y (mm)
0.5
Material: ST14, 0.8 mm; laser power: 2–3 kW Nd: YAG; focus diameter: 600 µm; feed rate: 5.5 m/min; gas flow: He 15 l/min.
14.29 Camera-based in-process monitoring system for inspection of weld bead showing (a) the detected image, (b) perspective intensity plots and (c) centerline intensity plots for partial and full penetration conditions.
had been used. Figure 14.31 shows the flow chart of that adaptive controller, with an algorithm based on detected thermal radiation. An Nd:YAG laser for which pulse duration and peak power were electronically controllable was employed. The laser’s pulse width is controlled so as to control the time during which the radiation intensity exceeds 1.7 µW. Thus a perforated upper sheet (that is, an unbridged weld) is “repaired” by metal from a keyhole formed in the lower plate by increasing pulse power from 1.07 to 1.39 kW. Figure 14.22 (right) shows the surface and cross section of a laser microspot weld produced in A3003 alloy under adaptive control for hole repair, and Fig. 14.32 shows the effect on the weld’s shear strength due to implementation of the monitoring of heat radiation and adaptive control of the laser spot welder’s power. In lap welding of metal with high reflectivity such as copper and aluminum, it is difficult to maintain weld quality, since their reflectivity is so high that weld quality is sensitively affected by the surface condition, fixturing, and heat capacity of the weld part. Under such conditions, as just seen for Al A3003, adaptive control based on in-process monitoring signals can provide a solution to problems of variable weld quality.
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6
4 3 2 1 0 0
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14.30 Relationship between gap of lap joint of 0.8 mm–0.8 mm thickness and mean square of a.c. signal (v = 3 m/min, 3 kW).
An adaptive control system has also been developed for microwelding of copper thin sheets of 0.1 mm thickness using a system combining an acoustooptical Q-switched SHG Nd:YAG laser with a fundamental Nd:YAG laser, whose pulse width is controlled at 0.15 ms intervals during laser irradiation.65,66 This example was shown in Fig. 14.6. The objective was to consistently produce fully-penetrated welds. Welding of 200 samples resulted in incomplete melting of 7 samples without adaptive control. With adaptive control, no incomplete melting was detected for a second 200 sample lot. Incomplete melting was detected in both cases by a neural network reacting to reflected laser beam and thermal radiation monitoring signals.
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Set laser peak power at 1.07 kW YES Heat radiation level > 1.7 µW Re-joint YES
NO
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Hole check NO
Period of heat radiation level over 1.7 µW ≥ 1.5 ms YES
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14.31 Flow chart of adaptive control based on intensity level and period of heat radiation. Conventional control Adaptive control (stabilization of joint strength) Adaptive control (hole repair + stabilization of joint strength)
Shear strength of lap welded joint (N)
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14.32 Relationship between period of heat radiation power level over 1.7 mJ and shear strength of lap-welded joint made under adaptive control for hole repair in addition to data under conventional control.
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14.5.3 System for microwelding Enough aspects of laser microwelding technology have been presented that it is possible to note the features that would be desirable in an “ideal” laser microwelding system. Clearly, a near-diffraction limited laser and beam delivery system capable of producing focused laser spots of <100 µm diameter (perhaps as low as 10 µm) will be at the heart of the system. In order to take advantage of the small spot size, a motion control system, whether it moves the laser spot, the part, or both, should be of at least an order of magnitude better positional accuracy than the spot diameter. The motion control will need to be computercontrolled unless only spot welding is contemplated. Further, it is difficult or even impossible to make reliable microwelds at such size scales unless a high quality vision system of appropriate magnification is incorporated into the work station. Finally, depending upon the motion system’s velocity capability, an average laser power requirement may be estimated. The faster the motion system, the higher the power may be employed. Approximating a 100 µm diameter hemispherical fusion zone as the size needing to be melted (though at rapid speeds it will of course be elongated into a tear drop shape), the volume of such a shape is ~2 × 10–3 mm3 (the volume of a hemisphere of radius r is 2πr3/3). Melting metals requires ~10 J/mm3. Assuming that the melting efficiency is about 50%, and the laser absorptivity is 0.5, about 8 mJ of laser output would be needed. Further assuming that the fusion zone travel speed is given by X µm/s, the power required is estimated by calculating the time that a moving spot would take to traverse 100 µm (the dwell time = 2 × 100/X), multiplying by 2 (since the average power over the reference spot will increase from zero to 100% and then decrease back to zero, we approximate its average power as 50%, though this will not occur linearly) and dividing 8 mJ by that time. Table 14.4 summarizes the average power needed for a variety of travel speeds.
Table 14.4 Approximate laser power requirements for a microwelding system Travel speed
Power required
1 (µm/s) 10 100 1 (mm/s) 10 100 1 (m/s)
1.6 1.6 1.6 1.6 1.6 1.6 1.6
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It is clear that not much average power is needed for microwelding, even at fast speeds. At the slower travel speeds, assuming a melting efficiency of 50% is probably not justified, so the requirements should be adjusted upwards. Nevertheless, for practical purposes only a few tens of watts of average laser energy is needed. Relative to the accuracy of Table 14.4, one of the authors30 successfully demonstrated full penetration bead-on-plate fusion zones at speeds up to 500 mm/s in 0.3 mm thick 304 SS with a 200 W fiber laser. Based upon the above considerations, a typical laser microwelding system will include a near-diffraction limited laser of perhaps 50W maximum, an integrated computer controlled motion control system (with peripheral part loading/unloading if high throughput is desired), a high quality vision system, some to-be-determined form of sensor or sensors to monitor the process and implement adaptive control features, all packaged into a vibration-controlled, laser-safe containment, which would also provide an environment that prevents particulate contamination and filters any hazardous fumes created in welding. This fully automated hardware system would also require software which integrates the laser control, part motion control and quality control features. Acquisition cost of such a system should range upwards from about 150K USD (ca. 2007 USD), depending upon options incorporated. The template for such a system originates with the micromachining stations which have become available in recent years, with the substitution of cw or ms-pulse lasers instead of the Q-switched (or shorter pulse length) lasers that a micromachining station would employ. Because the low power welding lasers envisioned do not employ exotic gases (such as the fluorine-containing gas needed by some excimer lasers), are quite efficient (wall plug power is adequate, and bulky cooling systems are unnecessary), do not need specialty lenses or mirrors (such as the ZnSe lenses needed for far-IR lasers), their footprint, upkeep and maintenance costs will be minimal. Travel speeds for the motion control would need to be much faster than those typical of micromachining, however. Another facility which would be needed which is not yet well in hand is some form of tool which would measure the intensity distribution of the focused laser beam, and its power or energy (both CW and per pulse). While CW and averaged pulsed power measurements are not difficult, time resolved ultrashort pulse energy measurement instruments are still laboratory curiosities. Similarly, while mm size laser spots can have their intensities profiled, profiling tools for 10–100 µm size spots are not readily available, except by drilling small holes or measuring spot size at a distance from the focal spot. The system described above is envisioned to use a near-IR wavelength laser. This is suitable for welding most metals, opaque polymers and some glasses. With special techniques transparent materials may also be joined (laser-opaque material is typically introduced at the joint faying surface, or absorptivity modifiers added to the polymer matrix). By employing other
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wavelengths, a broader selection of transparent materials may be joined. In particular, fused silica is eminently weldable, via either ultraviolet (UV) or far IR to avoid its transparent wavelength region. Welding lasers are not readily available in the UV (where 10–100 ns pulse duration Q-switched, frequency multiplied Nd:YAG or gas-based excimers predominate); however, CO2 gas lasers at 10.6 µm wavelength are. While the diffraction limit for focusing is proportional to wavelength, and theoretically is about equal to the laser’s wavelength, in practice, the limit tends to be about an order of magnitude higher. Thus the 10.6 µm CO2 laser should be focusable to <100 µm diameter. Alternatively, it has recently been found that non-linear absorption mechanisms (see Section 14.8.5) encountered when using pico-second and shorter pulse lengths allow precisely controlled heating of unmodified glass.29 Since transparent materials are fundamental to the rapidly-growing technologies of photonics, microbiology and microanalytical chemistry (“lab on a chip”), a microjoining tool capable of welding them would seem useful. Thus important options for the above system would be to substitute alternative wavelength and/or pico-second pulse lasers.
14.6
Evaluation of microweld joints
14.6.1 Visualization Conceptually, once a microweld is made, it can be evaluated in much the same way as a macroweld; that is, it can be checked for hermeticity, strength, cosmetic appearance, and metallographic sections can be performed. However, the difference between concept and actual practice for objects thinner than the human hair is another matter altogether. Let’s start with the simplest and most widely used quality measure: visual observation of the surface of the weld. Assuming a video screen of perhaps 300 mm diagonal extent, as might be used with a typical laser welding system, the magnification needed to display a 100 µ feature at a convenient size, say 25 mm, would be x250. As anyone who has spent time looking at mm-size laser welds through a binocular stereo microscope can testify, observing a 3-D surface beyond about x50 is increasingly problematic because of difficulties in illumination, depth of field and ultimately, resolution. And that’s using the amazing pattern recognition capability of the trained human brain to interpret the visual clues and information, which information is greatly degraded on a video screen (even at HD resolution, should it become available). At the very least, a digital depth-of-field-enhancing optical microscope would be needed,67 and more likely an SEM. (As an aside, if you need to use the SEM for observation, then you might consider µE-beam welding as an alternative – see Chapter 15.)
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14.6.2 Metallography Metallographic methods, though slow, are certainly applicable to evaluation of microwelds. Exceptional care to section at the appropriate location (particularly if other than a transverse section of a seam weld is desired) and to prevent edge rounding would of course be necessary. Since the parts are typically quite small to start with, the metallographic personnel would have to have suitable handling equipment and a deft and delicate touch to avoid damaging or simply losing the parts, both before and after sectioning. Preencapsulation before sectioning would undoubtedly be employed.
14.6.3 Leak checking Leak checking is a very sensitive tool for detecting cracks and other throughweld defects (assuming that the geometry of the weld allows it), but it, too, has difficulties. Particularly if the weld being evaluated is a final closure weld, then there will be no way to attach a vacuum leak detector to it directly. Bomb methods would have to be used,68 which involve applying external pressure to infiltrate a marker gas (typically He) through any defects, and then placing the sealed object into an evacuable chamber attached to a mass spectrometer. Such techniques require both gross and fine leak detection to assure that the marker gas does not leak out too quickly for the post bomb checking step to be valid, and the volume of the object and the time at pressure are also important parameters. Very small volumes require very small leak rates to maintain the internal environment expected over long periods of time. Very small leak rates require extreme cleanliness and steps to ensure that surface desorption is not providing a “virtual” leak instead of an actual leak. Although, one could envision a process where the virtual leak rate associated with a partial-penetration crack or surface pore would aid in its discovery.
14.6.4 Mechanical testing Mechanical testing of 25 µm diameter wire welds has been accomplished for many years in the microelectronics industry. The technologies of pull- and/ or shear-testing of wire (and chip-to-substrate) bonds are quite well developed.69 Commercial equipment is readily available, sized to measure small load values, typically in the ~100 mN range. However, for geometries that are not amenable to this equipment, the difficulty level rises appreciably. For example, while a lap or fillet geometry has a “lip” which a shear tool can push against, a butt weld geometry often does not. It is possible to apply loads by using adhesives or by melting and solidifying wax or “fusible” metal (e.g.: Wood’s metal, a Bi-Pb-Sn-Cd eutectic) to attach the load train to the part to be tested.
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Such techniques would of course require some development, and would be more or less easily implemented depending upon the size of the part that needs to be tested. Should it be that both parts to be tested are small (unlike the wire bond scenario, where the substrate being bonded to is typically macroscopic, and easily held) the difficulty level rises dramatically. When the individual workpieces to be welded are fragile to begin with, joining multiple parts may improve their durability if the load paths are reinforced. However, it is also possible for the opposite to be true if the moment arm of external loading, such as occurs in handling, increases (e.g. the welded part is doubled in length), actually decreasing durability. Even the strongest microweld, by virtue of its size can only resist minute forces! For example, a spot weld of 100 µm diameter in a material of 690 MPa tensile strength, can only resist about 2.7 N of shear force. For weaker materials, and smaller welds, the forces are reduced commensurately. Needless to say, micromanipulation is an active area of robotics research, in which issues of locating parts to be manipulated, picking them up while avoiding damage and/or contamination, transporting them from initial to final location/ orientation, and then finally releasing them, are all under intense investigation. A similar, if not quite so mature, situation exists for hardness testing on small scales. Microhardness testing has of course been available for many years, and nanohardness testing has also become routine over the past decade. Both techniques are suitable for automation. However, for both techniques, the issue of specimen preparation is a critical one. Both techniques are intended for use on smooth or pre-polished surfaces. If used on an as-solidified surface, the technique chosen will need to result in indentations that are small relative to the surface roughness, which may be set by the overall geometry or by the nature of solidification. For pure materials, this can be an appreciable fraction of the molten zone width. For alloys (which solidify dendritically), the limitation will be given by the dendrite size. As noted in Section 14.3.3 on thermal gradients, this might range from 0.4 to 1.1 µ. Finally, the nanohardness technique, which samples a thinner layer than microhardness, is more likely to be compromised if surface contamination leads to differential hardness relative to the fusion zone interior.
14.6.5 Other techniques Other conventional techniques for defect assessment using X-rays, ultrasonic waves, dye penetrants, magnetic particles, and eddy currents are conceptually applicable, but again, their practical realization is not yet available. For example, synchrotron-generated X-rays are used to produce the molds from which LIGA parts are made, so one could envision their use in detecting defects of the same size scale. However, synchrotron X-ray sources are not yet common in industry. Microbeam X-ray techniques combined with computed
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tomography are just now being put into practice for mm-size welds; there is no reason why this trend may not be extended to smaller size scales once industrial demand is present. Ultrasonic techniques are limited by the wavelength of the sound waves used and by the ability to detect small changes in amplitude. Also, there must be a means of introducing and detecting the waves. Conventional contacttype ultrasonic sensors are too large; however, recent progress in laser-based non-contact methods have been applied to mm-size welds.70 Future developments may allow progress to smaller sizes. C-SAM (C-mode scanning acoustic microscope) instruments are presently available that use focusing sensors at up to ~GHz frequencies that are capable of detecting indications of a few microns size. These require immersion of the sample in a coupling medium (usually water), and though commercially available, are not exactly common in industry. Among other applications they have been used to examine defects in Ball Grid Array solder joints for the micropackaging industry.71 Fluorescent dye penetrants may be useful, particularly if they can be applied on one side and detected on the opposite side, as capillarity forces are quite effective at wicking the dye through small, tight defects, and optical detection of fluorescing dye is quite practical. If the standard PT approach of: (1) dye application, (2) superficial surface cleaning, (3) development, must be applied, the technique will be much less attractive. A combined leak detection and fluorescent dye approach could be quite powerful, to first detect the presence of a leak and then to localize it for further analysis. Conventional dry magnetic particle techniques are probably not useful, as most applications for microdevices need to be contaminant-free, and the particles would have to be guaranteed to not be pyrophoric at the small sizes needed. Also, it is hard to visualize how a completely particle-free specimen could be guaranteed after inspection, particularly after a magnetization field is applied. On the other hand, wet techniques, and especially ferro-fluids have been used in optical metallography to determine domains in magnetic materials and to elucidate delta ferrite in austenitic stainless steel welds, so the requisite resolution is possible. Again, the question about post-inspection removal of particulate must be addressed. At least wet methods will not have to worry about pyrophoricity of the particles. Eddy currents, capacitive field, and electrical resistivity techniques could be useful in selected materials, but would require development for smaller sizes. Similar to ultrasonic techniques, the limitation would be on the sensor sizes, though capacitive fringe sensors have been developed which are capable of detecting quite small features, such as machining burrs. Whether or not any of the above techniques is suitable for feedback control, or at least real-time inspection depends upon how rapidly the return signal can be evaluated. If a travel speed of 1mm/s combined with a laser power which produces a weld width of ~50 µm, the distance traveled in
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50 ms is approximately equal to the weld width. This would allow ample time for real time control of the process, if a suitable sensor can be found. On the other hand, if the travel speed increased by two orders of magnitude, the required control loop time would decrease inversely by the same magnitude, and the controller would have to be quite fast to react. X-ray techniques requiring multi-axis irradiation and computational processing (i.e. computed tomography) are clearly not suitable, even if it were possible to integrate an X-ray source with a welding station. Neither are techniques such as dye penetrant or magnetic particle techniques, as it is hard to imagine how they could be integrated into a welding process controller. Conceptually techniques that employ pitch-catch or pulse-echo processes, such as ultrasonic, eddy current, capacitive field or resistivity methods could be employed. However, for applications where the laser beam is scanned rather than moving the part, either a compatible sensor scanning technique (e.g. a phased array sensor) would be needed or else a supplemental sensor motion system would have to be added. Needless to say, moving a sensor would be very much slower than the motion of a scanned laser beam, slowing throughput dramatically. The final approach to weld evaluation again takes a page from the microelectronics industry, and that is simply to test the assembled part and compare it with expected behavior (the “known good die” approach). With highly-automated testers and relatively few parameters to check, this becomes feasible.
14.7
Applications of laser microwelding
Laser microwelding already plays an important role in joining microparts for the electronic, electrical, automotive and medical industries. Industrial applications of laser microwelding started as an alternative to precision resistance spot welding. Pulsed Nd:YAG lasers in particular became accepted by industry due to their non-contact processing, high precision, minimal heat affected zone (HAZ) and thermal distortion, and for imposing less restriction on the workpieces’ shapes.
14.7.1 Industrial base of microwelding lasers Figure 14.33 shows the trend of Nd:YAG lasers for microwelding sold in Japan, which is one of the most important markets for microwelding.72 Lamppumped Nd:YAG lasers are mainly used because of their high peak power. Lamp-pumped pulsed Nd:YAG lasers with average power less than 100 W are most widely used for microwelding, and approximately 900 units were installed in 2006. Nd:YAG lasers with higher average power, 100~600 W, are also increasingly finding applications where higher throughput is desired. More than 400 units were installed in 2006. The main applications of
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1200 PW < 100W PW 100–600W CW < 1000W
Number of units
1000
800
600
400
200
0
2003
2004
2005
2006
2007 (expected)
Year
14.33 Number of Nd:YAG microwelding lasers sold in Japanese market vs calendar year (PW: pulsed, CW: continuous power waveform).
microwelding by Nd:YAG lasers include spot welding of a variety of miniature motors, electronic parts, batteries and optical parts.
14.7.2 Examples of laser microwelding applications Lithium ion batteries are widely used for hand-held phones and laptop computers, and recent trends toward greater capabilities along with miniaturization require welding of thinner and lighter structures with high mechanical strength at high throughput. Figure 14.34 shows an example of a lithium ion battery with a 0.2 mm thick aluminum alloy case welded to the 1 mm thick end cap. Welding is performed at average power of 300 W (1.5J/ pulse, 200 Hz) at 30 mm/s through a stepped index fiber with a diameter of 0.4 mm. Small size motors are widely used for a variety of electronic devices including CD players, DVD/HDD recorders, hard disk drives for personal computers, hand-held phones, printers and so on. Pulsed Nd:YAG laser welding has played an important role in precision assembly of miniature motors with high levels of productivity. Figure 14.35 shows an application of the pulsed Nd:YAG laser to assembling vibration motor parts. Pressed metal sheets are stacked and welded by a pulsed Nd:YAG laser beam of average power 200 W, which is split into three beams, producing three seam welds simultaneously. Several assemblies are processed each second. Since the memory density of computer memory hard disks is doubling
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14.34 Welding of lithium ion battery by pulsed Nd:YAG laser (a) top view, (b) butt joint of aluminum alloy of 0.2 mm thick case to 1 mm thick cap. Welding was performed at 35 mm/s using SI fiber with 0.4 mm in diameter (Courtesy of Miyachi Technos). Welding
Welding
14.35 Laser microwelding of small motor (left) side view, (right) top view.
every year, so must the precision requirements being imposed during their manufacture. In response to this trend, laser microspot welding has been used for high-throughput assembly of hard disk suspensions. A welding system containing a single-mode pulsed Nd:YAG laser, with a 200 µm diameter graded index (GI) optical fiber delivery, galvano-scanner and imaging lens was developed for spot-welding of these hard drive suspensions, which consist of thin stainless steel foils.73 Since the thin metal foil is easily deformed by weld-induced shrinkage, precise control of the thermal deformation is one of the most important tasks in laser welding of the hard disk suspension. Thermal distortion in spot welding is proportional to the molten volume as shown in
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Fig. 14.36(a). A special requirement to provide 10~20 µm spot welding thermal distortion indicated a melt diameter of 100–150 µm would be required. The above-mentioned system was used to provide 120 µm melt diameter spot welds. Figure 14.36(b) shows the suspension for a hard disk where two stainless steel foils of 20 µm and 40 µm thickness are lap-welded at a rate of 150 spots/s. The lateral focus point location of the galvano-scanner-steered laser beam can be maintained within 10 µm of the desired aim point in continuous 24 hour/day operation. Laser micro bending technology is also used for roll and pitch adjustment of magnetic head suspensions, and flattening or crowning of the air bearing surface of a magnetic head slider.74 Laser welding of optical fiber connectors is another major application of laser microwelding, producing a strong weld joint with minimal thermal distortion. Figure 14.37 shows a system for welding optical fiber connectors. Multiple location welding is performed simultaneously by split pulsed Nd:YAG laser beams to compensate for thermal deformation.75 In this system, no mechanical force is applied in the direction of the optical axis in order to allow free thermal expansion of the joined parts, dramatically reducing the thermal deformation to 1/5th of the previous conventional welding system. Thermal deformation was monitored during the welding development process by six-axis force and moment sensors with high time resolution. Although appreciable deformation was observed in the direction of the optical axis immediately after the laser irradiation, it decreased to negligible values within 30–40 ms. On the other hand, the bending moment increased with time, saturating after 100–120 ms. These measurements played an important role in understanding the mechanism of the thermal deformation thereby allowing development of a holding tool for the welding system, which minimized the thermal deformation while reducing processing time. The previous conventional Required melt diameter (µm) Required deformation (µm)
Deformation (µm)
40 30
20
10
0 0
50 100 150 200 Diameter of molten zone (µm) (a)
250 (b)
14.36 Laser microwelding of hard disk suspension (a) melt diameter vs thermal deformation, (b) appearance of laser welded suspension.
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Optical fiber Lens and isolator
Laser diode Alignment
Laser welding
Deformation
(a)
0.2 0 –0.2 –0.4 –0.6 –0.8 –1.0 –1.2 –1.4 7400
Mx Fx Fy
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Moment Mx My Mz (kgcm)
Deformation Fx Fy Fz (kgf)
Deformation at simultaneous welding of three spots (sampling period: 1 ms) 0.6 0.4
(b)
14.37 (a) Laser welding system for development of optical fiber connector welds with 6-axis force and moment sensors and results showing the change in (b) bending force and moment vs time.
system prevented thermal deformation in the z-direction, and the resulting excessive thermal distortion had to be adjusted empirically, afterwards. Conduction mode laser welding has found an interesting application for kitchen equipment. Figure 14.38 shows the cross-section of stainless steel sheets of thickness 0.9 mm and 0.8 mm welded by means of a diode laser of 2.5 kW output power. The conduction welding process results in smooth seams, and has led to the very first industrial application of high power diode lasers for kitchen sink manufacture.76 This application used to be implemented
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14.38 Cosmetic weld of stainless steel sheets (upper: 0.9 mm, lower: 0.8 mm in thickness) with high power laser diode.
by conventional gas-tungsten arc (GTA) welding, which required tedious post processing grinding or repair work, while the diode laser welding provides an optically perfect “cosmetic” weld seam with minimal post-process polishing. This led to an overall cost benefit, even though the investment for a 2.5 kW diode laser was much higher than for a GTA welding machine. Transmission laser welding of polymers was first described in 1985 for welding automotive components.77 A contemporary application of this process being applied to a part in mass production is for a remote entry keyless module for a Mercedes-Benz automobile78 as shown in Fig. 14.39. In this application, carbon black pigment was used as the additive colorant to enhance absorption for transmission laser welding. The black housing is laser welded in an overlap joint configuration. The requirement for two different colored materials at the joint is a limitation for the process in applications where appearance is important. Dyes mixed into the bulk of the plastic have been used that appear black to the eye but transmit infrared, which allows allblack components to be welded. In this production application, laser welding exhibits superior weld quality and reproducibility compared to conventional plastic welding techniques.
14.8
Novel laser microwelding technologies
As described in the previous section, laser microwelding has been finding a variety of industrial applications in different fields. In addition, some innovative laser microwelding-based joining technologies have been recently developed,
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14.39 Laser welding of plastic for remote entry keyless module for Mercedes-Benz automobile.
and are introduced in the following sections. From both technological and fundamental view points, they are expected to play an important role in the near future.
14.8.1 Laser droplet welding Laser droplet welding (LDW) has been developed to overcome existing laser welding drawbacks concerning gap bridging, joining of highly reflective materials, thin metallizations and heat sensitive materials. In LDW, a liquid metal droplet is created at the end of a wire by pulsed laser irradiation, and subsequently placed on the parts to be joined to achieve a weld. Figure 14.40(a) shows schematically an LDW setup, which consists of a pulsed Nd:YAG laser with triple optical beam splitting, a wire feed system, a target positioning system, shielding gas supply and a mechanical positioning system for the laser optics.79,80 The LDW process is split into five phases; droplet creation, droplet detachment, droplet flight, droplet landing and droplet solidification. Figure 14.40(b), shows the optimized pulse profile of the laser beam for 0.6 mm Niwire. The droplet is created at laser power P1 with duration of t1 to melt the wire. Then the laser power is decreased to P2 for time t2, during which the wire is moved to position the liquid droplet below the focus of the laser. Then the droplet is detached from the wire by a short and comparatively high peak P3. During phase t2, the wire is moved to avoid having the pulse peak P3 hit directly on the molten metal and produce a splash. The flight time is short enough for the droplet to retain its high temperature. The positional accuracy of the droplets is important for joining small
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t3
2500 2000 1500
Wire feed continues
p1
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p3
500 Pulse profile in principle
Target material
p2
0 Target position
0
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10 Pulse time (ms)
15
14.40 Schematic drawing of laser droplet welding set-up (b) schematic pulse profile for 0.6 mm Ni-wire.
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parts, and a sequence of deposited droplets made at 3 Hz is shown in Fig. 14.41. One can see the positional variation is +/–0.5 mm, which is independent of the droplet transit distance of between 1.3–4 mm for bead (droplet) on plate welds. The contact angle of ~90° on landing is reduced as the droplet wets the base material. During the final weld phase, solidification (and joining) between the droplet material and the work pieces occurs at a high cooling rate, resulting in a very narrow HAZ. Figure 14.42 shows an example of joining two stainless steel sheets of 200 µm thickness and 200 µm gap with nickel droplets. Joining was carried out with two droplets with very high positioning accuracy along the center of the gap (because of self alignment via the combined effects of capillarity
14.41 Sequence of droplets deposited by laser droplet welding.
Drops
14.42 Two 200 µm stainless steel sheets joined with Ni droplets via LDW.
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and laminar flow of the shield gas passing through the gap to precisely colocate the droplets). The plates are seen to have a narrow heat affected zone, but at the end of the plates the material has melted, effecting the joint. A local mixing of the materials occurs, which is seen in the right-hand image. Experience has shown that with accurate control of droplet energy, it is possible to join delicate components such as foils to thin metal coatings on a silicon substrate as shown in Fig. 14.43(a).81 In addition, LDW is very well suited to joining parts of different geometries, dimensions and materials, as shown in Fig. 14.43(b). Due to the droplet volume, LDW is also suitable for bridging gaps. LDW technology also appears promising for wire-to-foil connections.79 Successful LDW fusion welding of Ti to stainless steel has been reported, which is not possible by conventional welding techniques due to the large difference in thermal expansion coefficient α (αTi = 8.9 × 10–6/K, αSUS = 16 × 10–6/K) combined with the formation of brittle intermetallic compounds on solidification.82 The only successful welding techniques for this combination are solid state, e.g.: diffusion bonding, which requires careful surface preparation, high temperatures and pressures, and appreciable time, and friction welding, which is restricted as to size and geometry.
14.8.2 Laser spike welding As has been noted, dealing with gaps in microwelds is problematic for conventional laser welding procedures. Recently, a technique was developed which takes advantage of recoil pressure-driven material flow to bridge gaps in lap geometry spot welds.42 The procedure involves melting a pool at a laser power that gives conduction mode in the upper layer, then, when it is sufficiently large, and completely penetrated, increasing the laser power to a
200 µm
500 µm (a)
(b)
14.43 (a) Cross-section of silver droplet joint of silver foil onto a silver-coated silicon surface (b) cross-section of a droplet joint of parts of different materials and dimensions.
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level and for a time that generates sufficient recoil pressure to propel the diaphragm-like liquid pool into contact with the lower layer. Contact with the lower layer must be maintained for long enough to cause adherence via either superficial surface melting, or if the lower surface is clean enough, by a braze-like adhesion. Difficulties occur if the lower layer is too conductive, like Cu. Nevertheless, as shown in Fig. 14.44, where the technique was applied to stainless steel material 250 µm thick, gaps of up to 100% of the top layer thickness can be successfully bridged. If the cosmetic appearance of the top surface of the “spike” weld is problematic, a second pass can be applied to smooth it.
14.8.3 SHADOW welding SHADOW (Stepless High-Speed Accurate and Discrete One-pulse Welding)83 is a technique which adapts macro laser welders to the microworld. In effect, by taking a single pulse from a pulsed Nd:YAG laser, combining it with the rapid motion of a small workpiece relative to the beam (either by moving the part, or by moving the beam with a scanner head), one transforms a macro laser spot welder into a micro laser seam welder. Since electronically controlled
(a)
(b)
14.44 Laser spike welds in stainless steel sheets of 250 µm thickness X8CrNi1812 Stainless steel showing effect of gap (a) top view, and (b) cross-sections of welds with increasing gap from left to right.
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flashlamp or diode pumped lasers typically have maximum pulse lengths of ~20 ms, the size range of SHADOW welds are constrained by travel speed/ laser power/molten zone size requirements. As seen in Section 14.5.3, the power requirements for a microwelding laser system are modest, so the only concern is if the system to be transformed can be controlled accurately enough in its low power range. Further, since high travel speed is needed, the geometry of the weld generally must be fairly simple, such as a straight line, a circle, or at most a polygon with rounded corners if the part is moved, though using high response rate mirrors in a galvano-controlled scanner to move the beam instead greatly mitigates this requirement. Initial application of this technique was in the mechanical watch industry, as shown in Fig. 14.45, where gear-to-shaft welds of extreme regularity and excellent cosmetics were demonstrated.84 The first cosmetic advantage comes from not having a rippled weld bead caused by repeated overlapping spot welds, as is typical for most “micro” welds made using pulsed Nd:YAG welders. As has been known for many years, the total energy used for overlapping spot welds is many times that needed for a CW weld of the same size, simply because the same volume of molten metal is melted and resolidified multiple times as the seam is built up from individual spots. Since the metallic vapor soot deposited and spatter likelihood is proportional to the amount of energy put into the weld, reducing the overall heat input via the SHADOW technique also reduces these contributions to cosmetic unsightliness. A natural consequence of the rapid travel speed is that the melting efficiency is also quite high, as little time for thermal diffusion is allowed. This makes the process useful for high conductivity materials. Further, coupling of the beam to high conductivity material can be enhanced by modifying the initial shape of the pulse with a so-called leading edge spike. After the initial coupling occurs, the laser will generally maintain good coupling to the part. This is not always assured with an overlapping
(b)
(a)
14.45 Conventional pulsed laser seam weld (left) vs SHADOW weld (right) of watch gear to shaft.
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pulse seam, hence the reproducibility is enhanced as well. As a final benefit, since the geometry of the fusion zone is quite symmetrical, radial displacement of the shaft from the central location, and tilting of the shaft relative to the plane of the gear is also minimized. Some more recent applications of the technique include welding Cu to carbon die in a leaf spring and Cu-Be ball bearing assemblies (see Fig. 14.46).85
14.8.4 Microwelding of thin foil by single-mode fiber laser The single-mode fiber laser (SMFL) with its near-diffraction-limited beam can provide a very small focused spot diameter, essentially limited only by lens aberrations (which is not a serious problem with modern optics). It thereby provides extremely small keyhole diameters, enabling welding with very high spatial resolution. In addition, for both flow stability and surface energetic reasons, the keyhole has been found to become more stable as its size decreases, enabling welding with extremely high speed in very thin metal foil. Keyhole welding at high welding speeds normally results in a humped bead because the molten metal flow becomes unstable as described on pages 374–375. It was noted that welding with a small beam diameter is calculated to decrease the flow velocity at the keyhole perimeter, which should suppress humping. Figure 14.47 shows the cross-section of a weld bead using a singlemode fiber laser made at a welding speed of 1.5 m/s, which is the highest welding speed ever attained without humping by a single translating heat source,22,86 conclusively verifying the calculations. It was noted on pages 375–378 that lap welding of very thin metal foil to bulk material is difficult, since welding results in non-bridged beads without special tooling. The relationship between the thickness of the foil and the diameter of the keyhole (or beam diameter) h < 2r has been shown to be a critical parameter because of its link to dS/dr via equation 14.15. Since the SMFL has an extremely small value for 2r, it follows that it will be useful for extremely small h materials as well. Figure 14.48 shows an example of lap welding of thin foils of 20 µm and 30 µm by using an SMFL with a focus diameter of 10 µm. Sound lap welding was obtained using a simple fixuring tool, since dS/dr > 0 is satisfied.22,86 Lap joints between foils of 10 µm and 30 µm thickness resulted in a non-bridged bead when an appreciable gap existed between foils, since no common keyhole could be produced. It is seen, however, that sound welding was obtained at zero gap (again without precise fixturing and tooling), since the keyhole was sustained stably enough that that a common keyhole could be produced between the two foils.
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100 µm
Best welding results are obtained with a feed rate of v = 650 mm/s, a pulse length of τH = 14.5 ms and a pulse energy of Ep = 28 J. Pulse forming is necessary to avoid cavities. (a)
Ball bearing welded with SHADOW® Inner cage d = 0.6 mm Outer cage d = 3.0 mm The semi-circular shape indicates heat conduction welding. The gap between the inner and the outer cage reaches into the weld. The SHADOW® technique offers some advantages: The strength of the SHADOW® welded ball bearings is Fs = 90 N while for the flanged ball bearings stand FF = 75 N. (b)
14.46 SHADOW welds of (a) Cu leaf spring to carbon die, (b) Cu-Be ball bearing journal (cage) to shaft.
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Left leaf spring welded to a carbon die Right: cross-section
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14.47 Bead cross-section of stainless steel foil made by single-mode fiber laser (thickness: 40 µm, welding velocity: 1.5 m/s).
(a)
(b)
(c)
14.48 Lap welding of thin stainless steel foils by single-mode fiber laser with focus spot size of 10 µm (a) welding of 20 µm/30 µm in thickness (25 W, 1 m/s), (b) welding 10 µm/30 µm in thickness with clearance, (c) welding of 10 µm/30 µm thickness with contact.
14.8.5 Ultrashort pulse fusion welding of glass Glass has many potential optics, MEMS, electronics, and biomedical applications due to its excellent optical, mechanical, electrical and chemical properties. However, fusion welding of glass is very difficult because of its brittleness, requiring high heat input, low thermal gradient glass-blowing techniques. Only glass with a very small thermal expansion coefficient like fused silica can be laser-welded,87 and thus most glass joining has depended on adhesive bonding with relatively poor mechanical, thermal and chemical durability.
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Recently, a novel technique using ultrashort laser pulses has been developed which lap welds glass plates at their interface. This technique utilizes a laser wavelength normally transparent to the glass material given by λ > hc/Eg (where Eg = band gap energy, h = Planck’s constant, c = speed of light); the difference from previous transparent material welding techniques is that the laser energy is absorbed by a nonlinear process. No addition of colorant or absorbing material as is done for transmission polymer welding is needed.27,88 Table 14.5 shows glass welding reported so far using ultrashort pulse lasers, whose pulse durations ranged between 85fs and 10ps. The welding speed has increased significantly by using lasers with higher repetition rates and higher average power, compared with welding speeds limited to a few µm/s in initial work which used a Ti-sapphire laser with a pulse repetition rate of only 1 kHz. Figure 14.49(a) shows a schematic illustration of the lap welding of two glass plates. When the ultrashort pulse laser is focused at the interface of two glass plates, laser energy absorbed by a nonlinear process is transferred to the atoms of the glass material after the laser pulse, and then the surrounding area is heated by thermal conduction. It should be noted that only the interface of the two plates is heated and melted by the nonlinear process since the laser wavelength is transparent to the glass material except very near sharp focus. Figure 14.49(b) shows the cross-section of the weld bead obtained by moving the focused laser beam spot in a direction perpendicular to the laser beam axis along the interface. Figure 14.50(a) shows the effect of welding velocity on the dimensions of the melted zone by a laser of pulse duration 10 ps pulsed at 500 kHz. A melt width from 20 µm to 100 µm is obtained at travel velocities between 1 mm/ s and 100 mm/s at an average laser power of 700 mW. Solid lines calculated by a thermal conduction model agree well with the measured values. Figure 14.50(b) shows joining efficiency (defined by S/W where S is joint area created per unit time and W laser power). It is noted that joining efficiency exceeds the highest value obtained so far in metal welding, and is higher than the joining efficiency obtained with 400 fs laser pulses, also plotted. Although lasers with ultrashort pulse durations from fs to ps are available, it should be noted that ps lasers are more practical for fusion welding of glass, because they provide high nonlinear absorptivity28 and a simpler, more reliable system in comparison with fs lasers. Figure 14.51 shows SEM fractographs indicating the appearance of the peeled-apart weld joint. The separated sample exhibits brittle fracture along an irregular surface away from the original fusion zone interface, indicating that the two plates were strongly joined together. A mechanical strength higher than 50 MPa has been obtained at optimized conditions.94
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Table 14.5 Laser microwelding of glass Max average power
Welding velocity
Pulse repetition rate
Pulse duration
Material*
Year
Ref
800 nm 800 nm 1558 nm 1045 nm 1064 nm 1045 nm 1064 nm
0.001W 0.012W 0.4W 0.55W 0.5W 0.7W 6W
0.005 mm/s 0.1–1 mm/s 0.02–0.2 mm/s 0.2–125 mm/s 0.5–200 mm/s 1–100 mm/s 5–500 mm/s
1 kHz 1 kHz 500 kHz 1 MHz 0.1 and 1 MHz 100 ~ 640 kHz 0.1 ~ 1 MHz
85 fs 85 fs 950 fs 360 fs 325 fs, 400 fs 10, 15 ps 10 ps
F B-F A A-silicon B B B FB
2005 2006 2006 2006 2007 2007 2007
88 89 90 91 92 28 94
*: F = fused silica, B = borosilicate glass, A = non-alkali aluminosilicate glass
Laser microwelding
Wavelength
407
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Focusing lens
Welding n directio
Glass A
Glass A Glass B 25 µm Weld seam Glass B (a)
(b)
14.49 (a) Schematic diagram of seam welding of two glass plates. The interface of the glass plates is selectively melted, (b) crosssection of lap weld.
14.9
Future trends
With the advent of readily-achievable near-diffraction-limited laser spot sizes, the future prospects of laser microwelding, at least down to sizes of a few microns extent, will be limited by details of fixturing, handling, quality control, and the imagination of designers and technologists tasked with making their ideas work. As is always the case, and never more so than in the microworld, the devil is in the details! Whether or not lasers will play a role in the nanoworld is another question. Again, we are confronted with a semantic problem: what constitutes a nanoweld? Going by analogy with the definition of a microweld, something on the order of 100 nm will be implied. UV lasers are presently available down to <200 nm wavelengths, and could produce spot sizes that might be considered on the borderline of nano. Further, methods of producing subwavelength of light spots have been reported which take into account nearfield enhancement of intensity.95,96 To the authors’ knowledge, however, these methods require close contact with the surface being irradiated. Finally, the beam penetration of laser light into a material starts to become comparable with these distances, turning it into a volume source, rather than a surface source. Nevertheless, the need to communicate with nanostructures from the macroworld suggests that there will be a need to join at such size scales. Whether high energy sources like lasers will be needed, or whether surface energy considerations will allow touch/stick methods to suffice is for the future to disclose.
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α = 0.003 cm2/s
Melt width Ht (µm)
120 0.004 cm2/s
100
y
Vt
0.005 cm2/s
80
Ht
60 40 20 0
1
10 Traveling velocity v (mm/s) (a)
Joining efficiency R (mm2/J)
7
100
10ps 500 kHz–1.4 µJ
6 5 4 3
Metal welding with single-mode fiber laser
2
406 fs 1MHzHz–0.5 µJ
1 0 1
10 Traveling velocity v (mm/s) (b)
100
14.50 Effect of welding speed on (a) melt width and (b) joining efficiency at average power of 0.7 W (500 kHz and 1.4 mJ). Solid lines are calculated by thermal conduction model.
14.51 SEM fractograph of peeled-off weld joint.
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14.10 Acknowledgements Sandia is a multiprogram laboratory operated by Sandia Corporation, a Lockheed Martin Company, for the United States Department of Energy’s National Nuclear Security Administration under contract DE-AC0494AL85000. We would also like to thank Mr. D. O. MacCallum and Mr. G. A. Pressly, both of Sandia National Laboratories, for reviewing the text.
14.11 Equation annex Equation 1:97 Instantaneous stationary point source at origin, with total quantity of heat introduced Qρc: T=
2 Q exp – r 3/2 4αt 8( πα t )
where r2 = x2 + y2 + z2 or the continuous point source q per unit time: q r erfc 4πα r 4αt Equation 2:98 Instantaneous line source, with total quantity of heat liberated Q, parallel to z axis, passing through x′, y′: T=
T=
2 2 Q exp – {x – x ′} + {y – y ′} 4πα t 4 αt
or the continuous line source q per unit length per unit time:
T=
q 4 πα
∫
∞
r 2 /4 αt
e – u du u
99
Equation 3: a continuous point source q moving at velocity v in the x direction
T=
q – v( r – x ) exp 2 πkr 2α
r as defined in Equation 1 (except take square root) Equation 4:100 moving round Gaussian distributed heat source on a surface z = 0:
T=
∫
t
0
q t ′′ –1/2 ⋅ πρc(4 π a )1/2 2 at ′′ + σ 2
2 w 2 + y 2 + 2 wvt ′′ + v 2 t ′′ 2 × exp – z dt ′′ 2 4 at ′′ 4 at ′′ + 2 σ
where t″ = t – t′, x – vt′ = w + vt″ and w = x – vt
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Equation 5:101 moving rectangular shape (2b × 2l) top-hat distributed heat source on surface z = 0, moving at velocity U in the x-direction:
T=
Qα 4 kU 2 π
∫
∞
0
2 exp – Z ⋅ erf Y + L – erf Y – L 2u 2 2 u u
X+B+u × erf – erf X – B + u du 2u 2u u where: X = Ux/2α, Y = Uy/2α, Z = Uz/2α, L = Ul/2α, and B = Ub/2α. Equation 6:102 a continuous moving line source of strength q′ per unit length per unit time:
T=
( )
q′ vr exp vx K 0 2 πk 2α 2α
where K0 is the modified Bessel function of the second kind of zero order. Equation 7:22 a moving uniformly distributed cylindrical heat source of power density w: R
∫ ∫
T ( x, y) = w 2 πKh
–R
–
R2 – x ′ 2 R2 – x ′ 2
v( x – x ′ ) 2 vr dx ′dy ′ × exp – K0 α 2 2α Equation 8:22 a moving uniformly distributed square columnar heat source of power density W gives an average temperature within the heat source:
T=
∫
W 32 cρa 4 h
∞
0
2 a + vt vt 4 α t erf – 2erf 4α t 4 t α
2 a – vt 2a + erf erf dt 4αt 4α t where erf is the error function and a is the half-width of the heat source. Equation 9:103 A double-ellipsoidally distributed volume heat source (different ellipsoidal shapes for leading and trailing portions of the heat source) traveling in the x direction at velocity v:
T=
3 3Q 2 ρcπ 3/2
∫
τ
0
dt ′ (12 α ( t – t ′ ) + a h2 )(12 α ( t – t ′ ) + bh2 )
A′ + × 2 (12 α ( t – t ′ ) + c hf )
B′ (12 α ( t – t ′ ) + c bf2 )
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where
3( x – vt ′ ) 2 3y2 – A ′ = r f exp – 2 12 α( t – t ′ ) + a h2 12 α ( t – t ′ ) + c hf
–
3z 2 2 12 α ( t – t ′ ) + bh
and 3( x – vt ′ ) 2 3y2 – B ′ = rb exp – 2 12 α ( t – t ′ ) + c hb 12 α( t – t ′ ) + a h2
–
3z 2 12 α ( t – t ′ ) + bh2
where ah, bh, chf, chb are the ellipsoidal minor and major axes in the y, z, x forward and x backward directions, respectively, and rb and rf are heat apportionment coefficients to the back and front of the heat source (and sum to 2), such that rf = 2chf/(chf + chb) and rb = 2chb/(chf + chb).
14.11 References 1. Einstein, A., “Quantentheorie der Strahlung” (On the quantum mechanics of radiation), Physikalische Zeitschrift, 18, 1917, pp. 121–128. 2. Maiman, T.H., “Stimulated Optical Radiation in Ruby,” Nature, 187, 1960, pp. 493–494. 3. Ready, J.F., Industrial Applications of Lasers, Academic Press, New York, 1978, pp. 374–375, 387, 394–395. 4. Keister, F.Z., Engquist, R.D., Holley, J.H., “Interconnection Techniques for Micro circuits,” IEEE Transactions on Component Parts, 11(1) 1964, pp. 33–41. 5. Fritzke, G.P., “Joining Electronic Materials,” Report No D 8-23.1 of the American Society for Metals, 1968, SLAC-PUB-468, August 1968. 6. Rischall, H., Shackleton, J.R., “Laser Welding for Microelectronic Interconnections,” IEEE Transactions on Component Parts, 11(3), 1964 pp. 145–151. 7. Rasera, R.L., Bernstein, J.B., “Laser Linking of Metal Interconnect: Linking Dynamics and Failure Analysis,” IEEE Transactions on Components, Packaging and Manufacturing Technology, Part A, 19, Dec. 1996, pp. 554–561. 8. Latham, W.P., Kar, A., Handbook of Laser Materials Processing, J.F. Ready, ed., Laser Institute of America, Magnolia Publishing, 2001, p. 105. 9. ISO 11146 Lasers and laser-related equipment – Test methods for laser beam parameters – Beam widths, divergence angle and beam propagation factor. ISO 11146:1999(E). 10. Hagan, D.J., Kik, P.G., Class notes OSE5312 – Fundamentals of Optical Science, CREOL, University of Central Florida, Fall 2006. 11. Milonni, P.W., Eberly, J.H., Lasers, John Wiley and Sons, New York, 1988, p. 15.
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12. Modest, M.F., Handbook of Laser Materials Processing, J.F. Ready, ed., Laser Institute of America, Magnolia Publishing, 2001, pp. 175–183. 13. Ready, J. F., ibid., p. 341 (Note: our figure showing absorptivity was obtained by inverting the original reflectivity plot). 14. Hummel, R.E., ibid., p. 167. 15. Ibid., p. 169. 16. Modest, M.F., ibid., p. 182. 17. Ready, J.F., ibid., p. 356. 18. Qiu, T.Q., Tien, C.L., “Heat transfer mechanisms during short-pulse laser heating of metals,” ASME J. Heat Transfer, 115, 1993, pp. 835–841. 19. DebRoy, T., David, S.A., “Physical processes in fusion welding,” Rev. of Mod. Phys., 67, 1995, pp. 85–116. 20. Miyamoto I, Maruo H, “Spatial and temporal characteristics of laser-induced plasma in CO2 laser welding”, 2nd International Congress on Laser Advanced Materials Processing Conference (LAMP’92), Nagaoka, Japan Laser Processing Society, 1992, pp. 311–316. 21. Bohren, C.F., Huffman, D., Absorption and Scattering of Light by Small Particles, John Wiley & Sons, New York, 1983. 22. Miyamoto, I., Park, S-J., Ooie, T., “High-speed microwelding by single-mode fiber laser,” Proceedings of the (Conference on) Laser Assisted Net Shape Engineering 4, 2004, pp. 55–66. 23. Arata, Y., Miyamoto, I., “Laser welding” Technocrat, 11 (5), 1978, pp. 33–42. 24. Miyamoto, I., Maruo, H., Arata, Y., “Mechanism of bead-transition in laser welding”, Int Conference Welding Research in the 1980s, Osaka, The High Temperature Society, 1980, pp. 103–108. 25. Pittaway, L.G., “The temperature distributions in thin foil and semi-infinite targets bombarded by an electron beam”, Brit. J. Appl. Phys., 15, 1964, pp. 967–982. 26. Jones. I. A., Hilton, P. A., Sallabanti, R., “Use of infrared dyes for transmission laser welding of plastics”, 18th International Conference on Applications of Lasers and Electro-optics (ICALEO’99), San Diego, Laser Institute of America, 1999, pp. B71–B79. 27. Stuart, B.C., Feit, M.D., Herman, S., Rubenchik, A.M., Shore, B.W., Perry, M.D., “Nanosecond-to-femtosecond laser-induced breakdown in dielectrics”, Phys. Rev. B53, 1996, pp. 1749–1761. 28. Miyamoto, I., Horn, A., Gottmann, J. “Local melting of glass material and its application to direct fusion welding by ps-laser pulses”, JLMN-J. Laser Micro/ Nano and Microengineering, 2007, 2, 02, [06–072]. 29. Benter, C., Petring, D., Poprawe, R., “Investigation of the transition from heat conduction to deep penetration welding with high power diode lasers”, 3rd International WLT – Conference on Lasers in Manufacturing, Munich, WLT, 2005, pp. 67–72. 30. Miyamoto, I., Park, Seo-Jeong, Kosumi, T., Ooie, T., “Ultrafine-keyhole welding process using single-mode fiber laser”, 22nd International Conference on Applications of Lasers and Electro-optics (ICALEO’03), San Diego, Laser Institute of America, 2003, #1010. 31. Deininger, C., Mueller-Borhanian, J., Dausinger, F., Huegel, H. “Development of multi-detector systems for the process monitoring of laser beam welding capable for industrial use”, Proc. LANE 2004, Erlangen, CIRP, 2004, pp. 107–117. 32. Seibold, G., Dausinger, F., Huegel, H., “Absorptivity of Nd:YAG laser radiation on
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Microjoining and nanojoining ion and steel depending on temperature and surface condition”, Proc. ICALEO2000, Laser Institute of America, Dearborn, 2000, pp. E125–E132. Kroos, J. “Dynamic behavior of the keyhole in laser welding”, J. Appl. Phys., 26, 1993, pp. 481–486. Miyamoto, I., “Fundamentals of laser welding”, Proc. 26th Japan Laser Society Symposium, Japan Laser Processing Society, Osaka, 1991, pp. 1–17 (in Japanese). Miyamoto, I., Uchida, T., Maruo, H., Arai, T., “Mechanism of soot deposition in laser welding”, 13rd International Conference Applications of Lasers and Electrooptics (CALEO’94), Orlando, Laser Institute of America, 1994, pp. 293–302. Kim, K.R., Farson, D.F., “CO2 laser-plume interaction in materials processing,” Journal of Applied Physics, 89, 2001, pp. 681–688. Swift-Hook, D.T., Gick, A.E.F., “Penetration welding with lasers,” 52, Welding Journal, 1973, pp. 492s–499s. Andrews, J.G., Atthey, D.R., “Hydrodynamic limit to penetration of a material by a high-power beam”, J. Phys. D: Appl. Phys., 9, 1976, pp. 2181–2194. Mazumder, J., Steen, W.M., “Heat transfer model for CW laser material processing”, J. Appl Phys., 51, 1980, pp. 941–947. Zacharia, T., David, S.A., Vitek, J.M., Debroy, T., “Heat transfer during Nd:YAG pulsed laser welding and its effect on solidification structure of austenitic stainless steel”, Metall. Trans. A, 20A, 1989, pp. 957–967. Ki, H., Mohanty, P.S., Mazumder, J., “Modeling of laser keyhole welding: Part I – Mathematical modeling, numerical methodology, role of recoil pressure, multiple reflections, and free surface evolution”, Metall. & Mater. Trans A, 33A, 2002, pp. 1817–30; Part II – Simulation of keyhole evolution, velocity, temperature profile, and experimental verification”, ibid, pp. 1831–1842. Dijken, D.K., Hoving, W., De Hossen, J.T.M., Laser penetration spike welding: A microlaser welding technique enabling novel product designs and constructions, Journal of Laser Applications, 15, 2003, pp. 11–18. Vanderlee, K.A., Porter, R.B., Kulterman, R.W., Chalco, P.A., “Lasersonic bonding of TAB components to epoxy-glass circuit boards”, IEEE Trans. on Comp. Pack, & Man. Tech. C, 19, 1996, pp 277–282. Nassar, R., Dai, W., Modelling of Microfabrication Systems, Springer-Verlag, Berlin, 2003, pp. 77–121. Fuerschbach, P.W., Eisler, R.G., SOAR shareware available under license from Sandia National Labs. O’Brien, R.L. (ed.), Welding Handbook, 8th Edition, Volume 2 Welding Processes, American Welding Society, Miami, FL, 1991, p. 727. DuPont, J.N., Marder, A.R., “Thermal efficiency of arc welding processes”, Welding Journal, 74, 1995, pp. 406s–416s. Fuerschbach, P.W., “Measurement and prediction of energy transfer efficiency in laser beam welding”, Welding Journal, 75, 1996, pp. 24s–34s. Elmer, J.W., The Influence of Cooling Rate on the Microstructure of Stainless Steel Alloys, Sc.D. thesis, M.I.T., Sept. 1988, p. 93. Katayama, S., Matsunawa, A., “Solidification Microstructure of Laser Welded Stainless Steels”, Proc ICALEO, 1984, p. 60. Elmer, J.W., The Influence of Cooling Rate on the Microstructure of Stainless Steel Alloys, Sc.D. thesis, M.I.T., Sept. 1988, p. 214. ibid., pp. 149–172.
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53. Lippold, J.C., Kotecki, D.J., Welding Metallurgy and Weldability of Stainless Steels, Wiley-Interscience, Hoboken, New Jersey, 2005, pp. 182–186. 54. Knorovsky, G.A., MacCallum, D.O., “The effect of laser welding on LIGA materials”, Proceedings of ICALEO2001, Laser Institute of America, 2001, M802. 55. Katayama, S., Matsunawa, A., “Formation mechanism and prevention of defects in laser welding of aluminum alloy”, 6th International Conference on Welding and Melting by Electron and Laser Beams (CISFFEL 6), Toulon, Institut de Soudure, 1998, pp. 215–222. 56. Katayama, S., Kohsaka, S., Mizutani, M., Nishizawa, K., Matsunawa, A., “Pulse shape optimization for defect prevention in pulsed laser welding of stainless steels”, 12th International Conference Applications of Lasers and Electro-optics (ICALEO’93), Jacksonville, Laser Institute of America, 2001, pp. 487–497. 57. Herziger, G., “Laser materials processing in Europe”, 2nd International Congress on Laser Advanced Materials Processing (LAMP’92), Nagaoka, Japan Laser Processing Society, 1992, pp. 547–556. 58. Gorman, E.F., “New developments and applications in manual plasma arc welding”, Welding Journal, 48, 1969, pp. 547–556. 59. Lingenfelter, A.C., “Laser welding thin cross sections”, 1st International Congress on Laser Advanced Materials Processing Conference (LAMP’87), Osaka, Japan Laser Processing Society, 1987, pp. 211–216. 60. Kawahito, Y., Katayama, S., “In-process monitoring and adaptive control for stable production of sound welds in laser microspot lap welding of aluminum alloy”, J. Laser Applications, 17, 2005, pp. 30–37. 61. Miyamoto, I., Maruo, H., Arata, Y., “The role of assist gas in CO2 laser welding”, 3rd International Conference Applications of Lasers and Electro-optics (ICALEO’84), Boston, Laser Institute of America, 2001, pp. 68–75. 62. Petereit, J., Abels, P., Kaierle, S., Kratzsch, C., Kreutz, E.W., “Failure recognition and online process control in laser beam welding”, 24th International Conference Applications of Lasers and Electro-optics (ICALEO’05), Miami, Laser Institute of America, 2005, pp. 101–105. 63. Mikame, K., “Applications of laser material processing in Toyota Motor Corporation”, 2nd International Congress on Laser Advanced Materials Processing Conference (LAMP’92), Nagaoka, Japan Laser Processing Society, 1992, pp. 947–952. 64. Mori, K., Miyamoto, I. “In-process monitoring in laser welding by analyzing ripple of plasma emission”, J. Laser Applications, 9, 1997, pp. 155–159. 65. Kawahito, Y., Okada, T., “Intelligent laser process control in the micro spot welding for copper”, Proc. ICALEO2001, Laser Institute of America, 2001, #M803. 66. Kawahito, Y., Funami, K., Okada, T., Katayama, S., “In-process monitoring and adaptive control in laser micro-spot welding of copper”, The Review of Laser Engineering, 31, 2003, pp. 231–235 (in Japanese). 67. Weiss, P., “Pictures only a computer could love”, Science News, March 29, 2003, 163, (13), p. 200. 68. MIL-STD-202, MIL-STD-883, MIL-STD-1576. 69. Harmon, G.G., Wire Bonding in Microelectronics, 2nd edn, McGraw-Hill, New York, 1997, pp. 67–113. 70. Kercel, S.W., Kisner, R.A., Klein, M.B., Bacher, G.D., Pouet, B., “In-process detection of weld defects using laser-based ultrasound”, Proc. SPIE, 3852, 1999, pp. 81–92. 71. Harsanyi, G., Semmens, J.E., Martell, S.R., Percsi, L,. Toth, E., “Analyzing thick
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80.
81.
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83. 84.
85.
86.
87.
88.
Microjoining and nanojoining film multilayer defects using acoustic micro imaging”, Intl J of Microcircuits & Elec. Pack., 22, 1999, pp. 388–394. Simpo Corporation (2006), 2006/2007 Japan Laser World & Trend, Tokyo (in Japanese). Matsushita, N., Imakado, M., Tokura, Y., Furui, T., Ushimaru, A.,“Laser applications in magnetic disk production”, Welding Technologies, 52, 2004, pp. 67–71 (in Japanese). Matsushita, N., “Laser micro-bending for precise micro-fabricatioon of magnetic disk-drive components”, Int’l Symp 4th Laser Precision Microfabrication (LPM’03), Nagaoka, Japan Laser Processing Society, 2003, pp. 24–29. Matsushita, N., Yanagida, Y., Tsukahara, T., “Elaborate precision machining technologies for creating high added value at low cost”, Fijitsu Sci. Tech., 43, 1, 2007, pp. 67–75. Bachmann, F., “The impact of laser diodes to solid state lasers and materials processing applications”, Proceedings of the 61st Japan Laser Processing Society Meeting, Osaka, Japan Laser Processing Society, 2004, pp. 16–29. Toyota Jidosha. K.K., “Laser beam welding of plastic plates” Patent Application JP85213304, 26 Sep 1985. Brettschneider, C., “Im Vordergrund steht die Asthetik”, Laser-Praxis 2, 98, (10) 1998. Hoving, W., Jahrsdorfer, B. “Laser droplet weld – Ein innovatives Fuegeverharen”, Tagungsband Laser in der Elektronikproduktion & Feinwerktechnik – LEF2001, Meienbach, Bamberg, 2001, pp. 21–32. Jahrsdoerfer, B., Esser, G., Geiger, M., Govekar, E., “Laser droplet weld – an innovative joining technology opens new application possibilities”, SPIE Vol. 4977, Photon Processing in Microelectronics II, 2003, pp. 518–529. Govekar, E., Klemen, J., Kokalj, T., Schmidt, M., Kastens, M., “Progress in laser droplet formation and welding”, Laser in der Elektronikproduktion & Feinwektechnik, LEF 2007. Jahrsdörfer, B., Schmidt, B., Geiger, M. “Laser droplet welding and its potential for joining dissimilar materials”, Proceedings of LANE2004, 2004, pp. 651–659. Olowinsky, A.M., Klages, K., Gedicke, J., “SHADOW a new welding technique: basics and applications”, Proc. SPIE, 5662, 2004, pp. 291–299. Olowinsky, A., Kramer, T., Dumont, N., Hanebuth, H., “New applications of laser beam micro welding”, Proceedings ICALEO 2001, Laser Institute of America, 2001. Gillner, A., Olowinsky, A., Klages, K., Gedicke, J., Sari, F., “High precision and high speed laser microjoining for electronics and microsystems”, Proc. SPIE, 6161 2006, pp. 616102-1–616102-11. Miyamoto, I., Park, Seo-heong, Ooie, T., “Application of single-mode fiber-lasers to novel microwelding”, Proc. 4th Laser Precision Microwelding, Japan Laser Society, Nara, 2003, pp. 507–514. Arata, Y., Miyamoto, I., Takeuchi, S., “Dynamic behavior of laser welding and cutting”, Proc. Int. Conf. Electron and Ion Beam Science and Technology, 1976, pp. 111–128. Nolte, S., Will, M., Burghoff, J., Tuennermann, A. “Ultrafast laser processing: new options for three-dimensional photonic structures”, Modern Optics, 10, 2004, pp. 2533–2542.
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89. Tamaki, T., Watanabe, W., Nishii, J., Itoh, K. “Welding of transparent materials using femtosecond laser pulses”, Jpn. J. Appl. Phys. Part 2, 44, 2005, p. L587. 90. Watanabe, W., Onda, S., Tamaki, T., Itoh, K., Nishii, J., “Space-selective laser joining of dissimilar transparent materials using femtosecond laser pulses”, Appl. Phys. Lett. 89, 2006, p. 021106. 91. Tamaki, T., Watanabe, W., Itoh, K., “Laser micro-welding of transparent materials by a localized heat accumulation effect using a femtosecond laser at a 1558 nm”, Opt. Express, 14, 2006, p. 10460. 92. Bovatsek, J., Arai, A., Schaffer, B.C., “Three-dimensional micromachining inside transparent materials using femtosecond laser pulses”, New Applications, presented at CLEO/QELS and PhAST 2006, California, 2006. 93. Miyamoto, I., Horn, A., Wortmann, D., Gottmann, J., Yoshino, F., “Fusion welding of glass using femtosecond laser pulses”, J. Laser Micro/Nano engineering, 2, 2007, pp. 57–63. 94. Miyamoto, I., Herrmann, T., “Characteristics of internal melting of glass for fusion welding using ps laser pulses with average power up to 8W”, Proc. LPM2007, 2007. 95. Lecler, S., Takakura, Y., Meyrueis, P., “Properties of a three-dimensional photonic jet”, Optics Letters, 30, 2005, pp. 2641–2643. 96. Pendry, J.B., “Negative Refraction Makes a Perfect Lens”, Physical Review Letters, 85, 2000, pp. 3966–3969. 97. Carslaw, H.S., Jaeger, J. C., Conduction of Heat in Solids, 2nd edn, Clarendon Press, Oxford, UK, 1959, pp. 256–257 (instantaneous), p. 261 (continuous). 98. Ibid., p. 258 (instantaneous), p. 261 (continuous). 99. Rosenthal, D., “The theory of moving sources of heat and its application to metal treatments”, Trans Am. Soc. Mech. Eng., 68, 1946, pp. 849–866. 100. Eagar, T.W., Tsai, N.S., “Temperature fields reproduced by traveling distributed heat source”, Weld. J., 62, 1983, pp. 346s–355s. 101. Carslaw, H.S., Jaeger, J.C., Conduction of Heat in Solids, 2nd edn, Clarendon Press, Oxford, UK, 1959, p. 279. 102. Ibid., p. 267. 103. Goldak, J., Chakravarti, A. Bibby, M., “A double ellipsoid finite element model for welding heat sources”, 1985 IIW Doc Number 212-603-85.
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15 Electron beam microwelding G A K N O R O V S K Y, Sandia National Laboratories, USA, T D O R F M Ü L L E R , U D I L T H E Y and K W O E S T E , RWTH – Aachen University, Germany
15.1
Introduction
Electron beam technology is based on the fact that charged electrons accelerated by a high voltage, suitably focused and accurately directed, can be used as a tool for material thermal treatment, especially welding. Nowadays, electron beam (E-beam) welding is firmly established in many industries and is generally accepted for its reliability and efficiency. The range of joining tasks reaches from welding foil in thicknesses of just a few micrometers right up to thick plate welding with achievable weld depths of more than 150 mm in steel alloys and more than 300 mm in aluminium alloys. Moreover, almost all electrically conductive materials are weldable, and many may be joined in dissimilar material combinations. The high power density (up to 107 W/cm2) available with E-beam welding and the resulting weld depth-to-width ratio (up to 50:1) allow a large variety of possible applications of this joining process. The development of highly integrated micro systems will have to be accompanied by parallel advances of requisite micro assembly and joining techniques. Frequently, existing technology from the field of macro applications is used for this purpose. With careful consideration of the specific demands, this technology can, by downscaling, often be successfully utilized for the assembly of hybrid microsystems. E-beam welding is an example of one such macroscale technology gaining in importance for micro range applications. The ability to precisely focus the beam to a few micrometers in diameter, makes the application of this method particularly interesting for microdimensions. Another advantage of this joining technology is the almost inertia-free manipulation of the electron beam via electro-magnetic coils. This allows extremely rapid beam movements combined with very accurate control of the energy input and its precise location, which promise high welding productivity while maintaining high quality levels. Finally, because techniques developed in the electron optics industry for high resolution, high depth-of-field imaging are easily and naturally combined with those needed 418 WPNL2204
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for precise welding, the micro electron beam welder becomes an ideal instrument for assuring quality control of the welds produced, by providing an in-situ visual inspection capability. Recent progress in the fields of control technique and hardware performance has extended the multitude of E-beam welding applications; there is little doubt that more advances and broader application of E-beam welding, especially in the microworld, are in the offing. In this chapter, after a brief description of the state-of-the-art of conventional macroscopic E-beam welding, we will describe the similarities and differences in its application to the micro- and nano-worlds. We will present work describing how the process is characterized, how the beam interacts with the workpiece, and practical suggestions for fixturing and part manipulation. Along the way we will provide illustrative examples.
15.2
Technology
15.2.1 Electron beam generation and guidance Generation of the electron beam In modern E-beam welding machines, triode systems are invariably used for beam generation. These systems are composed of an anode, a cathode and a control electrode (Wehnelt cylinder). The electrons which compose the beam are emitted from the cathode by thermionic emission. The cathode is made out of a material with a small work function (the amount of energy needed for an electron to be emitted from the material). According to Richardson’s Law, the current density of emitted electrons exponentially increases with temperature. Therefore, the cathode material must not only show a high electron emission rate (i.e. have a low work function) but also has to be high-temperature resistant to guarantee a relatively long cathode life. Appropriate materials are tungsten and tantalum. Additionally, the mass of the cathode is kept low in order to rapidly cool down after switching off the heating current; this reduces oxidation of the cathode in a system which must be frequently and rapidly vented to atmosphere [1]. The heating of the cathode may be direct or indirect. Directly heated cathodes are heated by Joule (resistance) heating. Indirectly heated cathodes are heated by electron bombardment from an auxiliary cathode. By the application of a high voltage electric field between the anode and cathode, electrons are first extracted from the electron cloud around the cathode and subsequently accelerated to attain a high value of kinetic energy [2]. Depending on the magnitude of the applied voltage, the electrons may be accelerated up to 2/3 of the speed of light. By applying a control voltage between the cathode and a control electrode, a barrier field is generated which forces some of the emitted electrons back
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to the cathode. Thus the beam current is regulated by this control voltage. Decreasing the magnitude of the control voltage allows more electrons past the barrier field towards the anode. Due to its particular shape, which can be compared to a concave mirror used in light optics, the control electrode effects an electrostatic focusing of the electron beam. In this manner, a beam of several tens of mA with an accelerating voltage of up to 150 kV is generated. Having passed the anode, the electrons have achieved their final speed; subsequently the electron beam is focused and deflected by means of electromagnetic focusing lenses which do not effect the beam’s kinetic energy. The focusing effect leads to the constriction of the electron beam, the socalled “crossover” condition. Beam manipulation The electron beam, which diverges slightly after having passed the pierced anode is focused to a spot diameter of from 0.1 to 1.0 mm by the following beam manipulation system in order to reach the necessary power density of 106 to 107 W/cm2. First, the beam is guided through an alignment coil onto the optical axis of the focusing objectives. One or several electromagnetic lenses aim the beam at the workpiece inside the vacuum chamber. A focusing coil adjusts the focal distance and therefore the beam’s effective diameter on the workpiece surface. Several E-beam welding machines even use a double focusing system to realize even smaller beam diameters and higher power densities. A stigmator lens which consists of two or more pairs of opposing magnetic coils is included which compensates magnetic and electric perturbations on the beam and corrects unsymmetrical beam geometries. Because of the electric charge of the electron beam, it can easily be aimed by magnetic fields, generated by deflection coils. The very low mass of the electrons (only 9.1×10 –28 g) enables virtually inertia-free, very rapid manipulation of them. Perpendicular axis deflection coils which are positioned at various parts of the electron beam column assist in the creation of various programmed motions of the electron beam around its aim point. Circular, bar, and figure-8 patterns are common, but any arbitrary pattern is possible, depending upon how the voltages to the deflection coils are programmed. A schematic representation of a conventional E-beam welding machine is depicted in Fig. 15.1.
15.2.2 Electron beam welding machine peripherals The E-beam welding machine is composed of a multitude of individual components. The basic component of the machine is the electron beam generator where the electron beam is generated in high vacuum, influenced by electromagnetic deflection coils and then focused onto the workpiece in the
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High voltage supply
Working chamber
Beam forming and guidance
Beam generation
Cathode Control electrode Anode Adjustment coil Valve
To vacuum pump
Viewing optics
Stigmator Focusing coil Deflection coil
Workpiece Workpiece handling
To vacuum pump
Chamber door
15.1 Schematic diagram of a conventional E-beam welding machine.
vacuum chamber. As the electron beam in free atmosphere diverges strongly due to collisions with air molecules and thus loses power density, welding is generally carried out in high vacuum inside a working chamber. Different vacuum pumps are typically used for vacuum generation in the beam generator and in the working chamber. While in the beam generator a high vacuum (p < 10–5 mbar) is indispensable for prevention of electrical breakdown (arcing or “flashover”) and oxidation of the cathode, the possible range of allowable pressures in the working chamber varies between moderately high vacuum (p < 10–4 mbar) and atmospheric pressure. Besides the above-mentioned modules, a high voltage supply and its controls, vacuum pumps, a workpiece motion system (rectilinear and rotational) and its controller, and an operator interface are necessary. The equipment is controlled from the operator interface console where all relevant process parameters are set and monitored. For the determination of optimum welding parameters for process control and also for the adjustment of the electron beam on the workpiece, a workpiece viewing system is used. While traditionally
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these viewing systems were optical in nature, recently electron-optical systems have been introduced where the electron beam scans the workpiece surface with a very low power. The secondary electrons show, as in an SEM, an image of the workpiece surface.
15.2.3 Conversion of an SEM into a micro electron beam welder Column modifications For the practical realization of a suitable machine for welding of micro components, conventional E-beam welding machines may be eliminated. Their beam powers, which lie between 100W and ~5 kW are simply too high for the welding of microscale components. However, simple calculations show that the use of a scanning electron microscope (SEM) as a microwelding tool is much more promising. This concept combines two basic functions in a single tool: observation and welding. The SEM which has been modified at the Welding and Joining Institute (ISF, RWTH Aachen) is made by LEO Zeiss company and has a maximum beam generator power of 6 W. Power losses which are due to screening by apertures and also beam scattering in the column reduce the power down to approximately 3–4 W at the workpiece. For a 30 kV accelerating voltage, this corresponds to ~100 µA probe current. Figure 15.2 shows the simplified beam path of the converted SEM in both operating conditions: (a) observation, and (b) welding. The integration of d0
d0 Crossover
Condensor lens 1
d1
fK1 fK1
Condensor lens 2
d1
d2 Aperture
Objective lens
dF
Beam focus
Observing mode
dF Welding mode
15.2 Modification of the SEM electron optics for welding operation.
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both functions in a single piece of equipment puts, for the most part, opposed process demands on the equipment. The observation and the analysis of substrates require the lowest possible energy input and a high resolution. Multiple small diameter apertures for the screening of off-axis electrons and two condenser coils for the reduction of excess electrons in the beam provide an extremely low power beam with extremely small focused spot diameter on the sample surface. In welding, on the other hand, a power which is in most cases higher by several orders of magnitude than that for observing must be provided. The modifications needed to convert to a welding mode relate to changes of the beam column components. Among those alterations is the removal of two apertures at the lower end of the liner tube, right below the second condenser lens. These apertures are normally necessary to prevent stray electrons from contaminating the lenses. In the welding mode, these apertures reduce the effective beam-power. Therefore, the removable endpiece of the liner tube, containing these two apertures, is replaced with an “open” endpiece. Additionally, the objective (final) aperture is removed, which is fitted onto a slider which can be easily pulled out of the beam path. Also, the lower condenser lens is switched off before starting the SEM operating system. As those modifications are reversible, the existing equipment may be used as both an observation tool and as a welding tool, though the removal/ insertion of the liner tube apertures requires opening the column. For those less willing to modify their SEM, it has been found that simply removing the objective (final) aperture (which is usually done by an alreadyprovided micrometer adjustment knob), adjustment of the condenser lens currents to maximize the beam current and using a commercially-available 2% thoriated W cathode filament (instead of the usual pure W) can provide ~1 W of beam power, certainly less than that achievable by the above modifications, but sufficient for smaller applications. The advantage here is that the column does not need to be opened to switch between modes. Positioning unit modification As subsequent examples welded with the micro E-beam welding system will show, varying arrangements of micro components may be joined to one another. The quality of a micro E-beam welded joint is found to be strongly dependent upon (a) the accuracy of the joint preparation, and (b) the adjustment precision of the components which are to be joined. For instance, a slight angular deviation of two components which are to be joined with a square butt joint leads to a gap which is no longer bridged by the minute beam diameter of just a few micrometers. Faulty joints are a consequence. Due to the excellent imaging provided by the SEM, highly precise part alignment may be carried out in its working chamber if the existing sample manipulation system is supplemented with a second independent five-axis manipulator.
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Such an addition allows correction of the component alignment without repeated withdrawal of the components from the working chamber. The adjustment device is composed of three linear adjusters, two tilting axles and a control unit. The vacuum suitability of the components allows the maintenance of the vacuum which is necessary for the welding process and also the reliable operation of the motors. Self-locking of the gear drives allows one to switch off the controls after the desired position has been reached, which also leaves the electron beam uninfluenced by the motors’ electromagnetic fields and vibrations. When choosing and integrating the supplemental adjustment unit into the vacuum chamber of the SEM, several essential factors must be considered: • • • • • •
vacuum suitability of the mechanical and the driving components high plane-parallelism of adapter plates collision avoidance with wall, electron gun and detectors during operation central positioning of all axes below the exit outlet of the electron beam column z-position of the joining plane must be located in the region of the working distance of the electron beam vacuum feedthrough for the electrical wiring of the motors.
The system design shown in Fig. 15.3 allows one workpiece component to be moved by the five-axis-positioner independently from the second. By using the additional rotation axis of the SEM’s own positioning system ΘZ this component offers six motional degrees of freedom. Additionally, the entire assembly can be aligned with the beam axis by using the SEM’s existing positioning system. The tilting actuator provides inclination in two axes in order to realize contact pressure or to compensate improper angle positions. Together, the two units provide a flexible positioning- and handlingsystem, essential for processing complex joining problems. In addition to the stage motion, control of the beam motion is equally important. A beam controller modification will be treated in a later section.
15.2.4 Beam characterization methods (Faraday cup) Consistent reproducible quality of a weld is defined by the workpiece dimensions, fitup and material properties, and by the properties of the electron beam. Similarly to the field of metal-cutting manufacturing, a precise characterization of the tool is needed to draw valid conclusions from phenomena observed during the welding process. The most important parameters of the electron beam are beam diameter, peak power density and the distribution of power density. Several methods for the measurement of these properties are not suitable if the beam is not rotationally symmetric. Assuming that the beam in not
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Chamber
Gun Clamping device 2 goniometer stages (positioning system) Θ x, Θ y
Θ
z y
Θx
x Support of the stationary clamping device
y
Z
x-, y-, z-axes (positioning system)
Y
Carrier
X ΘZ
3 linear stages and 1 rotating stage (SEM), X-, Y-, Y-axes, ΘZ
15.3 Positioning unit. Top: diagrammatic representation, bottom: installation in the SEM chamber.
symmetric, the characterization is possible by using an apertured sensor. This sensor consists of a metallic casing with an electrically insulated Faraday cup inside. On top of the casing above the Faraday-cup, a platinum faceplate with a hole of 10 µm in diameter is embedded. A part of the beam electrons
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can pass through this hole and impinge on the Faraday cup. The diameter of the hole was chosen to be smaller than the smallest beam diameter at the focus point with the three apertures mentioned earlier removed. Beam diameters for currents ranging from 75 to 175 µA were measured. Both the signals generated by electrons hitting the Faraday cup as well as the casing are recorded for every X- and Y-position while scanning the sensor line by line. By logging and evaluating the signals for every beam position, a threedimensional image of the actual power density and distribution can be generated. The measurement technique and an example of the results obtained are depicted in Fig. 15.4.
15.2.5 Potential of fast beam controls and multi-beam technique As the electron beam is readily deflectable by relatively low current magnetic or no current electrostatic methods, it is possible to move the aim point of the beam with extremely high rapidity, far faster than the analogous beam scanning technique used in laser marking where galvo-controlled mirrors with small but finite mass and angular momentum are employed. Recent developments in the field of beam deflection allow varying the focus position rapidly; by simultaneous application of a suitable control electrode voltage, the beam intensity between the individual beam traces can be controlled (or even switched off) as well, effectively skipping the beam between several positions with a speed so high that the thermal influence on the structure is carried out at different points essentially simultaneously, effectively simulating a parallel rather than serial process. In Fig. 15.5 (left), five electron beams are pictured as simultaneously processing the material. This technique promises a future
Faceplate
355.0 310.0
Aperture 350 300 250 200 150 100 50 0 Direction of –30 beam deflection X- –20 co –10
Sign
265.0 220.0
V) al (m
175.0 130.0 85.00 40.00 –5.000
(µ m
)
30 20 10 or 0 h = 21.6 mm di 0 na 10 –10 Uam = 30 kV te 20 –20 (µ Iam = 125 mA m 30 Inl = 3.60 A 40 –30 ) Ikt = 270 mA Iobj = 1270 mA
Y–
co
or
di
na
te
Cross-section of electron beam in measuring layer
15.4 Measuring principle of the (apertured) diaphragm sensor and an example of results obtained at a beam current of 125 µA.
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15.5 Multi-beam technique.
potential for a multitude of applications, particularly in the world of integrated circuits, where the ability to parallel process on a wafer scale is highly desirable. The simplest application of this technique is to skip the beam alternately between two (or more) positions on an independently moving workpiece, thus producing two (or more) welds. For many years now, saw bands for band saws in conveyor units have been joined this way. These saw bands consist of a ductile backing layer in the middle and two hardened toothed edges, i.e., two parallel welds are necessary. A further interesting field of application for the multi-beam technology is the welding of concentric bodies, such as a gear to a shaft. As the material melts (and then solidifies) simultaneously at several points along the circular faying surface, the shrinkage stresses also occur symmetrically, thus avoiding misalignment of the axes. Through this, often costly and labour-consuming press fits for centering and avoiding axial misalignment may be dispensed with. Another application of this fast beam deflection is joining of material combinations. The multibeam technique allows, by varying holding times at different points, to supply one of the workpieces with a significantly higher energy than is the case for the second. For example: one workpiece is melted while the other one is not (effectively a braze). In this manner, it is possible to join material combinations where the joining members are metallurgically incompatible. This technique could also be used when the two pieces have very different thermal properties. When the multi-beam technique is not possible, offset of the beam impact point from the joint groove is required (and must be carefully determined). This demands the extremely precise positioning of the beam and is thus not easy to reproduce in high rate production.
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15.2.6 Clamping and guiding devices An often-occurring phenomenon during the welding of microparts is the failure of effectively joining the parts, due to either misalignment or the lack of sufficient material to bridge the welding gap. The consequences of surface tension of the melt and the capillary pressures which develop in the molten zone (p = 2σ/r, where p = pressure, σ = surface tension, and r is the radius of curvature of the molten surface) must be considered when microparts are joined. Both are relatively high because of the small dimensions (and resulting high radii of curvatures) of the joined parts. For example, with a typical metallic value for σ of ~ 1 N/m, at a radius of curvature of 100 µm, p = 20 kPa; at 1 µm, p = 2 MPa ~ 20 bar! When the melt forms, it attempts to minimize its surface area and tends to ball up, which tends to increase the pre-existing gap; then when cooling, thermal contraction causes it to retract from the joint. Thermal contraction can generate very large forces as well. One can readily estimate that under rigid restraint conditions, only about 100 K temperature change is needed to drive most metals to their yield point. This is particularly noticeable when joining wire-shaped geometries, as depicted in Fig. 15.6. This effect makes the process inconsistent and prevents sound conclusions when researching the joinability of different materials. By applying a compressive preload, both effects can be mitigated, however. Therefore, the use of clamping and guiding devices is essential to obtain reproducible welding results. At the ISF, for the welding of wire-shaped parts an assembly as shown in Fig. 15.7 is used. Like the components of the positioning unit, the assembly must be made of vacuum suitable and non-magnetizable materials and is often made of aluminium. The relatively high thermal conduction of aluminum and the limited beam power of the SEM can prevent the parts from becoming molten during the process, if the stick-out is insufficient. Of course, too much stick-out results in a buckling instability condition. The two clamping jaws of the device can be precisely positioned with the positioning units. While one jaw has a fixed position (1), the other is mounted on a linear slider (2) connected to an adjustable spring (3) to provide the proper preload and tracking during the welding process. A significant increase in reliability and reproducibility of the joining process is obtained by using this assembly. However, excessive preload also causes faulty welds. A similar assembly is used for the welding of sheet material, though the preload and guiding are less important than the prevention of alignment failures due to heat distortion during the process. These can be counteracted by using the multi-beam technique.
15.2.7 Waveform generator and teach-in programming In a standard E-beam welder, welds are typically produced by moving parts underneath the beam, which except for localized beam deflection is stationary.
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Ck 101 φ 175 µm
Ck 101 φ 175 µm
15.6 Joining of wire-shaped geometries.
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Ck 101 φ 175 µm
Excessive preload
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15.7 Guided clamping device with adjustable preload: 1. fixed jaw, 2. sliding jaw, 3. adjustable spring load.
SEMs vary greatly in their specimen stage motion control capabilities, but none have the ability to control the path and velocity needed for a joining process. Even if a CNC-controlled specimen stage were available, it would not have the flexibility and range of speeds accessible by beam scanning. As the images in a SEM are generated by rastered beam scanning of the specimen being observed, the capability of directing the beam over a wide area is inherently present. However, the possible heat source patterns created with a converted SEM are restricted to rectangles, straight lines and spots. Modification of the beam scanning control system with the addition of a programmable waveform generator is required for realization of arbitrary spot and seam geometries. By means of such an addition, the ISF micro Ebeam welding system can select standard geometries such as line, circular or ellipsoid segments. In addition, user-programmable functions are possible which allow consecutive welding of multiple parts with the same beam path or single parts with a complicated combination of weld paths. Another advantage of the beam scanning control system is that it enables the multibeam technique mentioned earlier, where if the beam deflection between different points is carried out rapidly enough, the simultaneous welding of several joining points is possible. For example, with a circular weld path geometry, the local heat input may be distributed symmetrically around the component which, as a rule, results in less distortion [3]. A waveform generator made by Pro-Beam AG & Co. (Fig. 15.8) converts and amplifies a digital input signal into analog voltages which drive the deflection coils and other beam controls of the SEM. This waveform generator is capable of controlling three output channels. Two channels are used to control the X- and Y-deflections for the positioning of the beam. The third channel is used either for control of a focusing lens (focus shift) or for the
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15.8 Waveform generator for beam pattern control.
control of the emission current (power modulation). Thus, the input thermal power can be adjusted either via beam defocusing or through the alteration of the beam power. The generator software which creates the digital input signal offers, moreover, the ability to modulate the beam position as it travels along the programmed welding path. This technique of “local” beam pattern motion is often used in conventional E-beam welders where small diameter circular or cross-weld bar patterns are commonly used. The local beam motion allows joining across gaps which are wider than the focused beam diameter, or can be used to improve the cosmetics of a rough as-welded surface (examples of this technique will be given in a later section). The signal flow connections of the generator components are shown in Fig. 15.9. At present, file management for the beam control generator is carried out with a Windows-based application program. The data format is ASCII, and the control files can be written with programs such as EXCEL or MathCAD. Alternatively, an SEM image of the components to be joined is loaded into the control software. The operator interface then allows teach-in programming of arbitrary travel paths through imposition of splines, lines, circles and free points. Input of power data and the determination of start and stop locations of the welding process are also provided. All these data are then converted into an ASCII file which is compatible with the control software of the waveform generator. Figure 15.10 depicts the process of this teach-in programming for a randomly chosen beam path.
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Digital/analogconverter
Function Amplitude Vectorization DC-offset generation adjustment
DA
X
DA
Y
X RAM Y
Function Amplitude Directional X /Y number vector and frequency
DC-X/–Y
Function generation
RAM
Function number and frequency
15.9 Signal generation for beam motion control. The top row of signal control blocks refers to the overall beam path; the bottom row refers to the local beam pattern.
As these are comparatively time-consuming and laborious processes which poorly utilize the SEM’s capability for observation of the workpiece, a software module is currently being developed which will make motion control of the beam and SEM operation much more user-friendly.
15.2.8 Energy transfer with the workpiece and penetration effect When the beam of electrons strikes the workpiece, its kinetic energy is mostly converted into thermal energy. The high electric voltage potential of the electrons allows, at an accelerating voltage of 150 kV, electron acceleration up to a speed of approximately 2×108 m/s, which is two-thirds of the speed of light. Not all beam electrons penetrate into the workpiece and release their energy to the material. A part of the striking electrons are emitted in the form of back-scattered electrons, secondary electrons and X-ray radiation
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15.10 Teach-in programming of beam pattern generator.
(Fig. 15.11). Additionally, some energy is lost to thermal radiation and by evaporation of vaporized material. Where the melting temperature is exceeded, a molten zone is formed. The coalescence of the separate pools created in the workpieces to be joined across the initial gap and subsequent solidification after removal of the beam is what creates the weld. For the accelerating voltages employed in a typical E-beam welder (60–150 kV), the electrons only penetrate ~150 µm, due to their low mass. This is referred to as conduction mode welding. In order to obtain large weld depths, a special effect is needed, the socalled deep penetration or keyhole effect. At sufficiently high beam power density, material is superheated and vaporized by the beam despite heat dissipation into the surrounding cold base material. This vapor may exceed a temperature of 2700 K in ferrous materials. At this temperature, the vapor pressure is sufficiently high to locally depress the molten metal surface, forcing molten metal outwards and upwards, producing a vapor cavity (the keyhole) which in its core consists of superheated vapor and is surrounded
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Back-scattered electrons
X-ray
Secondary electrons
Thermal radiation
Convection
Heat conduction
15.11 Energy transfer of beam with the workpiece.
by a shell of molten metal. This effect is maintained as long as the pressure from the developing vapor cavity and the surface tension of the molten pool are in equilibrium. The diameter of the vapor cavity corresponds approximately with the electron beam diameter. With a sufficiently high energy beam, the developing cavity penetrates through the entire workpiece [1]. The relative motion between workpiece and electron beam causes the material which has been molten at the front of the electron beam to flow around the cavity and to solidify at the backside. The formation of the vapor cavity is depicted in Fig. 15.12 [1, 2]. The pressure and temperature conditions inside the cavity and its instantaneous shape and depth are subject to extremely rapid changes. As a result of the dynamically changing geometry of the vapor cavity, welding faults such as porosity may occur when the welding parameters have been unfavorably chosen. Those faults may be avoided by the modification of a series of welding parameters and, in particular, by the selection of suitable beam deflection characteristics, as, for example circular, sinusoidal, rectangular or triangular functions. For the case of micro E-beam welding, the beam accelerating voltage is reduced to that typical of an SEM (1–40 kV) and of course is reduced in current from tens of mA to ~100 µA (up from nA in observation mode). At the same time, however, the spot size is considerably reduced from the 0.1– 1 mm diameter typical of a conventional E-beam welder to about 20 µm. Thus while the overall power is down many-fold, the power density of the
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Formation of a vapor capillary
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Solidified weld seam
15.12 Deep penetration or keyhole effect.
micro E-beam welder can be even greater than that of a high power conventional E-beam welder. Since capillary forces are inversely proportional to the radius of curvature of the liquid surface, and the keyhole size is comparable with the beam diameter, the capillary forces encountered in micro E-beam welding are increased substantially (5–50 times), so the ability to achieve a stable keyhole is lessened. Fortunately, in microjoining there is no need for keyholing, since we are dealing with small, thin workpieces. For this reason, in what follows we shall only treat the electron beam in non-keyhole mode.
15.2.9 Modeling of beam/workpiece interaction (Monte Carlo) Beam penetration and energy distribution calculations As noted above, when an electron beam strikes a sample, the beam penetrates below the surface, and is “diffused” by a variety of mechanisms [4]. (Note that this is unlike the case of a laser beam, which is adsorbed with a few atomic layers of the surface.) Simulation programs for the beam’s penetration behavior employing Monte Carlo techniques are publicly available. These have been used for many years for the simulation of electron trajectories in electron microscopy applications [5]. By employing one of these programs (we gratefully acknowledge the use herein of the CASINO code by investigators at the Université de Sherbrooke, Canada [6]), a great deal of insight into the beam’s behavior on µm-size samples becomes possible. While the accuracy of these calculations depend strongly on the models adopted for the scattering of electrons by atoms which then determine the cross sections for scattering into a particular angle, these cross sections have been extensively studied and are generally considered quite good for the medium accelerating voltage (AV) regime of 5–100 keV. The calculations’ accuracy is most often stated to
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be better than 5% relative, although the errors can sometimes be as high as 20% depending on the type of measurement used [7]. The electrons’ energy was tracked from its original value down to 0.05 keV. Otherwise the CASINO default model settings were used. In Fig. 15.13, the distribution of 500 electron trajectories (projected onto a plane) for two different size beams striking a Ni specimen are shown. In Fig. 15.13(a) a 100 nm diameter focused beam is used, and in Fig. 15.13(b) a 1 µm beam is shown, both at 30 kV AV. It is noticeable that for an order of magnitude change in the beam diameter, the outer contours of the electron trajectories are very similar. That is, the electronic penetration controls the effective size of the beam once it hits the sample far more than the focused diameter. Also noticeable is that the electrons travel in a roughly hemispherical region of about 1 µm radius. When a 10 kV AV is used instead, the behavior noted in Fig. 15.14 is seen (the same number of trajectories are plotted as in Fig. 15.13, and the same size scales for the axes are used). With the lower AV, the effect of focusing of the beam at these diameters is significant, since the electron “diffusion” distance is now less than the beam diameter, and the penetration depth is likewise greatly reduced. Note that since the same number of electron trajectories is plotted in Figs 15.13 and 15.14, they can be considered to represent a constant beam current. (In reality the CASINO calculation does not take time into consideration.) The electrons still penetrate a great deal more deeply relative to a laser’s photons, but it is clear that the concept of a surface heat source is acceptable, particularly for the beam diameters appropriate to micro E-beam welding (typically several times greater than 1 µm of the above examples). For the 10 kV example, though the AV has been reduced by a factor of three relative to the original 30 kV, the projected area of the “hemisphere” has decreased by a factor of ~4, and its depth by ~6 (for a volume reduction of ~24-fold), so the intensity (energy/volume) has actually increased by ~[10/30/(1/24)] = 8X. Further decreases in the AV will further increase the surface nature of the heat source. The maximum AV of commonly available SEMs is 40 kV (more often 30 kV). The beam trajectories calculated for this are given in Fig. 15.15. It is clear that a significant fraction of the electrons penetrate in excess of 2 µm. Furthermore, this calculation is for Ni, which is a relatively dense material of moderate Z. Many possible applications for µEB welding will involve Si, in particular MEMS devices, and Si has a much lower density and Z, which results in much deeper penetration for a given AV. MEMS devices employ multiple layers that are ~2 µm thick. If two parts in the same layer are to be joined, using a 40 kV AV would be contraindicated, as a significant portion of the beam would basically pass right through the layer to be welded, as shown in Fig. 15.16(a). On the other hand, if parts from two different layers are to be joined, perhaps in a lap joint, then the more highly
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(a)
(b)
15.13 Electron trajectories calculated for 30 kV beams of (a) 100 nm diameter and (b) 1 µm diameter in a Ni target. The black trajectories are back-scattered electrons, the grey are the primary beam electrons. 500 electron trajectories have been plotted.
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(a)
(b)
15.14 Electron trajectories calculated for 10 kV beams of (a) 100 nm diameter and (b) 1 µm diameter in a Ni target. Plotted to same scale as Fig. 15.13. The black trajectories are back-scattered electrons, the grey are the primary beam electrons. 500 electron trajectories have been plotted.
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15.15 Electron trajectories calculated for 40 kV beam of 1 µm diameter in a Ni target. The black trajectories are back-scattered electrons, the grey are the primary beam electrons. 500 electron trajectories have been plotted.
(a)
(b)
15.16 Electron trajectories (primary beam electrons only, 500 trajectories, note that figure is enlarged ~2× relative to Fig. 15.15) calculated for 40 kV beam of 1 µm diameter in (a) Si target 2000 nm thick, (b) Si target with two layers, each 2000 nm thick. Note that all trajectories are visible, i.e. they are projected onto the plane of the figure. WPNL2204
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penetrating AV might be useful, depositing significant amounts of its energy below the faying surface as in Fig. 15.16(b).
15.2.10 Conversion to energy vs position (surface vs volume heat source) While the electron trajectories are indicative, they do not show where the energy is deposited, or how much. Instead energy contours are needed. The energy contour plot corresponding to Fig. 15.16(a) is given in Fig. 15.17(a), and displays contour lines which show the percentage of the beam’s energy not yet absorbed. That is, the 10% contour shows the volume within which 90% of the beam’s energy has been absorbed. As suggested by the electron trajectories, a significant portion of the beam’s energy actually penetrates completely through the top layer. The energy contours are nearly cylindrical, and the spacing of the contours along the beam centerline indicates that the energy is deposited in a nearly linear manner except at the very top and bottom. For modeling purposes, a cylindrical broadened line source approximates these conditions reasonably well. For the two Si layer case of Fig. 15.16(b), while many of the electrons’ trajectories penetrate into the second layer, the energy deposition contours shown in Fig. 15.17(b) indicate that less energy penetrates to the second layer than might have otherwise been suggested by the trajectories. However, if instead of a relatively thin second layer, a thicker or even ‘massive’ substrate of Si is present, the energy absorbed in the top layer is higher because electrons which would have exited from the bottom of the second layer have a finite chance of being scattered back upwards into the top layer. This is shown in Fig. 15.18, where the interfacial energy density contour is now ~50% of the original beam intensity. When comparing Figs 15.17 and 15.18 it must be remembered that a larger amount of the beam’s energy has been absorbed in the latter two cases, so one cannot directly compare the same percentage contour lines. As will be seen later, the percentage amount of the total beam energy absorbed by a 2000 nm layer of Si only amounts to 12%, whereas for the 4000 nm layer the corresponding amount is 27%, and for a “thick” layer of Si, it would become slightly over 90%. Thus, to perform even a qualitative “equienergy” deposition comparison between the various thicknesses, the contour line value must be scaled by the total energy absorbed. As an example, in Fig. 15.17(a), the 25% contour line (75% deposition) nearly reaches the back surface. In Fig. 15.17(b), the 25% contour line has actually spread to the second layer, implying that a slightly higher percentage contour (and thus slightly lower energy deposition percentage) would just reach the back side of the top layer. In Fig. 15.18 the 50% line just reaches the interface. Thus to compare the energies deposited along the beam axis in the top layer requires comparing three different contour lines, of ~25%,
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(a)
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15.17 Electron energy contours through the specimen in (a) Fig. 15.16(a) and (b) Fig. 15.16(b). The contours show the percentage of the beam’s energy not yet absorbed. These figures are vertical slices (184 nm thick) through the specimens in Fig. 15.16.
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5.0% 10.0% 25.0% 50.0% 75.0% 90.0% ~ 4257.2 nm
0.0 nm
4257.2 nm
15.18 Electron energy contours. This figure is a slice through a specimen like that in Fig. 15.16(b) except the lower layer is 3000 nm thick.
~30% and ~50%, which represent ~75%, ~70% and ~50% of the total energies deposited, which are 12%, 27% and 38% of the beam energy, respectively. Thus the relative energy depositions above the interface actually are: in the no lower layer case: 0.75 × 12 ~ 9%; the 2 µm thick lower layer case: 0.70 × 27 ~ 19%; and the 3 µm thick lower layer case: 0.50 × 38 ~ 19%. Adding a second layer doubled the absorption into the top layer, but thickening the lower layer further had essentially no effect. A similar effect occurs if the lower layer is a denser material. For the lower layer, whatever the thickness, the cylindrical line source no longer makes much sense, as the contours are not even close to cylindrical. A layered elliptical geometry approximates the situation more closely.
15.2.11 Thermal effects in microscale parts (FEM calculations) Surface vs volume heat source A commercial FEM code (COSMOS/M V2.9, Structural Research and Analysis Corporation, Los Angeles, CA. USA) was used to calculate thermal contours
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of heat sources approximating surface and broadened line sources in MEMSlike tensile bars of Ni or Si. For the case of Ni a single bar was modeled, while for Si, two bars were overlapped. The results of the calculations are shown in Fig. 15.19. Each bar was 2 µm thick × 10 µm wide × 240 µm long. Because of symmetry, only 1/4 of the actual volume was modeled. For purposes of simplicity, only conduction heating using constant ambient values for the thermal properties (the heat of fusion was not incorporated) was modeled. All surfaces were treated as adiabatic except for the non-heated end, which was held at ambient. Volume heating within selected elements was used to introduce the heat into the specimen in both cases. For the “surface source” (Fig. 15.19(a)), one layer of elements 0.2 µm thick contained in an approximately 1 µm radius quarter circle was used, whereas for the “volume source” (Fig. 15.19(b)), all the elements below (and including) the set used for the surface source were used. Since the overall bar thickness was 10 times that of the surface layer, a volume heat intensity one tenth that used for the surface source (and over the same time period) gave the same total heat input. The values used were 1 or 10 × 109 Calcm–3s–1. This corresponds to an overall heat input of 60 mW. Of interest is that the surface source gives a slightly higher peak temperature and slightly steeper gradient for the same total heat input; however, considering that the heat intensity is 10 times greater, this is a very small (<3%) difference. Both exhibit very rapid evolution of the temperature field, which is essentially controlled by the ramp up of the heat source rather than by thermal diffusion (Fig. 15.19 et seq. portray the temperature at the end of the heat sources’ ramp-up and are essentially equivalent to steady state values). Given Ni’s value of thermal diffusivity of ~0.2 cm2 s–1, the diffusion “distance” (4αt)1/2 in 1 ms is ~300 µm. Under these thermal conditions, which approximate a stationary rather than a transient solution, the distance to the nearest heat sink controls the required heat input necessary for melting. A slightly smaller heat input would suffice to initiate melting at the end of the bar, as the melting isotherm, given approximately by the arrow, is ~20 µm wide. In the case of two layers of material, if the direct beam cannot penetrate through the top layer, then transfer of heat across the interface will limit the heat into the bottom layer. Modeling this is difficult, because it is difficult to know the interfacial thermal resistance. As the vacuum environment essentially precludes convection, until coalescence of the two parts or keyholing occurs, the only methods of heat transfer are radiation and conduction. (Keyholing has not been noted in any of our experiments.) Both of these mechanisms will tend to produce a reduced intensity heat source in the bottom layer, and thus will have difficulty raising its temperature to the molten state. Replacing the material properties in the Ni bar model with those for Si, and adjusting the heat input/volume for the cylindrical heated volume to obtain a suitable
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444 Microjoining and nanojoining 15.19 Calculated temperature contours in °C for (a) surface and (b) volume heat sources of equivalent total heat input in Ni bar. Arrows indicate equivalent distance (18 µm). Note different temperature scales.
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maximum temperature gives the results of Fig. 15.20. For the bottom layer (Fig. 15.20(b)), it was assumed that the beam broadened when passing through the first layer. A reduction in the volume heat source’s intensity by a factor of 3.8 was made to offset the volume increase of the heated elements, maintaining the total heat introduced into the bottom layer equal to that in the top layer. Nevertheless, the maximum temperature in the bottom layer only reached about two-thirds of the melting temperature (note that beam heating was the only source of heat introduced into the bottom layer). Similar calculations for thicker and wider Si bars (simulating a substrate) indicate that an energy source of approximately 10 W will be needed to achieve fusion.
15.3
Microwelding process
15.3.1 Process sequence As the modifications of the SEM are reversible, the functionality of the SEM as an analysis device is still available. Therefore, positioning, joining and analysis are possible in one piece of equipment, without the need of transportation or protection against contamination or oxidation between the joining process and the analysis. Figure 15.21 shows the chronological sequence of the welding process: as a first step, the components to be joined are adjusted exactly in relation to each other by using the two positioning systems in the vacuum chamber. Then, the electron beam is positioned on the joint. After the changeover to the welding mode the actual welding operation starts where, however, online process observation is not yet possible. After the welding process has been finished, the joining point may, after changeover to the observation mode, be subject to further analyses or measurements. All welding sequences in practice employ this step sequence.
15.3.2 Basic beam process variations The modified SEM offers the possibility of choosing from among several basic beam process variations, shown in Fig. 15.22. They basically differ in the type of beam manipulation on the substrate. In single scan mode the electron beam is guided once over the welding zone with a fixed welding speed. During multiple scan mode, it is oscillated back and forth for a defined period of time with a pre-selected deflection frequency. The method selection depends on the joining task. Prior tests showed that single scanning for the joining of foils led to very clean, sharply delineated weld edges without weld notches. Multiple scanning, however, showed weld edge regions which were remarkably uneven. The beam energy, which was being absorbed over a
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15.20 (a) Thermal contour (and location of heated elements) of top layer, (b) thermal contour for bottom layer (and heated elements) with identical total heat input in Si bars. Note different temperature scales.
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Evaluation of the position Adjustment of the components Beam positioning
Welding process
Evaluation of the surface structure
15.21 Process sequence in micro E-beam welding.
Single scan
Inverting points
Electron beam Base material Weld seam
Scanning of a layer
Multiple scan
vs
–vs
vs
Scanned area
Y
Applied layer of solder
X
15.22 Process variations.
longer period of time led to the partial evaporation of the stainless steel foil material. A further variation of energy input is the raster scanning of a larger substrate surface. The surface area is determined by the magnification setting of the SEM. In this mode, the electron beam is applied as the heat source for a soldering process with pre-placed low-melting solder on the micro components to be joined. The E-beam’s thermal effect on these materials leads to either direct fusion of the solder or liquation of the contacting materials which result in the joining of the components. Materials with high
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thermal conductivity, in particular, or non-metal materials may, under certain prerequisites, be joined by means of micro E-beam soldering.
15.3.3 Parameter choices for micro electron beam welding process schedules When developing a traditional arc or laser welding schedule, a number of questions are typically asked, then answered when choosing heat input and travel speed (or heat source on-time if a spot weld). Among the most important of these are: what material do I have, how thick is it, what is the geometry, are there any nearby features that are heat or strain sensitive (e.g. glass-tometal seals, pyrotechnic materials, etc.) and can filler material be added? While a similar process will obtain for the micro E-beam welding process, the approach is somewhat different. First, as noted above, the penetration ability of the electrons is strongly controlled by the material, so the beam’s AV should be matched to the thickness and geometry (one layer or more). After this primary concern, then issues of the heat input (controlled by beam current, travel speed or dwell time, and nearest heat sink distance) and concerns about damage to nearby features can be addressed. For a Si-based MEMS scenario, energy losses due to back-scattered and transmitted electrons were calculated [6] versus thickness and accelerating voltage and subtracted from the beam’s original energy to determine the percentage of energy absorbed. These are collected in Table 15.1. Some interesting trends may be seen. It is clear that a 1 µm thickness of Si is sufficient to absorb most of the energy from a 10 keV beam; however, as the voltage is increased above that level, an increasingly significant portion of energy is lost to transmitted electrons, until at 40 keV, most of the energy simply transmits through the sample because of the inverse dependence of the cross section on AV (the very slightly increased “thick” value at 40 keV is merely a statistical fluctuation). For 2 µm thickness, the transition occurs over a broader AV range, and for this MEMS device typical layer thickness, it appears that a useful range of AV’s of from ~15–25 keV would apply for single layer and two layer joints. For a thick specimen (e.g. the substrate), a high fraction is absorbed, though it is absorbed over a volume ~10 µm thick. Table 15.1 Calculated electron energy absorbed vs Si thickness and accelerating voltage Si thickness accelerating voltage
1 µm
2 µm
Thick
10 20 30 40
88% 21 7 4
90% 59 20 12
90% 90 90 92
keV keV keV keV
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Assuming a 10 µA beam, heat inputs from the conditions listed in Table 15.1 are listed in Table 15.2. It is notable that the AV for peak heat input depends upon the sample thickness. It is also notable that the values obtained in Table 15.2 are similar to that employed for the thermal calculation, suggesting that the model, while simplified, is adequate. Similarly for a Ni-based sample, energy losses due to back-scattered and transmitted electrons were calculated vs thickness and accelerating voltage and subtracted from the beam’s original energy to determine the percentage of energy absorbed. These are listed in Table 15.3. Clearly, Ni is not nearly as transparent to typical SEM accelerating voltage electrons as Si. At 40 keV substantial transmission of electrons occurs for <2 µm thick specimens; at 30 keV it occurs only for <1 µm thick specimens. Again, assuming a 10 µA beam, the corresponding heat inputs are noted in Table 15.4. The choice of Si and Ni as the materials of interest in this work was driven by work on Si MEMS and LIGA Ni microcomponents. However, they are prototypic of other materials with similar Z. Thus, Si behaves similarly to Al, Mg, and light metal alloys, while Ni behaves similarly to Fe- and Cubase materials. Table 15.2 Heat inputs for the conditions of Table 15.1 @ 10 µA beam current Si thickness accelerating voltage
1 µm
2 µm
Thick
10 20 30 40
88 mW 42 22 17
90 mW 118 60 48
90 mW 180 271 369
keV keV keV keV
Table 15.3 Calculated percentage of electron beam energy absorbed vs Ni thickness and accelerating voltage (5000 electrons) Ni thickness accelerating voltage 10 20 30 40
keV keV keV keV
1 µm
2 µm
3 µm
83% 79 53 28
83% 79 79 67
83% 79 79 81
Table 15.4 Heat inputs for the conditions of Table 15.3 @ 10 µA beam current Ni thickness accelerating voltage
1µm
2µm
3µm
10 20 30 40
83 mW 158 159 112
83 mW 158 237 268
83 mW 158 237 324
keV keV keV keV
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Micro electron beam welding parameter choice conclusions
A number of issues have been explored relative to heat transfer and energy absorption in the micro E-beam welding process. In summary, the small size of the joints possible with this technique relative to the characteristic heat transfer distance renders most thermal distributions effectively stationary rather than transient in nature, and hence the distance to the nearest effective heat sink and the material’s thermal conductivity control the needed energy input to cause melting. For relatively electron-transparent materials such as Si, it is conceptually possible to “tune” the depth of heat absorption by changing the accelerating voltage, which should aid in producing welds between layers. For materials beyond a few µm thick, the beam behaves as if a surface, rather than a volume heat source were being used. Because of the rapid heat transfer, this is only a minor consideration.
15.3.5 Advanced process variations (multi-beam and superposition) Bead-on-plate tests show the potential of fast beam control options combined with the possibility of superposition. Not only is it possible to precisely define the local energy input and thus create significantly larger pool-craters than the beam’s spot size without defocusing, but by varying frequencies and deflection patterns, the thermal conductive behavior of the material can be influenced. Figure 15.23 shows an example of this phenomenon on stainless steel sheet material (Material 1.4301 is equivalent to grade 304 stainless steel). In the five trials shown, the total energy input defined by beam power and process time is exactly the same. Two circular figures are superpositioned and their frequencies varied. Obviously, when using low frequencies, the heat can dissipate, resulting in a barely-influenced substrate. By raising the circular frequencies, the molten volume can be increased. High frequencies lead to very sharp and narrow weld seams, eventually leading to hot cracks of the material. This is a promising observation regarding the difficult weldability of materials with high thermal conductivity (like copper and aluminium) given the limitations on power of the SEM. The superposition of, for example, a step-figure and a circular figure, is shown in Fig. 15.24. Using low frequencies, the three circles are welded one after another in multiple scan mode. Raising the frequencies leads to the multi-beam effect as described earlier, resulting in three simultaneously welded seams. With even higher frequencies, all three circles appear to be simultaneously irradiated in their entirety. Analogous to the example in Fig. 15.23, the heat conductance seems to change with the frequencies used. Figures 15.25 and 15.26 show an example of the possibilities of superposition
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F1 = 1 Hz F = 100 Hz
F1 = 1000 Hz F = 1 Hz
F1 = 100 Hz F = 10 Hz
F1 = 1000 Hz F = 100 Hz
φ 38 µm (F)
φ 500 µm (F1) 20 µm
20 µm
20 µm
15.23 Superposition of two circular figures with varying frequencies.
20 µm
20 µm
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Material 1.4301 s = 30 µm PBeam = 2.7 W t = 10 s
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F1 = 0.1 Hz F = 100 Hz
F1 = 100 Hz F = 100 Hz
F1 = 1 Hz F = 100 Hz
Material 1.4301 s = 30 µm PBeam = 1.65 W t = 20 s
φ 500 µm (F1)
50 µm
50 µm
50 µm
50 µm
15.24 Superposition of a step-function and a circular figure with varying frequencies.
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Optical fiber
Carrier plate
Layer 3
250
µm
Layer 2 Layer 1
15.25 Optical fiber array.
15.26 Optical fiber in nickel silver structure (a) failed joint in line-scan mode, (b) successful joint with superpositioned beam-deflection figures.
while joining parts coated with a filler material. Here, the possible joining methods for an array of optical fibres for a plug-in connector are tested. Both the optic fiber (glass) and the nickel silver carrier plate are coated with a CuZn based filler material. While using the limited beam deflection possibilities of the SEM, it was not possible to evenly heat the joining partners, instead, due to the effects described in Section 15.2.6 (surface tension), the filler material pulled away from the joining partners and agglomerated. By superpositioning a linear and a circular function, a more evenly distributed heat input was possible, resulting in a successful joint.
15.3.6 Addition of filler material In the last decade, analytical instruments have become readily available that employ a focused ion beam (FIB) to observe, remove and deposit material. The FIB deposition technique of ion beam-assisted deposition (IBAD) is a
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form of catalyzed chemical vapor deposition which relies upon the focused ion beam to “crack” an organo-metallic gas which has been adsorbed onto a surface. The cracked adsorbed molecule leaves a metallic-like deposit behind. The gas is introduced at extremely low pressure (the instrument is very similar to an SEM) through a capillary tube aimed at the location of interest. Aiming of the capillary is non-critical, as long as the gas flow blankets the beam impingement area. The gas is generated by gentle resistive heating of an organo-metallic compound, which typically sublimes at slightly above ambient temperatures. Multiple deposit materials are available; common are Cu, Pt, and SiO2. Most of these instruments also have an electron beam column primarily used for sample observation in SEM mode, but which may be employed to deposit material as well, though the electron beam-assisted deposition (EBAD) process is less efficient than IBAD. The FIB used in the work described here produces a Pt-rich deposit from methylcyclopentadienyl trimethyl platinum. Literature surveys indicate that the FIB organometallic Pt deposit actually is less than half Pt (by atom fraction) with a similar amount of C (and appreciable Ga) content [8]. However, the deposit is fully dense and is metallic in nature (it exhibits metallic conduction). In addition to its near universal adoption as a technique for preparation of thin film transmission electron microscopy samples, various investigators have used this technique to “write” conductor patterns to repair microcircuits, and to build deposits of varying shape and complexity, in a manner very analogous with rapid prototype laser powder deposition (such as Laser Engineered Net Shaping: LENS) except on a much smaller size scale. By suitably manipulating the manner and amount of deposition, joining across gaps is also possible. Further, because the beam power is quite low (no melting is incurred, and extremely low temperatures are normally seen) thermal expansion and contraction, and microstructural changes of the base material are not an issue. Finally, since the beam is quite highly focused, deposition of material on a nanoscale is feasible. To produce a joint, the ion (or electron) beam must describe a pattern that will “grow” the deposit across a gap. FIB devices have the capability to “write” simple shapes via programmable control of the beam’s rastering. In addition to examining these patterns, beam control parameters were investigated to find optimal values for beam current, pixel dwell time and raster overlap. Non-optimal values can result in removal of material by sputtering, or rough as-deposited surfaces. In addition to developing an optimal beam deposition condition, the systematic study showed an apparent effect of fluid mechanics of the gas on the deposit morphology, with deposit thickness “ripples” emanating from corners upstream relative to the gas flow. Figure 15.27 shows the results of an examination of the deposition response as a function of the FIB operating ranges for pixel dwell and pixel spacing. For some values of parameter space, the ion beam’s effect is to remove material faster
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15.27 Behavior of deposit under varying dwell and raster overlap conditions. An apparent effect of gas flow also shown by the rippled surfaces.
than it deposits, and for other values, the deposit becomes highly non-uniform. Sputtering is not expected when using the electron beam; however, the deposition rate is much slower. Longer pixel dwell times during deposition with the ion beam require a larger spacing between pixels (negative overlap). Dwell times as long as 0.4 µs resulted in sputtering of the substrate for overlaps smaller than (less negative) –300. It appears that dwell times of 0.1 to 0.2 µs with an overlap of –200 to –400 can provide good deposits while maintaining acceptable deposition rates. One of the first practical applications of the technique was to produce “patches” over etch holes in a micromachined peristaltic pump. The etch holes are needed to produce the internal cavities of the pump body, but of course are direct leak and contamination paths. A FIB-produced patch is shown in Fig. 15.28. As noted above, building a matrix of deposits (with the standard Pt-rich molecule, methyl-cyclopentadienyl trimethyl Pt : (CH3)3(CH3C5H4)Pt) and filling gaps in MEMS die specimens showed that the deposit morphologies are dramatically affected by the direction of gas flow relative to the deposit, implying that there may be an interaction. While determining the actual pressure in the vicinity of the molecule/beam interaction region is difficult, it must be between the sublimation pressure of the molecule (~0.5T) and the chamber pressure ~1 µT. While the latter is clearly in a molecular flow regime (where the Reynolds number is meaningless), the former is in a viscous flow regime. Assuming a viscous regime, an estimate of the Reynolds number (VD/ν), where V = gas flow velocity, D = characteristic dimension and ν = kinematic viscosity, was made. An upper bound estimate of V was
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(b)
(a)
15.28 (a) Before, (b) after FIB deposit sealing of etch hole.
obtained by multiplying the molecular velocity of He at ambient (1197 m/s) by the square root of the ratio of the molecular weights of He and the FIB gas (4/319)0.5. The kinematic viscosity of the FIB gas was estimated as being equal that of a diffusion pump oil of similar molecular weight, and the characteristic dimension was set at 10 µm. The Reynolds number obtained is ~30, implying a laminar regime. Therefore, the ripple formation seen in the deposits cannot be due to turbulent flow; some other instability mechanism must be responsible. The earliest FIB deposit mechanical properties data was obtained from standard poly-Si MEMS test dies. In order to provide a “joint” a prismatic hole approximately two-thirds the width of the gauge section with approximately square top and a “stair-stepped” bottom was cut in the gauge length of a standard MEMS tensile specimen. After re-filling the hole with Pt-based FIB deposit (made at a beam current of 300 pA with a dwell of 0.1 µs per pixel and an overlap of –200), the poly-Si along the sides of the deposit was removed with the ion beam. This left a slightly narrowed single bevel FIB deposit which joined the two ends of the tensile bar as pictured in Fig. 15.29(a). This was then tested in a micro tester. Pull tests showed a maximum fracture strength only about 1/4 of that obtained from the poly-Si itself. The fracture location was through the FIB deposit, and is pictured in Fig. 15.29(b). Subsequently, a special MEMS polycrystalline-Si die was designed and fabricated using SUMMIT-V technology [9]. This die incorporated specimens for developing parameters and techniques to implement butt geometry joints, and to measure their mechanical properties. In order to assure alignment across the designed joint gaps, integral strongbacks were provided that were cut away prior to testing. An example of such a joint is shown in Fig. 15.30. Limitations in SUMMIT-V processing restricted the minimum gap width in the gage section to 2 microns. This was found excessive for efficient joining in our FIB facility, even using the optimal deposition conditions obtained above.
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(a)
15.29 (a) Pre-test single bevel geometry FIB weld specimen, (b) fracture surface.
(a)
5 µm (c)
1 µm
(b)
15.30 (a) Pre-test joint geometry of specialized tensile bar, (b) side view of half of post-test specimen (the alignment bars, ion-beam cut before testing, are seen), and (c) close-up of the fracture surface.
We therefore sought a less machine-time-intensive approach. Two approaches were attempted: in one we used the FIB ion beam to cut narrower (i.e. less than 2 µm) notches in the gauge length alignment bars and then filled them (first one side, then the other, to give a balanced geometry), and in the second, we built fillet joints across the specimen gauge length to its restraining tie bar to allow shear testing. Figures 15.31 and 15.32 show examples of these joints. Figures 15.31(a) and (b) illustrate a double-bevel geometry prepared on one of the alignment sidebars. However, it was found faster and easier to prepare a narrow gap single bevel joint geometry and
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(b)
(c)
(d)
15.31 (a) Double bevel preparation in alignment bar and (b) finished double bevel joint, (c) single bevel preparation in alignment bar, and (d) finished single bevel joint.
(a)
(b)
15.32 (a) General view (arrows show location of bevel joints shown in Fig. 15.31) and (b) close-up of fillet weld to tie bar.
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then rejoin it, as shown in Fig. 15.31(c) and (d). Figure 15.32 illustrates the fillet geometry joint between the specimen gauge length and the tie bar that keeps it restrained within a small angle of motion. Figure 15.32(a) also shows the location of the bevel joints in the alignment bars (arrowed). Frames from a movie made when microtesting one such specimen (shown in Fig. 15.33) illustrate the amount of deformation that occurs in the sidebars upon tensile loading. A COSMOS/M v2.9 FEM model of the specimen was built to analyze the bending stresses introduced by the strongback geometry. A plot of the principal tensile stress calculated is shown in Fig. 15.34. Mechanical properties obtained for these geometries are given in Table 15.5. The results obtained noted in Table 15.5 indicate that the FIB deposit joint strengths were quite variable. None of the failures exhibited any ductility. The ion beam deposits showed systematically better strengths than the electron beam deposits, despite the latter appearing to have a somewhat more uniform geometry (see Fig. 15.33(d) and (e)). The nominal poly-Si UTS measured on our die ranged from 300–720 MPa, approximately a factor of 10 greater than the FIB deposit (however, others have measured SUMMIT-V poly-Si strengths of up to 2GPa). A substantial portion of the reduction in strength values for the FIB deposits’ UTS is due to the mixed loading in the sidebars. A substantial bending component raises the actual maximum tensile stress in the sidebars relative to the loading in the main element. The FEM model determined a ratio of 5.6 for the maximum principal stress in the sidebar to that in the main gauge length away from the sidebars. Additionally, the roots of the single bevel joints are not well formed (see Fig. 15.31), introducing an additional stress riser. In the fillet geometry joints, where a more regular geometry was always obtained, the strength of the deposit was sufficient to crack or fracture the tie bar when the deposit pulled off. The range of nominal UTS values for this geometry (always done with the ion beam) ranged from 150–240 MPa, and in several cases the fracture path deviated from the FIB deposit into the poly-Si. Also, the simple load/area calculation is clearly violated in this triangular shaped deposit, and a substantial stress riser is also seen near the root of the fillet in the FEM calculation. While the single bevel results are rather poor, the fillet results seem to suggest that the strength of the FIB deposit can approach that of the poly-Si under favorable conditions. Further development of procedures to assure better deposit geometry would be needed, however.
15.4
Examples of joining and applications
15.4.1 Metallic Knowledge about the equipment gained in its development was then applied for the study of wire joints and metal foils in a variety of materials. Welding
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(b)
(c)
(d)
(e)
15.33 (a-c) Frames from video of FIB joint-containing microtensile test, (a) essentially unloaded, (b) max load before breaking, (c) after fracture (top sidebar has broken at joint) (d) fractography of (different) test specimen showing good reproducibility of joints (ion beam deposit) (e) fractography of FEB deposit also showing excellent reproducibility.
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High bending stress
15.34 Principal tensile stress calculated in test bar (units of psi). Note indications of elevated tensile stress due to bending at the single bevel joint location, at the juncture of the sidebar with the main gage length, and at the fillet geometry joint location in the tie bar, (see Fig. 15.33(a)–(c) for an actual specimen showing fractures at both high stress locations in the sidebar).
Table 15.5 Single bevel joint mechanical strength ID
Poly layer
Gage width (um)
Gage thickness (um)
Net failure load (mN)
UTS (MPa)
I2 I3 I4 I5 I6 E1 E2
1–2 1–2 1–2 1–2 1–2 1–2 1–2
4.14 4.14 4.14 4.14 4.14 4.14 4.14
2.5 2.5 2.5 2.5 2.5 2.5 2.5
0.47 0.51 0.4 1.34 1.24 0.27 0.3
45.4 49.3 38.6 129.5 119.8 26.1 29.0
I# Ion beam; E# Electron beam deposited
feasibility tests (bead-on-plate melt runs) were carried out for the various metals as listed in Table 15.6. Tests showed that the maximum beam power of the modified SEM is not sufficient to melt copper. No alteration of the surface could be observed. Obviously, the high thermal conductivity of copper dissipates the applied heat too rapidly. Aluminum also has a high thermal conductivity. In addition, aluminum reacts with the surrounding oxygen (even in the evacuated work chamber) and generates a layer of refractory oxide (melting temperature = 2030°C),
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Table 15.6 Tested metallic materials and their properties Material
Atomic number(s)
Aluminium 13 Copper 29 Nickel 28 Nickel/ 28/24 chromium 80%/20%
Melting temperature TE [°C]
Boiling temperature TS [°C]
Thermal conductivity λ [Wm–1K–1]
Thickness s [µm]
660 1083 1453 1400
2467 2567 2732 –
237 401 90.9 13.4
50 25 50 25
which has to be cracked by the beam in order to reach the material. On the other hand, aluminum generates fewer back-scattered electrons when struck by the E-beam; therefore, a higher effective beam power is available. All told, aluminum also proves to be a material which is difficult to weld, even if the results shown in Fig. 15.35 show molten areas. However, neither with high power parameters (left) nor with high intensity parameters (right) was complete root fusion possible. The use of the multi-beam technique in combination with a high intensity beam seems promising, as the oxide layer can be cracked and the material melted locally. The bead shape found is typical for aluminum. Nickel and a Ni/Cr alloy’s weldability were also tested. Condenser coil currents of both 290 mA and 260 mA proved the weldability of nickel with the modified SEM (Fig. 15.36). Measurements show that lowering the condenser coil current from 290 mA to 260 mA leads to a slightly lower effective beam-power, but also to a smaller effective spot size accompanied with a higher power-density. In all cases studied, a continuous welding bead, 20 to 25 µm wide, can be realized over the full length. The base material is visibly deformed in the rim of the weld, due to the heat impact on the grain boundaries. The surface of the solidified weld is very smooth. The heat influence on the specimen and its surroundings is very high, due to the long process time of 60 seconds. Alloying the material can improve the weldability in terms of decreasing the necessary amount of energy. In this case, chromium is a suitable alloying element. Furthermore, chromium increases the resistance to heat and corrosion. For further tests, NiCr sheet material with a Cr content of 20% was chosen. The results of alloying Ni with Cr are shown in Fig. 15.37. Despite the much lower energy used, the volume of molten material is visibly higher compared to the results with pure nickel. The effective beam power is reduced by more than 50%. Even a condenser coil current of 230 mA, which leads to a lower beam power-density compared to 260 mA, and a process time of only 20 seconds result in a 35 µm to 45 µm wide welding bead. This may be
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Welding mode: multiple scan ts = 30 s
Welding mode: multiple scan ts = 60 s
IK1 = 290 mA IEm = 160 µA PSt = 2.32 W
IK1 = 260 mA IEm = 160 µA PSt = 2.00 W
463
15.35 Weld test: aluminum 50 µm.
caused by the lower heat-conductivity as well as the smaller thickness of the sheet material. Figure 15.38 shows examples of joining. Joining of thermocouple elements made of NiCr/Ni wire combinations with a wire diameter of 70 µm each allows almost globular beads for the temperature measurement in the micro range. On the right the joining of Ck101 sheet material (equivalent to SAE1095 plain carbon steel) is shown. The following welds were made with several SEMs (ISI 90, LEO 1430 VP, JEOL 6400) with only the final aperture removed and the condenser lens’ current adjusted to provide maximum beam current to the specimen. In Fig. 15.39 a successful cross-wire weld between two wires of Tophet C (high resistivity bridgewire alloy) is shown. Critical to the success of the joint was the intimate contact of the two wires; without the contact, the individual wires were cut. For the geometry portrayed in Fig. 15.40, it was observed that the protruding pin did not melt until after a few seconds beam irradiation, but then melt back occurred rapidly. An explanation for the delayed melting suggested is that it takes some time to heat the whole assembly to near the melting range, and that once this ‘preheating’ phase has occurred, only the latent heat of melting needs to be supplied. In Fig. 15.40, the before welding picture was taken without the final aperture in the column, and shows the image quality which may be expected in this condition. Figure 15.41 shows a LIGA-produced Ni gear melted onto alloy steel shaft. This demonstrates the major effect of surface tension when parts are
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200 µm
200 µm
20 µm
10 µm
Welding mode: multiple scan ts = 60 s
Welding mode: multiple scan ts = 60 s
IK1 = 290 mA
IK1 = 260 mA
IEm = 160 µA
IEm = 160 µA
PSt = 2.19 W
PSt = 1.88 W
15.36 Weld test: nickel 50 µm.
not sufficiently restrained, similar to ‘tombstoning’ of soldered surface mounted devices. Figure 15.42 shows that it is possible to finely control the heat input. This picture shows a joint geometry intended to attach a rotating part to a clock plate. On the left the tool steel gauge pin is welded to the LIGA Ni disk, while on the right it is balled up to act as a rivet, allowing rotation. In Fig. 15.43 extensive annealing of the normally fine-grain LIGA Ni material has occurred (before weld hardness of >300 VHN), and indications of grain boundary separation and voiding are present (not atypical in electroplated materials which may have had “brighteners” like saccharine added to the plating bath). It is also noted that the steel pin (darker etching material) has melted and “brazed” to the Ni, which has apparently not melted. Figure 15.44 shows melting of a poly-Si MEMS test specimen, which did
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20 µm
20 µm
Welding mode: multiple scan ts = 20 s
Welding mode: multiple scan ts = 5 s
IK1 = 230 mA IEm = 60 µA PSt = 0.83 W
IK1 = 260 mA IEm = 60 µA PSt = 0.98 W
15.37 Weld test: nickel/chromium 25 µm.
Type K thermocouple, (=70 µm, PSt = 1.25W crossed configuration
Ck101, s = 2 × 50 µm, PSt = 1.65 W fillet weld at the lap joint
15.38 Examples of micro E-beam welded micro components.
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15.39 Cross-wire weld of 30 µm diameter Tophet C, performed on ISI Model 90A SEM. Inset shows pre-weld geometry, with wires in intimate contact before melting.
15.40 Welding of LIGA (nearly pure Ni) gear to tool steel gage pin. Figure on left before welding is taken without final aperture in column, and shows the image quality which may be expected in this condition (LEO 1430VP).
not bridge the gap. Figure 15.45 shows melted Si fibers of ~1µm width which did successfully coalesce. Figure 15.46 shows the major effect of a heat sink. The left side of the MEMS Si compact tension fracture toughness-like specimen was “floating” while the right side is anchored to the substrate. The sample was melted by reducing the rastered area from the full field to a small rectangle centered on
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15.41 LIGA-produced Ni gear melted onto alloy steel shaft, showing strong effect of surface tension.
15.42 Pin/rivet geometry intended to attach a rotating part to a clock plate. On the left the tool steel gauge pin is welded to the LIGA (Ni) disk while on the right it is balled up to act as a rivet (LEO 1430VP).
100 µm
50 µm
15.43 Metallographic section of the gear to shaft ‘weld’ of Fig. 15.42. Post-weld hardness of LIGA Ni was 122–128 VHN.
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15.44 Melting of poly-Si MEMS test specimen, which did not bridge gap (JEOL 6400). Post-weld hardness of LIGA Ni was 122–128 VHN.
10 µm
15.45 Melted Si fibers of ~1 µm width which did successfully coalesce (JEOL 6400).
the “notch”. The diagonal cut from the upper left to the center was caused by the progression of the interscan dwell location as the scan area was reduced.
15.4.2 Non-metallic In the field of plastics, the two semi-crystalline thermoplastic resins polyethylene (PE) and polyacetal (POM) as well as the amorphous polymethylmethacrylate (PMMA) were investigated (Table 15.7).
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15.46 Shows major effect of heat sink. Left side of MEMS Si CT specimen is ‘floating’ while right side is anchored to substrate. Table 15.7 Tested polymeric materials and their properties Material
Structure
Melting temperature TE [°C]
Boiling temperature TS [°C]
Thermal conductivity λ [Wm–1K–1]
Thickness s [µm]
PE
Semicrystalline Amorphous Semicrystalline
220–280
–
0.29–0.51
100
220–250 205–215
– –
0.18–0.19 0.29–0.40
500 500
PMMA POM
Plastics are usually not joined by electron beam welding since the lack of electrical conductivity prevents the conduction of electrons. This leads to charging of the material, vaporization and uncontrollable discharge, resulting in a flashover of the beam. Weld feasibility tests with gold plated plastic components were carried out. The gold plating has a thickness of only a few µm. The results are shown in Fig. 15.47. Only PE shows reasonable results, though with a scaly structure to the weld bead. The rest of the gold plating shows no signs of heat-induced
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PMMA
200 µm
PSt = 1.9 W vs = 3 mm/min
POM
100 µm
PSt = 1.6 W
15.47 Weld tests on plastics.
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cracks. The structure of the weld leads to the assumption of a foam-like consistency. The resin expands explosively due to the massive heat input, leading to the volatilization of contained solvents. With longer process times, the resulting contamination of the vacuum led to the shutoff of the SEM. Obviously, the chamber pumping capacity would need to be increased to use this process for PE. With PMMA, there is little reaction visible. Only the gold plating is sublimated and there may be slight alteration of the upper layers.
15.5
Summary
Micro E-beam welding is in its infancy [10–13]. The work reported here however shows that a great deal of promise is exhibited by this technology, even using simply converted existing SEMs. We have reported a few steps taken toward the understanding of the processes in the areas of beam characterization, beam-material interactions, thermal control, fixturing and the integration of electron-optical imaging into a microwelding processes. Much of this work has been in spot welds, but great strides are being made in the incorporation of real welding controls (i.e. motion and beam controls) which will facilitate research into welding of seam geometries, which will greatly extend the application of the technique. A small number of materials have been investigated, and certainly there is much work to do here. Issues of microstructure and the resulting mechanical properties of the relatively large grain size fusion zones which result have barely been touched. Finally, we introduced in FIB joining a technique that may be capable of reducing the size scale of joining to below the micron size scale.
15.6
Acknowledgements
Sandia is a multiprogram laboratory operated by Sandia Corporation, a Lockheed Martin Company, for the United States Department of Energy’s National Nuclear Security Administration under contract DE-AC0494AL85000. We would also like to thank Dr J.R. Michael and Mr D.O. MacCallum, both of Sandia National Laboratories for reviewing the text. We would also like to thank the German Society for Research (DFG) for financing the Collaborative Research Centre 440 “Assembly of Hybrid Microsystems”, which funded the project of electron beam microwelding investigations at Aachen University.
15.7
References
1. Schultz, H., Elektronenstrahlschweißen, Deutscher Verlag für Schweißtechnik (DVS) GmbH, Düsseldorf, 2000.
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2. Schiller, S. et al., Elektronenstrahltechnologie, Wissenschaftliche Verlagsgesellschaft mbH, Stuttgart, 1977. 3. Janssen, W., Verbesserung des Elektronenstrahlschweißens mit Hilfe der flexiblen Doppelfokussierung. Dissertation RWTH Aachen, VDI Verlag, Reihe 2: Fertigungstechnik, Nr. 230, 1991. 4. Goldstein, J.I., Newbury, D.E., Joy, D.C., Lyman, C.E., Echlin, P., Lifshin, E., Sawyer, L., Michael J.R., Scanning Electron Microscopy and X-Ray Microanalysis, Kluwer Academic/Plenum Publishers, New York, 2003, 61–98. 5. Heinrich, K. F. J., Newbury D. E., Yakowitz, H., (eds), “Use of Monte Carlo Calculations in Electron Probe Microanalysis and Scanning Electron Microscopy”, National Bureau of Standards Special Publication 460, 1976. 6. Gauvin, R., Drouin, D., Couture, A.R., Casino v2.42, available at www.gel.usherb.ca/ casino/ 7. Joy, D.C. Monte Carlo Modeling for Electron Microscopy and Microanalysis, Oxford University Press, New York, 1995, p. 82. 8. Tao, T., Wilkinson, W., Melngailis, J., J. Vac. Sci. Technol. B9, 162, (1991). 9. SUMMIT-V: for information visit http://mems.sandia.gov/tech-info/summit-v.html. 10. Dilthey, U., Brandenburg, A., Moller, M., Smolka, G., “Joining of Miniature Components”, Welding and Cutting, 2005, 52 No. 7, E143–E148. 11. Hwang, I.-L., Na, S.-J., “A Study on Heat Source Modeling of Scanning Electron Microscopy Modified for Material Processing”, Metallurgical and Materials Trans. B, 2005, 36B, 133–139. 12. Knorovsky, G.A., Nowak-Neeley, B.M., Holm, E.A., “Microjoining with a Scanning Electron Microscope”, Science and Technology of Welding and Joining, Vol. 11, No. 6, 2006 (electronic journal). 13. Knorovsky, G.A., MacCallum, D.O., Meyers, M.T., “Selection of Parameters for µE-Beam Welding”, Science and Technology of Welding and Joining, Vol. 11, No 6, 2006 (electronic journal).
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16 Resistance microwelding S F U K U M O T O , University of Hyogo, Japan, Y Z H O U , University of Waterloo, Canada and W TAN, Medtronic Inc., USA
16.1
Introduction
Resistance microwelding (also termed small-scale or micro-resistance welding) is a group of microjoining processes in which the coalescence of metals is produced at the faying interfaces, similar to large scale resistance welding, by the heat generated by the resistance of the workpieces to the passage of electric current [1]. Resistance microwelding is increasingly used in the fabrication of electronic components and devices (such as batteries, cellular phones, interconnections in printed circuit boards, relays, sensors, air-bag diffuser screens, and medical devices [2–7]). In this chapter, we will first introduce the principles of resistance microwelding and their differences compared to regular or large-scale resistance welding. We will then discuss the process variations of resistance microwelding and their bonding mechanisms, especially in terms of metallurgical changes. We will then cover process conditions and their influences on welding process and joint quality. Several process parameters, including the key variables of welding current, electrode force, and weld time, and surface conditions of materials need to be closely controlled to obtain sound and consistent joints. Finally we will discuss dynamic resistance and its application in process monitoring and control, and effects of power supplies and electrodes on welding results.
16.2
Fundamentals
To better understand the principles of resistance microwelding, resistance spot microwelding through melting will be used as an example. Suppose two metal sheets are to be welded (Fig. 16.1). First, two electrodes are pressed against the sheets squeezing them together. Electric current is then passed from electrode to electrode through the sheets, causing heating and eventually localized melting and coalescence of a small volume of the sheet materials, due to the resistance heating caused by the passage of electric current. The volume of material melted and subsequently resolidified depends 473 WPNL2204
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Electrode
Molten nugget
Solidified nugget
Workpiece
I Electrode
16.1 Schematic of resistance spot microwelding through melting and resolidification [11].
on the amount of heat generated and directly affects weld strength. The heat delivered to the assembly of sheets and electrodes can be expressed as
Q = I2 R t
(1)
where, Q is the heat generation, I is the welding current, R is the resistance of workpieces, and t is the duration of the current (weld time). The resistance, R, includes contact resistance at the electrode/workpiece interface and at the faying interface between the two workpieces, and bulk resistances of workpieces and electrodes. For the two-sheet joint, these are the resistances at: • • • • • • •
upper electrode (R1) contact between upper electrode and upper sheet (R2) body of upper sheet (R3) contact between upper and lower sheets (faying interfaces) (R4) body of lower sheet (R5) contact between lower sheet and lower electrode (R6) and lower electrode (R7) (Fig. 16.2) [1].
Each of these resistances changes dynamically during the process and their relative magnitudes control the process. Among them, the contact resistance at the faying surfaces, which is influenced by surface characteristics (such as cleanliness, roughness, hardness and presence of coatings or plating materials) and electrode force (or pressure), are critical to the process, especially in the early stages. Usually, the electrodes are made of high thermally and electrically conductive copper alloy, so that electrode bulk resistance should be the lowest among all components. Preferably, the highest resistance component should be at the faying surface between the sheets to be joined. If heat generation is suitably concentrated at the faying surface region, it will lead to coalescence and bonding by transient melting in this region, leaving a fusion nugget (Fig. 16.3). Ideally, a sound nugget of greater than a specified
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Load
R1
R2 R3 R5
R4 R6
R7
16.2 Resistances involved in resistance spot microwelding.
HAZ
Indentation
Surface flash
Weld nugget Sheet separation
Expulsion
Porosity
Corona bond
Crack
16.3 Illustration of cross-sectional view of resistance microwelding.
minimum size is needed, centered on the electrode axis and the sheet thicknesses, and without any defects, surface splash, electrode sticking or weld metal expulsion. Generally, the process parameter tolerance range must be established prior to production welding. The range of welding current and time combinations over which acceptable nugget characteristics may be obtained for a given equipment and material combination is generally referred to as the weld lobe [4]. While resistance welds normally involve localized melting of the base metals, it should also be noted that the process may be used to make solid-state bonded or brazed/soldered resistance microwelds [2–4, 8–10], e.g., between Au-plated nickel sheets [11, 12]. In contrast to the extensive work on regular or large-scale resistance
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welding as used in automotive and appliance industries for parts with a sheet thickness or wire diameter larger than 0.5–1.0 mm, only limited systematic investigations have been carried out in resistance microwelding. For example, recommendations for process parameter selection have been provided for resistance welding of steels and, to a lesser extent, of some other alloys [1]; however, very limited information is available for resistance microwelding. On the other hand, resistance microwelding is not simply a result of downsizing from regular resistance welding (Table 16.1) [10, 11]. There are many differences between these two groups of processes. First of all, metals to be welded in resistance microwelding are mostly non-ferrous such as copper [13, 14], Kovar [15, 16], nickel [12, 15, 16], platinum [17], brass [13, 14], aluminum [13, 14], nickel-copper alloys [16], titanium, silver, etc., while the workpieces in regular resistance welding are mainly steels. Moreover, sheets and wires for electrical components in resistance microwelding are often coated with other metals such as Au, Ag, Ni, Sn, etc., whereas materials joined by large scale resistance welding are usually uncoated or coated with Zn. Workpieces’ (base metals and coatings) physical properties influence their weldability. Electrical resistivity and thermal conductivity should be the most important physical properties in terms of Eq. (1). But other physical properties such as melting point, latent heat of fusion, and specific heat are important as well. For example, one study indicates that the weldability, evaluated as current magnitude for a given nugget diameter, of Al, brass and Cu with a.c. welding equipment can be listed in a decreasing order of Al, brass, and Cu, which is not exactly in the same order of their electrical resistivity or thermal conductivity (i.e., brass < Al < Cu) [14]. In this case, a very simplified model that neglects contact and electrode resistance can account for the weldability by considering melting temperature, specific heat, and latent heat of fusion. Table 16.1 Differences between small- and large-scale resistance spot welding
Sheet thickness Electrode force Welding current Electrode cooling Metals to be welded Plating materials Applications
Small-scale resistance welding
Large-scale resistance welding
< 0.2–0.4 mm < 100–200 N < 2,000–4,000 A No Mainly non-ferrous metals (Kovar, Ni, Ti, etc.) Ag, Au, Ni, Pb, etc. Electronic and medical components
> 0.6–0.8 mm > 1,000–2,000 N > 6,000–10,000 A Yes Mainly steels and Al-alloys
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The second difference between resistance microwelding and larger scale welding is due to the different magnitude of electrode forces used. Specifically, the electrode pressure is much lower in resistance microwelding. Because of this difference, in regular resistance welding, the maximum nugget diameter can be almost the same as the electrode tip diameter. Therefore, when weld metal expulsion occurs, the electrode tip will indent into the workpiece causing a decrease of joint strength. In resistance microwelding, the ratio of maximum nugget diameter (Dn) to electrode tip diameter (De) is about 13 to less than one depending on electrode pressure (Fig. 16.4), and therefore the indentation of electrode tip into workpiece is less likely to occur. Therefore, weld metal expulsion has less influence on joint strength in resistance microwelding [18]. Third, the risk of electrode sticking during resistance microwelding is much greater than with resistance welding. In resistance microwelding, electrodes are not internally water cooled due to their small size, leading to higher electrode tip temperatures and a risk of incipient welding of electrodes to worksheet surfaces. In addition, the much lower electrode pressures result in much higher contact resistance, which reduces welding current thresholds during resistance microwelding but also promotes electrode sticking.
16.3
Process variations
Some of the process variations such as resistance spot, crossed wire, seam, flash, upset, and projection welding, etc., [1] in regular resistance welding are also found in resistance microwelding.
Large scale
Small scale
16.4 Comparison of nugget size between large-scale and small-scale resistance spot welding [18].
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16.3.1 Sheet to sheet Basically, sheet-to-sheet joints can be produced by direct (Fig. 16.1) or series-mode (Fig. 16.5) resistance spot microwelding processes. For example, series-mode resistance spot microwelding is applied to fabrication of cellphone battery packages (Fig. 16.6) [19]. Since weld nuggets tend to be formed predominantly in the upper sheet due to the effect of current shunting in the series-mode resistance welding process, engineered adjustments should be made to adjust heat balance for instance by changing sheet thicknesses, electrode positions, or electrode materials, during resistance spot microwelding. Fusion bonding is observed in resistance spot microwelding of most bare metals [11]. At the initial stage, partial surface melting first occurs in an annular ring-like area, which is a result of the current density at this region being much higher than that of the central region, due to current constriction [18]. Then surface melting spreads towards the center of the faying surface. Subsequently, surface melting covers the whole area. As the temperature
Load
Workpiece
Load
Nugget
Nugget
16.5 Schematic of series-mode resistance spot microwelding.
16.6 Electric contacts in battery for cellular phone. Note resistance spot microwelds; black is as welded, white is after peel testing.
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builds up, the fusion nugget starts to grow from the center of the faying surface (Fig. 16.7). If too much energy was input to the weld, weld metal expulsion will occur. In fact, surface melting is preceded by oxide film breakdown, which will be described later on.
16.3.2 Crossed wire In resistance cross-wire microwelding two wires are joined to each other, usually at a right angle [20]. In large-scale welding, cross-wire welded products (mainly steel) include such items as stove and refrigerator racks, lamp shades, baskets, fencing, concrete reinforcing, etc. On the other hand, cross-wire microwelding is commonly used in electronics, and medical and instrument components, mainly for electrical interconnections [21, 22]. This set-up of resistance microwelding (Fig. 16.8) is very similar to projection welding because two crossing wires form a point contact at the faying surfaces [23]. Crossed-wire and projection welding differ significantly from flat sheet welding because of the progressive set-down/collapse of the joint assembly during the welding operation and the resultant increase of contact area. For fine nickel wires, it has been suggested that resistance microwelding includes the following stages: • • • •
cold-wire collapse surface melting molten-phase squeeze-out and solid-state bonding (Fig. 16.9) [24].
The joint quality in cross-wire microwelding can be evaluated using tensile shear tests (Fig. 16.10). The value of ‘set-down’ (embedment) (Fig. 16.11) has also been suggested as an indicator of joint quality [16, 25]. It has been proposed that the measurement of set-down could offer a non-destructive method of monitoring joint quality. But the work on fine nickel wires suggested that the set-down alone was not a good quality indicator. Figure 16.12 shows the cross sections of joints with different welding conditions but similar setPartial surface melting
Fusion nugget
16.7 Schematic of nugget formation during resistance microwelding of bare nickel sheets [11].
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(a)
(b)
16.8 Setup of cross-wire microwelding: (a) illustration, (b) Nitinol cross-wire joint.
down values. Joint B showed much higher joint strength than Joint A despite having the same set-down. A recrystallized microstructure was associated with a strong bond interface (Fig. 16.12(b)), while cold-pressed microstructure with little recrystallization was associated with a weak bond interface (Fig. 16.12(a)). Therefore, not only a sufficient set-down, but also an appropriate local temperature history associated with this set-down are important to produce a strong bond interface.
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Stage 1
Liquid phase
Stage 2
Squeezedout liquid
Recrystallization Stage 3
Flash
Stage 4
16.9 Main stages in resistance microwelding of crossed fine nickel wires [24].
16.3.3 Wire to sheet Fine Al, Au, or Cu wires are also joined to other electronic components such as housings, printed circuit boards, integrated circuits, power devices, cables and so on. For example, wires are connected to printed circuit boards using parallel-gap microwelding, which is a variation of resistance microwelding
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16.10 Tensile shear test for cross-wire joint.
B
A
A
Set-down = (A – B)/A × 100 (%)
16.11 Definition of set-down (embedment).
(Fig. 16.13). This is very similar to series-mode resistance spot microwelding but two nuggets are formed in series-mode resistance microwelding (Fig. 16.5) and only a single nugget is formed in parallel-gap microwelding. The process may be solid state (e.g. Pt wire to Pt pad [26]) or with a transient liquid phase similar to cross-wire microwelding. In practice, both substrate and wire are often plated with other materials such as Au, Ni, Pb-Sn alloy or Ag. Those plating metals often have a lower melting point or can react with the base metals to form a braze alloy. The braze fillet can strengthen the joint. The effect of plating on both bonding mechanisms and joint quality is one of the significant factors in resistance microwelding, which will be further described in Section 16.4.2.
16.4
Process conditions: welding parameters
16.4.1 Welding current Welding current is the most significant variable in resistance microwelding; having a power of two (I2) it influences the rate of heat generation, as shown
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100 µm (a)
100 µm (b)
16.12 Typical cross sections of cold or hot collapse. Crossed-wire joints with similar set-down but different joint strength [24]. (a) Joint A: Cold collapse (set-down of 78%, joint breaking force of 2 N), (b) Joint B: hot collapse (set-down of 84%, joint breaking force of 42 N).
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Power supply (constant voltage)
Electrode force Electrode
Wire Pattern Substrate Weld
16.13 Schematic of parallel-gap microwelding for wire to board.
in Eq. (1). Therefore, current is the variable that needs to be controlled most carefully. Various types of welding current waveform may be utilized, depending on the type of available power supply, including line-frequency or high-frequency alternating current (a.c.), direct current (d.c.) from three-phase rectified supply, capacitor discharge (CD), or d.c. from rectified high frequency inverter supply. Usually the root mean square (RMS) values of the welding current are used as machine settings and also for process control. In all cases, it should be noted that the size of weld nugget increases with increasing welding current for a given weld time regardless of the type of power source in resistance microwelding. Since the joint strength defined by joint breaking force is usually roughly proportional to nugget size, joint strength tends to increase with welding current as well (Fig. 16.14 [15]). Below a critical welding current or current threshold, no weld nugget is formed. Above the threshold, the nugget grows rapidly with increasing welding current, but often decreases in volume after reaching the peak value due to weld metal (liquid) expulsion. Excessive energy input may also reduce weld strength due to severe softening of the HAZ. Therefore welding current needs to be optimized to ensure the formation of a sound nugget without any defects, weld metal expulsion, electrode sticking, severely softened HAZ or other quality problems.
16.4.2 Weld time According to Eq. (1), the total heat generation should be proportional to the weld time. However, heat losses due to conduction into the electrodes,
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160 140
22 N 44 N 66 N
Peel force (N)
120 100 80
22 N
60 Electrode force = 66 N 40 44 N 20 0 0.2
0.4
0.6
0.8 1.0 Welding current (kA)
1.2
1.4
1.6
1.2
1.4
1.6
(a) 1.6
Nugget diameter (mm)
1.4
22 N 44 N 66 N
1.2 1.0 0.8
22 N
0.6 0.4 0.2 0.0 0.2
Electrode force = 66 N
44 N
0.4
0.6
0.8 1.0 Welding current (kA) (b)
16.14 (a) Peel force (b) nugget diameter versus welding current at different electrode forces for the 0.2 mm Kovar joints using the a.c. power supply and 8-cycle weld time [15].
surrounding workpieces and air must be considered. Although these losses increase with increasing weld time, they are essentially difficult to control [1]. Increasing the weld time could decrease the current threshold to form a weld nugget [11]. However, heat generation is not actually directly proportional
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to the weld time because both contact resistance at the faying surfaces and bulk resistance of workpieces change drastically during the welding process (Fig. 16.15) [11, 13, 27]. When the weld time is very short (such as when using a CD power supply), more of the heat is generated at the faying surfaces because at the beginning of the process the contact resistances are high and bulk resistances are low. Although electrical resistivity of each material increases with increasing weld time, due to increasing temperature, the effect of time on heat generation is less significant than that of welding current. Therefore a long weld time is not an effective way to generate local heat. If the welding current is too low, simply increasing the weld time may not make a weld. Conversely, a very short weld time may also lead to weld quality problems, due to the high welding current required.
16.4.3 Electrode force The electrode force is a very important variable in resistance microwelding and particularly in cross-wire microwelding. The role of electrode force will be described separately for resistance spot and cross-wire microwelding. Sheet to sheet resistance spot welding The electrode force influences the resistance microwelding process mainly through its effects on the contact resistances and on the contact areas. As described above, the influences of electrode force in resistance microwelding are more significant than in large-scale resistance welding. The contact resistances decrease as electrode force increases [28]. Increasing electrode force will increase the contact radius at the contact interfaces, decreasing the welding current density, and hence delaying nugget initiation and growth for a given current value. A longer threshold weld time is also needed for nugget
Bulk
Resistance
Contact
Weld time
16.15 Schematic showing the change in resistance during resistance microwelding [13].
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initiation when the electrode force is increased. Therefore, increasing electrode force often needs to be accompanied by increased welding current and/or weld time to compensate for the reduced resistance. Increasing electrode force also decreases the cooling rate at the nugget center after the welding current is turned off [29]. If electrode force is too low, weld metal expulsion may occur immediately after starting the welding current due to excessive faying surface contact resistance, resulting in very rapid local heat generation. Crossed-wire welding Electrode force must be balanced in accordance with welding current especially in cross-wire microwelding. As described in Section 16.3.2, the balance of deformation and heat generation must be controlled to form an adequate microstructure, leading to high joint strength. An optimum electrode force may be determined for each material and geometry, as shown in Fig. 16.16 [24]. Usually joint strength in cross-wire microwelding is evaluated by joint breaking force, which is the product of bonded area and interfacial strength. An increase in electrode force, if it did not affect heat generation too much, would increase the bonded area (through hot collapse), resulting in an increase of joint breaking force. However, a further increase in electrode force could reduce resistance and heat generation too much to achieve high interfacial strength, resulting in a low joint breaking force, even though a relatively large wire deformation is achieved (cold collapse). Therefore, optimized process parameters need to produce sufficient surface melting at the very beginning of the welding sequence. This molten metal, carrying the surface contamination, needs to be squeezed out to produce ‘clean’ surfaces for a strong bond. A relatively large bonded area is also required, since the load100
90
4 80 3 70 2
Set down [O]; (%)
Joint strength (kg)
5
0.8
0.7
0.6
0.5 60
1
tw = 6 cycle
0 0
5 Welding force: Fw (kg)
50
Static resistance [+]; (mΩ)
6
0.4
10
16.16 Effect of welding electrode force (Fw) on joint strength [24].
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carrying capacity of the joint is determined by both interfacial strength and bonded area. Either a high welding current or low welding force would seem to be effective ways to produce sufficient molten phase. However, excessive welding current would cause electrode sticking and degrade the wire properties in the heat affected zone, and insufficient welding force cannot produce a sufficiently large bonded area. On the other hand, although a large welding force can increase the bonded area, it would also reduce the initial contact resistance and, hence, heat generation, which produces insufficient surface melting. Thus the ideal situation would be to use a lower electrode force at the beginning of welding to obtain sufficient heat generation and, subsequently, to increase the force to obtain a large bonded area. If the electrode force increases to its final value over a period of time, it is preferable to initiate the welding current when a ‘firing force’ is reached so that the initial part of the weld cycle takes place while force is still increasing toward the welding force, if the welding system is able to do that. A relatively low firing force, that is, initiating the welding current well before the electrode force increases to its set point, will result in high initial heat generation to produce sufficient surface melting, and the larger welding force will then work as a forging force to enlarge the bonded area while squeezing molten metal out. This method could improve both the interfacial strength and bonded area. This differs significantly from the general industry practice where firing force (point of current initiation) is usually set at 90 to 97% of the full welding force [24].
16.5
Process conditions: surface
Every resistance microwelding process makes active use of electrical resistances, i.e., contact resistance and bulk resistance as shown in Eq. (1). Surface conditions and topography at the faying surfaces as well as at electrode to sheet interfaces will influence the contact resistances. Contact resistances are normally considered to comprise two separate components: i.e., constriction and film resistances. In considering the effects of surface conditions on contact resistance, the following factors need to be considered.
16.5.1 Surface roughness Even if two metallic bodies having very flat surfaces are pressed together under a certain load, the actual atomic contact occurs only at a number of local contact points microscopically. As the force increases, contacting regions are compressed elastically and then plastically. In either case, the contact resistance could be expressed in terms of the force (F), and the materials’ properties such as modulus of elasticity (E), contact or penetration hardness (H) , and electrical resistivity (ρ): as shown in Eq. (2) and Eq. (3) [30, 31]:
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Resistance microwelding
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1
Rc =
0.57ρ E 3 n FD
: Elastic case
(2)
: Plastic case
(3)
1
ξH 2 Rc = 0.89 ρ nF
where n is number of contact spots, and ξ is the pressure factor. Either a larger force or a larger number of contact spots will lead to lower contact resistance according to those equations. For example, static resistance is normally found to decrease as force increases in resistance welding [28] and in cross-wire microwelding as shown in Fig. 16.16 [24]. Furthermore, for the same material, a change in surface roughness leads to different contact resistance. The surface roughness needs to be controlled more carefully in resistance microwelding than in large-scale resistance welding since relative surface roughness compared to workpiece dimensions would be larger. As a general rule, the peak-to-peak surface finish variations for small- and microscale parts should not exceed 10% of the part diameter or thickness (Fig. 16.17) [10].
16.5.2 Oxide film Most metals have some non-metallic reaction film, usually oxide, on the surface, and may in addition have extraneous surface contaminants. Materials forming the surface films are typically non-conductors or at least have much lower electrical conductivity than the underlying metals. The film resistance RF can be written as [30, 31], RF =
ρt ξ H F
(4)
where ρt, the film resistance per unit area is called tunnel resistivity [30]. In order to pass current through workpieces to generate heat and make a nugget at the faying interface, film breakdown is necessary during the process. Diameter (d )
Surface finish variations < 10% of d
Top part
Bottom part
16.17 Acceptable surface finish variations in relation to the part thickness [10].
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Two methods are recognized for the generation of conducting pathways through oxide films, i.e., electrical breakdown and mechanical breakdown. In the case of resistance microwelding, mechanical film breakdown would not be expected to be of great importance due to the relatively low electrode force. So the electrical breakdown mechanism probably best fits the actual physical situation as it occurs in the early stage of the welding sequence [32]. In most cases the existence of surface oxide films and contamination is undesirable because electrical resistance increases with increasing film thickness. This may in turn generate larger heat at electrode/sheet interfaces, hence leading to electrode sticking. In practice, if materials have contamination or thick oxide films, they should be cleaned by mechanical or chemical operations. Wire brush abrading is the most popular method to remove oxide film. However, it should be gentle enough to prevent formation of an excessively rough or scratched surface because large surface roughness will increase contact resistance as described above. Chemical cleaning methods are widely used in large-scale resistance welding (LSRW). Since the most suitable etchant or cleaning agent varies with material composition and prior history, the selected cleaning process and chemistry needs to be specifically developed and tested for each application.
16.5.3 Plating Surface coatings are applied to improve corrosion resistance or obtain a unique combination of mechanical, thermal, and electrical characteristics. For example, in large-scale resistance welding, an increasing fraction of production work involves steels with zinc or zinc alloy coatings for corrosion protection in automotive applications. In parts joined by resistance microwelding, coatings or platings are often applied for improvement not only of corrosion but also electrical properties and these surface coatings often complicate the welding process. Surface coatings or platings usually significantly alter the bonding mechanisms compared to resistance microwelding of bare base metals. Figure 16.18 shows cross sections of resistance microwelded gold-plated nickel sheets joined with different welding currents [12]. Since gold does not form an oxide film and the electrical resistivity and hardness of gold are low, solid-state bonding between gold platings occurred initially in this example. Then melting initiated at the Au/Ni interface, resulting in a Au-Ni alloy braze. Finally, a Ni fusion nugget formed. In this case, gold plating increased the threshold welding current required to form a joint due to the reduction of contact resistance. However, the brazed joints between gold plated nickel sheets were stronger than comparable welds between bare nickel sheets even without formation of a fusion nugget.
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491
(a) I
400 µm 1622 A (b)
II
I
I
400 µm 2198 A (c)
I
II
II
III
I
400 µm
2563 A (d)
I
III
II
II
I
400 µm 2933 A (e)
III
400 µm
2120 A
16.18 Cross sections of (a) through (d) Au plated Ni and (e) bare Ni joints made with different welding current showing different type of bonds. I: solid state bond, II: braze, and III: fusion nugget. Large voids are indicated by arrows and part of the fusion boundaries are indicated by dash lines [12].
In cross-wire microwelding, changes in bonding mechanism could also occur with plating [33]. Figure 16.19 shows cross-wire joints for nickel wires with and without gold plating before and after load testing. Although the Ni wire joints showed initial necking on the base wire, the bonded interface had already separated upon loading due to the presence of a notch. In contrast, the fillet formed in gold plated Ni wires joint provided a larger
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16.19 Cross-wire joints of (a) bare Ni wires, (b) Au plated Ni wires welded at 600 A for 1 ms (1: as-welded joints, 2: after loading) [33].
bonded area and a smooth transition from one wire to the other. No separation at the bonded interface occurred, and the necking was clearly observed at the edge of the fillet. On the other hand, soft plating could promote electrode sticking. Interactions between the platings and electrodes will frequently accelerate the deterioration of the electrodes. Trials of process and electrode material changes have shown that electrode tip life and sticking tendencies can both be improved in resistance microwelding as well as resistance welding. For example, reducing welding current and weld time, and increasing electrode force and spacing could reduce electrode sticking. New electrodes with TiC metal matrix composite coatings have been developed and have demonstrated superior performance in welding plated materials [19, 34].
16.6
Dynamic resistance and process control
Studies of dynamic resistance during resistance welding, between electrodes and/or sheets, have been carried out in order to find ways to monitor and control the joining process [35, 36]. For example, in 1980, Dickinson proposed that dynamic resistance during resistance welding of steels could be broken down into recognizable stages, namely, surface breakdown, asperity softening, increasing bulk resistivity resulting from heating, initial melting, nugget growth, mechanical collapse and weld metal expulsion. Tracking of realtime resistance changes could thus be used to monitor and control the process
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493
[27]. Similarly, Tan et al. [11, 32] proposed that dynamic resistance in resistance microwelding of thin nickel sheets could also be used to monitor and control the process, using an analysis summarized as follows: dynamic resistance (R) between sheets consists of constriction resistance (RC), film resistances (RF) and bulk resistance. (RB) as shown in Eq. (5). R = RC + RF + F B
(5)
Resistance
The dominance of each component will usually change during the welding sequence. Each component is an independent variable: a function of temperature, material properties, electrode force and/or geometry. Figure 16.20 shows a schematic illustration of a typical sheet-to-sheet dynamic resistance curve during resistance microwelding of Ni (solid line) including the component resistances (broken lines) [32]. Four main stages are shown in the curve; the second stage involves another three sub-stages. Stage 1 is an asperity heating stage; resistance would increase as temperature increases since resistivity increases with temperature. Although resistance shows a monotone decrease at Stage 2, each component shows different behavior. In Stage 2a, the dynamic resistance drops because of electrical breakdown of surface films. The magnitude of decrease would depend on the thickness of oxide films. In Stage 2b, surface asperities begin to soften and constriction resistance to decrease shortly after the moment of film breakdown. In Stage 2c, the dynamic resistance, apparently affected increasingly by bulk resistance, continues to drop but at much lower rate and eventually reaches its minimum. At the B valley, faying surfaces start to melt. After the B valley, most contact resistance (film plus constriction) should have disappeared and the dynamic resistance begins to reflect essentially the behavior of bulk resistance. Since there is large resistivity difference between liquid and solid Ni phases, as the fusion nugget forms and grows, the dynamic resistance increases correspondingly until the nugget reaches its maximum diameter and/or
A
C
RF B
RC RB
1 2a 2b
2c
3
4
Time
16.20 Schematic showing a typical sheet-to-sheet dynamic resistance curve during resistance microwelding of bare Ni (solid line) and its components (broken line) [32].
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thickness at the C instant (Stage 3). In the Stage 4, the dynamic resistance starts to drop again since temperature starts to drop. When energy input is too high, the surrounding solid material cannot hold the liquid phase and weld metal expulsion occurs. The bulk resistance shows a sudden drop due to loss of material causing the shortening of the current path. As a result, a second peak in the dynamic resistance curve indicates the formation of fusion nugget, and this could be used as quality control input variable in resistance microwelding of Ni sheets. The first peak in dynamic resistance is not seen in resistance welding of steels [27]. However, the characteristic nature of dynamic resistance varies greatly depending on the material combinations of base metals and platings to be welded. For example, Fig. 16.21 shows the dynamic resistance curve at the faying interface during resistance microwelding of gold plated Ni sheets, which indicates the effect of gold plating on dynamic resistance as compared to Fig. 16.20 [11]. The curves can be divided to the following stages: asperity softening (2b), resistivity increase of Au and Ni due to temperature increase (2d), solid-state bonding (2e), brazing (2f) and nugget growth (3).
16.7
Equipment
16.7.1 Power supplies Compared to the dominant use of line-frequency controlled alternating welding current (a.c.) in resistance welding, resistance microwelding systems use a wider variety of power supplies [10]: line or high frequency a.c., capacitor discharge (CD), high frequency (HF) inverter, and direct current (d.c.). Microwelding inherently requires much more precise control of not only welding force but also electrical power, compared to large-scale applications. When an a.c. power supply is used (Fig. 16.22), the thermal power is controlled Resistance
C
RC
B
RB 2b
2d
2e
2f
3
Time
16.21 Schematic showing a typical sheet to sheet dynamic resistance curve during resistance microwelding of Au plated Ni (solid line) and its components (broken line) [11].
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CD HF
0
Φ
Time a.c.
16.22 Schematic of current waveforms of CD, HF and a.c. power supplies. The solid lines are the result of switching off current for a portion of each cycle, which compares the dashed a.c. lines with 100% heat (Φ = 0) [15].
by changing voltage and switching off the current for a portion of each cycle [15]. When using thyristor-based controllers, the minimum controllable heating unit is 12 cycle, that is, 8.3 ms or 10 ms duration, depending on a.c. frequency (60 or 50 Hz). Control resolution of other types of power supplies can be much better than with a.c. When a CD power supply is used, the energy is provided by a charged capacitor bank and the amount delivered is determined by the amplitude and duration of the current pulse (pulse width), which can be manipulated to any required resolution by changing circuit characteristics and charging voltage. Typical pulse width is on the order of 1–5 ms. The current waveform for HF inverter systems is typically d.c. with a superimposed high frequency, low amplitude a.c. ripple. Typical welding time is 10–30 ms and typical control resolution is 1 ms. Most metals could be welded successfully by all of the above types of power supplies, although the effect of power supply characteristics on the magnitude of welding current required to produce a given size of weld nugget can be significant [13, 15]. For example, the minimum welding current to produce 400 µm-diameter of weld nuggets are shown in Tables 16.2, 16.3 and 16.4 with different power supplies using Class 2 (chromium copper alloy) and Class 14 (molybdenum) electrodes [13, 14].
16.7.2 Electrodes The electrode materials used for resistance microwelding are generally precipitation-strengthened copper alloys such as Cu-Cr, Cu-Zr or Cu-Cr-Zr alloy, similar to large scale resistance welding. Others, such as Cu-W, CuMo alloys[13, 14]), are also used where resistance heating of the electrodes
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Table 16.2 Welding current (kA) for 0.2 mm thick sheet metals using the a.c. power supplier Sheet metals
Electrodes
Minimum nugget*
Weld expulsion
Electrode sticking
Suggested range
Al
Class Class Class Class Class Class
1.1 0.7 1.6 1.2 ≥3.5 ≥2.2
2.0 1.0 2.0 >1.8 3.8
>2.1 ~1.1 2.6 ~1.4 2.8 2.0
1.1–2.1 0.7–1.0 1.6–2.6 1.2–1.4
Brass Cu
2 14 2 14 2 14
Note: The minimum current is determined to produce 0.4 mm-diameter of weld nuggets. Electrode force is 4.5 kgf and weld time is 8 cycles.
Table 16.3 Welding current (kA) for 0.2 mm thick sheet metals using the HF inverter power supply Sheet metals
Electrodes
Al
Class Class Class Class Class Class
Brass Cu
2 14 2 14 2 14
Minimum nugget* 1.9 1.2 2.6 1.6
Weld expulsion
Electrode sticking
Suggested range
3.0 1.6 3.0 2.0
>3.0 1.4 >3.2 2.0
1.9–3.0 1.2–1.4 2.6–3.2 1.6–1.0
>3.3
*Note: The minimum current is determined to produce 0.4 mm-diameter of weld nuggets. Electrode force is 4.5 kgf and weld time is 20 ms.
Table 16.4 Welding energy (joules) for 0.2 mm thick sheet metals using the CD power supply Sheet metals
Electrodes
Minimum nugget*
Weld expulsion
Electrode sticking
Suggested range
Al
Class Class Class Class Class Class
30 30 35 35 >125 >125
70 70 70 70
100 50 80 40
30–100 30–50 35–80 35–40
Brass Cu
2 14 2 14 2 14
*Note: The minimum current is determined to produce 0.4 mm diameter of weld nuggets. Electrode force is 4.5 kgf and pulse width is 2.0 ms.
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is needed to increase the overall heat generation. To reduce electrode sticking and or improve electrode tip life, dispersion strengthened electrode material (such as Cu-Al2O3) or TiC composite-coated electrodes may also be used [18, 34]. Flat or radius ended electrodes are often used in resistance microwelding. Geometry and dimensions of the electrodes control the current density and the resulting feasible size range of the weld nugget. Thus small changes in electrode diameter would influence current density more significantly in comparison with large scale resistance welding. Therefore, electrode tip faces must be precisely machined and maintained. In addition, compared to large-scale welding where the electrodes are internally water cooled, the heat build-up at the electrode/sheet interfaces is worse in resistance microwelding since no water cooling is used and current density is generally higher [13, 14, 19]. Thus operators or maintenance technicians must be more aware of the implications of electrode sticking and resultant decrease of electrode tip life. In large-scale welding, when electrodes become deformed or alloyed with workpiece materials, tip faces are often cleaned using a hand file or sandpaper, which could alter the tip geometry and surface texture, causing inconsistent welding. In resistance microwelding, it is very important to use well-designed supporting blocks or fixtures when electrode tip faces are cleaned to ensure maintenance of a very consistent tip geometry [10].
16.8
Summary and future trends
This chapter has attempted to introduce various aspects of resistance microwelding. However, since the microwelding field is immature compared to resistance welding at larger scales, the information in this chapter is far from comprehensive. In this regard, much more systematic research and development work is needed to study process, materials and control issues in resistance microwelding. As precision components, devices and systems are further miniaturized, the parts to be welded and the joints in them continue to reduce in scale. This places ever-increasing requirements for better resolution (both electrical and mechanical) on machines, including power supplies and monitoring systems, and also more accurate controls on materials (such as surface roughness). In addition, resistance microwelding of dissimilar and especially biocompatible materials is increasingly required, and is introducing unique metallurgical challenges. For instance, resistance microwelding of TiNi alloys to stainless steel wires produces fusion nuggets of very problematic properties. For resistance microwelding, specific international standards on process, materials, and equipment are mostly lacking and urgently needed as well.
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16.9
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References
1. R.L. O’Brien: Welding Handbook, Vol. 2, 8th edn, American Welding Society, NY, 1991. 2. J.W.A. Hunt: Metal Construction and British Welding Journal, 5(6), 1973, 205–213. 3. G. Sitte: Welding and Cutting, 21(2), 1990, 13–22. 4. ASM International Handbook Committee: ASM Handbook, Vol. 6: Welding, Brazing, and Soldering, ASM International, The Materials Information Society, Materials Park, OH, 1993, p. 226. 5. W.R. Bratschum and J.L. Leicht, IEEE Transactions on Components, Hybrids, and Manufacturing Technology, 15 (6), 1992, 931–937. 6. H. Sasaki, S. Kazui, T. Shida and T. Shibata, Welding International, 5 (9), 1991, 697–700. 7. C.-L. Chu, Proceedings of the 2nd Microelectronic Packaging Technology Materials and Processes, Philadelphia, Pennsylvania, April 24–28, ASM International, OH, 1989. 8. N.B. Potluri: Welding Journal, 78(3), 1999, 39–42. 9. W.A. Baeslack III, P.S. Liu, P.R. Smith and J. Gould: Materials Characterization, 41(1), 1998, 41–51. 10. D. Steinmeier: Welding Journal, 77(7), 1998, 39–47. 11. Wen Tan: Small-Scale Reisistance Spot Welding of Thin Nickel Sheets, PhD. Thesis, University of Waterloo, Ontario, Canada, 2004. 12. W. Tan, Y. Zhou and H.W. Kerr: Metallurgical and Materials Transactions A, 33A 2002, 2667–2676. 13. Y. Zhou, S.J. Dong and K.J. Ely: Journal of Electronic Materials, 30(8), 2001, 1012–1020. 14. Y. Zhou, P. Gorman, W. Tan and K.J. Ely: Journal of Electronic Materials, 29(9) 2000, 1090–1099. 15. K.J. Ely and Y. Zhou: Science and Technology of Welding and Joining, 6(2), 2001, 63–72. 16. James A. Munford: NASA Tech Note, D3714 (1966), p. 1–23. 17. John J. Fendrock and Lazaro M. Hong: IEEE Transactions on Components, Hybrids and Manufacturing Technology, 13(2), 1990, 376–382. 18. B.H. Chang, M.V. Li and Y. Zhou: Science and Technology of Welding and Joining, 6(5), 2001, 273–280. 19. S.J. Dong, G.P. Kelkar and Y. Zhou: IEEE Transactions On Electronics Packaging Manufacturing, 25(4), 2002, 355–361. 20. Resistance Welder Manufactures’ Association (RWMA): Resistance Welding Manual, 4th edn, George H. Buchanan Co., Philadelphia, PA, 1989, Sec.1. 21. V.E. Ataush, E.G. Moskvin and V.P. Leonov: Welding International, 6(8), 1992, 624–627. 22. V.I. Stroev, V.E. Ataush and Y.A. Rudzit: Welding International, 14(6), 2000, 63–72. 23. ASM International Handbook Committee: ASM Handbook, Vol. 6: Welding, Brazing, and Soldering, ASM International, The Materials Information Society, Materials Park, OH, 1993, p. 230. 24. S. Fukumoto and Y. Zhou: Metallurgical and Materials Transactions A, 35A(10), 2004, 3165–3176. 25. V.E. Moravskii, V.N. Korzh and S.P. Svidergol: Automatic Welding, 33(9), 1980, 24–26.
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26. John J. Fendrock and Lazoro M. Hong: IEEE Transactions on Components, Hybrids and Manufacturing Technology, 13(2), 1990, 376–382. 27. D.W. Dickinison, J.E. Franklin and A. Stanya: Welding Research Supplement, 59(6), 1980, 170s–176s. 28. J.G. Kaiser, G.J. Dunn and T.W. Eagar: Welding Research Supplement, 61(6), 1982, 167s–174s. 29. B.H. Chang and Y. Zhou: Journal of Materials Processing Technology, 139, 2003, 635–641. 30. D. Baker et al.: Physical Design of Electronic System, Vol. 3, Prentice-Hall Inc., Englewood Cliffs, NJ, 1971. 31. R. Holm: Electric Contacts: Theory and Application, 4th edn, Springer-Verlag, New York, 1967. 32. W. Tan, Y. Zhou, H.W. Kerr and S. Lawson: Journal of Physics D: Applied Physics, 37, 2004, 1998–2008. 33. S. Fukumoto, Z. Chen and Y. Zhou: Metallurgical and Materials Transactions A, 36A(10), 2005, 2717–2724. 34. S.J. Dong and Y. Zhou: Metallurgical and Materials Transactions A, 34A(7), 2003, 1501–1511. 35. W.F. Hess and R.L. Ringer: Welding Journal, 17, 1938, 39s–48s. 36. W.L. Roberts, Welding Journal, 30, 1951, 1004–1019.
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17 Adhesive bonding S B Ö H M , E S TA M M E N , G H E M K E N , M W A G N E R , Technical University Braunschweig, Germany
17.1
Introduction
In precision engineering and micro system technology adhesive bonding is becoming more and more of a key technology. Adhesive bonding has a lot of potential in comparison to other joining methods, but the use of adhesives in microjoining still lags behind other methods. One of the most important reasons for this is the fact, that everyone thinks they can use adhesives, but in practice adhesive bonding demands considerable skill and knowledge. Adhesive bonds often fail because of the wrong process or the wrong adhesive selection. In this chapter, first the adhesive bond technology is described including the classification of adhesives dependent on the chemical base and the properties of adhesives. Then the joining process with surface treatment, dispensing and mixing of the adhesives are described. The final part of the chapter deals with information about the quality control of adhesive bonds and the selection of adhesives in association with the application, future trends and standards for adhesive bonding.
17.2
Adhesive bonding as joining technology
17.2.1 Definition of adhesive bonding According to DIN EN 1692 an adhesive is a non-metal substance, which is capable of holding two other materials together because of wetting the surfaces (adhesion) and its inner strength (cohesion) while resisting separation. Forces can be transmitted from one material to another by bonding substrates adhesively.
17.2.2 Properties and benefits of adhesive bonds Using adhesive bonding techniques offers the possibility to join dissimilar materials even if they are small, extremely thin or without thermal stability. 500 WPNL2204
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Table 17.1 shows the advantages and disadvantages of this technology [Hab06], also relevant for microbonding applications. In the past adhesives were designed with UV-stability, high transparency or a low thermal expansion coefficient. Nearly every material can be bonded in an acceptable and even automated way. Mechanical fastening is often limited by design performance or mechanical behaviour. Compared to conventional joining technologies, adhesive bonding offers many advantages for micro system production. Most types of materials can be joined to each other with reduced or even no stress. Usually a low process temperature reduces the load on the components while creating high temperature resistant bonds. Adhesive bonding is a flexible technique which can add additional functions via the adhesive. So, adhesive bonding can provide almost unlimited technical and economic advantages.
17.2.3 Principle of adhesion and cohesion The cohesion of an adhesive depends on its chemical nature, especially the nature of polymer, and its interactions between atoms and molecules within itself (in contrast to the adhesion). Cross-linked adhesives such as duromers show a higher cohesion than thermoplastic systems; the temperature dependence of mechanical behaviour is a consequence of the basic bonding type. Several theories for adhesion mechanisms have been proposed, but none is capable of explaining all types of adhesive interactions and bonding [Bro05, Yac02]. The adhesive mechanisms include adsorption theory, chemical bonding, Table 17.1 Advantages and disadvantages of adhesive bonding Advantages
Disadvantages
Bonding of dissimilar materials possible
Pre-treatment of substrate surfaces necessary Limited thermal stability
Low process temperatures; reduction of load on components; especially for thermosensitve materials Joint with reduced/no stress; consistent tension distribution upright to stress Introduction of additional functions via adhesive (e.g. conductivity); multifunctionality of joints Bonding of extremely thin/small materials possible High dynamic stability; high vibration absorption Fast processes; automation possible; flexible technique Freedom of design and innovative technolgy
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Low peel strength; tendency to creep Limited repair options
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diffusion theory, electrostatic attraction, mechanical interlocking and weak boundary layer theory. Additionally, various defects and contaminations might significantly affect adhesion characteristics. Bonding forces, coming from adhesion forces in the borderlines and cohesive forces in the adhesive generate the joint between two or more atomic groups in the molecules and the phase interfaces (Fig. 17.1). The physical and chemical properties of an adhesive in such an interfacial region may substantially differ from properties of the adhesive away from this contact area. The composition of this interphase region determines the strength and the durability of the resulting bond. The adsorption theory relates the adhesion to interatomic or intermolecular attractive forces between substrate and adhesive. According to this theory, the wetting of the substrate surface by the adhesive is a key factor (see below). So far, this theory has substantial experimental support. The chemical bonding theory proposes that all properties of the adhesive bond result from interfacial forces, valence bonds or van-der-Waals forces. According to this theory, bonding is a result of chemical reactions of adhesive molecules which are adsorbed on the substrate surface. Electrostatic forces develop at an interface between materials of different electronic band structures. If adhesives or substrates contain polar groups or permanent dipoles, electrostatic bonding is possible. The diffusion theory of adhesion is reliable only for polymer substrates and adhesives, because there is the need for interdiffusion of the polymer chains of molecules. This theory does not apply to ceramics and metals. The weak boundary layer theory relates weak adhesive strength and interfacial failure to the characteristics of the first few atomic layers around the interface. In microscopic scale the typical substrate surfaces are usually
Substrate 1
Borderline
Adhesion
Adhesive
Cohesion Borderline
Adhesion
Substrate 2
17.1 Structure of an adhesive bond.
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rough. This theory explains the failure not at the adhesion interface but within the adhesive or substrate. For this the formation of a weak boundary layer is necessary, which contains impurities, causing adverse conditions for chemical reactions. The above theories are combined in the term ‘specific adhesion’ and take place in interfaces between the substrate and adhesive of about 0.2–1.0 nm. It is necessary for the substrate and adhesive to come into very close contact because of this small region of influence, otherwise no interaction can take place. For this reason, contaminants have to be removed and the wetting properties of the substrate surface must be improved to generate an optimized bond. Good wetting properties are achieved by surface pre-treatments as described in Section 17.3.2. Wetting is defined as the process of spreading a liquid on a solid surface. This process is mainly controlled by the relationship between interfacial tensions at a point on a three-phase contact line, described by Young in 1805 [Mit93] (see Fig. 17.2): γsv = γsl + γlv cos θ (Young Equation; for equilibrium and θ > 0°) where s = solid; l = liquid; v = vapour; γsl, γsv, γ lv = interfacial tension between the two phases and θ = contact angle corresponding to the angle between vectors γlv and γsl. The interrelationship of wetting properties and substrate surface properties can be described by the contact angle measurement. The lower the contact angle, the greater the tendency for the liquid to wet the solid, until complete wetting occurs at an angle of θ = 0, (cos θ = 1). The surface tension of the liquid is then equal to the critical surface tension of the substrate. Large contact angles are associated with poor wettability [Bro05] (see Fig. 17.3). Two methods are commonly used for measuring contact angles on solid surfaces; the static method (known as sessile drop method or goniometry) and the dynamic method (known as Wilhelmy method or tensiometry). The static method requires the observation of a drop of test liquid on a solid substrate. In contrast the dynamic method measures forces of interaction as the solid is immersed in a test liquid. The methods are sensitive to interactions between liquid and solid, including swelling of the solid surfaces, roughness γlv Gas phase
γsv
θ
γsl
Liquid Solid
17.2 Relationship between interfacial tensions described by Young [Mit93].
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θ > 90°
Poor wetting
θ = 0–90°
Good wetting
θ = 0°
Complete wetting
17.3 Wetting of substrates.
and porosity [Hab06, Bro05, Dun03]. Another influence is the change of surface tensions of many test liquids caused by small quantities of surface active impurities which requires careful working and storage at constant test conditions, particularly measurement at defined room temperature. The ‘mechanical adhesion’ or mechanical interlocking, e.g. the positive fit between the liquid phase and macroscopic surface effects like pores, capillaries and undercut, is not nearly as reliable for surfaces at the microscopic scale.
17.2.4 Classification of adhesives Generally, adhesives may be classified in many different ways, such as by: • • • • • • •
origin (natural or synthetic adhesives) physical form or aggregate state (e.g., pastes, dispersions, liquids, oneor multipart components, films, tapes) chemical base or type (organic or inorganic adhesives; epoxy, silicone, polyurethane, or acrylic type) mode of reaction (chemical reaction such as polymerisation, polyaddition or polycondensation; physical reaction such as hotmelt, PSA or contact adhesive) curing method (e.g., heat curing, moisture curing, radiation curing) strength (structural, semi-structural or as sealant) end-use function (MEMS, medical applications, optics, electrically or thermally conductive).
Some classifications will be described in detail, but usually most of them are part of the classification by the mode of reaction. Classification by the mode of reaction The most common classification of adhesives is shown in Fig. 17.4. The class of physical setting adhesives include the important group of hotmelt
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Adhesive
Reactive hotmelt
Chemical reaction Curing at room temperature
Physically bonding
Hot curing
Contact adhesive (PSA) Hotmelt
Polymerization
Dispersion, solution
Polycondensation
Solvent borne adhesive
Polyaddition
Plastisol
17.4 Organogram: classification of adhesives.
adhesives. These are well known for rapid micro-assembly but have the disadvantage of low strength and limited temperature resistance of the joint. This is a result of thermoplastic behaviour which allows melting and solidification by the influence of heat and cooling energy. Most common thermoplastic hotmelt adhesives are polyamides, saturated polyesters, polyolefines, ethylene-vinylacetate-copolymers, blockpolymers such as styrenebutadiene-styrene and polyimides [Hab06, Bro05]. Pressure-sensitive adhesives are permanently in a tacky form at room temperature and can be used to join various materials with the application of moderate pressure. Pressure-sensitive adhesives are not normally used in sustained-load-bearing applications and can often be removed without leaving a residue. Most pressure-sensitive adhesives are based on elastomers (natural or synthetic rubbers), acrylics or hotmelt thermoplastics, tackifiers, or antioxidants. Single- or double-coated pressure-sensitive adhesives are employed in many medical applications as well as in the electronic industry because of the thin bondline and good flexibility [Tav01]. Contact adhesives are high molecular weight polymers without chemical crosslinking. Through solvent use (organic solution or water dispersion) they are given a low viscosity which is able to wet the substrate surface. Both bonding partners have to be wetted by the adhesive; after drying the two parts are joined together while interdiffusion between the adhesive layers forms the joint [Hab06]. Plastisol adhesives are PVC polymers, dispersed in a high boiling-point liquid (diluent). They form a paste with low molecular weight which can be heat activated. Plastisols are applied like hotmelt adhesives. The polymer particles dissolve in the diluent and form a viscoplastic synthetic material. Their typical use is bonding metals in automotive applications [Hab06].
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In contrast, the class of reactive hotmelts are a link between physical setting and chemical reaction in the adhesive. Through the crosslinking reaction, the hotmelt becomes infusible and insoluble. Polyurethanes with solid polyester polyols as the main component for the prepolymer reaction are a typical example. By the influence of humidity, crosslinking occurs. A new class are the pre-applicable polyurethane hotmelt films containing solid isocyanate crosslinker which are activated by heat above the melting temperature [Bue05]. Another example is the reaction between thermoplastic polyamide and isocyanate crosslinkers into humidity-resistant networks. Classification by curing mode Polymerization, polyaddition and polycondensation are the three types of different adhesive curing modes [Hab06]. One or two component polymerization reactions are induced by radicals, catalysts or irradiation such as UV or electron beam which start a radical, anionic or cationic curing mechanism. Monomers with carbon-carbon double bonds (vinyl-groups or others) are bifunctional through this activation. They are known as acrylic-based adhesives or acrylates. There is a wide range of acrylic based adhesives that join a variety of similar and dissimilar materials. The main types of acrylic based adhesives are cyanoacrylates, anaerobics, and modified acrylics. These adhesives are usually available as solvent-based liquids, emulsions, tapes, or monomer-polymer mixtures (one- or two-part components), with liquid or powder curing agents. Acrylic-based adhesives may be polymerized or cured using moisture, catalysts, heat, ultraviolet or visible light, or other sources of radiation. Because of the excellent electrical properties that are stable with respect to ageing, acrylics are of interest in the electrical and electronics industries as well as in the medical industry. Cyanoacrylates, an important class of acrylic adhesives, are one-component adhesives and are commonly known as superglue. Cyanoacrylates are ester derivates of α-cyanoacrylic acid. Because of the cyano- and ester function, the charge equation of the double bond is displaced and negative atoms coming from dissociated water, surrounding moisture or weak bases in the atmosphere or on the substrate surface, can start the ionic polymerization. Because of their fast setting, cyanoacrylates filled with silver or nickel, can be used as electrically conductive adhesives [Hab06]. Anaerobic adhesives are one-component formulations that cure in the absence of oxygen under the influence of metal ions. They are used for bonding metal substrates. Activation energy is low so curing reactions start at room temperature. Typical use is bonding of ferrite cores in the electronic industry [Del07]. For radiation curing, monomers or pre-polymers with active carbon-
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carbon double bonds like vinyl groups (acrylate, methacrylate) are essential. The steps of this curing mode are as follows: 1. emission of radiation 2. energy absorption through monomer molecules or photoinitiators (with photochemical primary reaction) 3. creation of free radicals R• 4. initiation of reactive system, creation of start radicals R• + CH2 = CH2 → R-CH2- CH •2 5. start of chain reaction R-CH2- CH •2 + CH2=CH2 → R-CH2-CH2-CH2- CH •2 6. propagation to crosslinked adhesive 7. end of propagation through combination of polymer molecules or recombination of free radicals. Monomer reactivity determines the rate of reaction. Free radicals of oxygen can work as inhibitor. Usually UV light works in a range of 230–400 nm, depending on the source of light. UV curing is easy to use, especially for glass bonding. A combination of UV curing and other curing methods (such as thermal, anaerobe) offers the possibility of decreasing disadvantages including shadow effects [Hab06]. Typical two-component adhesives with polymerization curing modes are methacrylates with a hardener component instead of the photoinitiator. Physical setting adhesives with polymerization reactions are polyvinylalcohol, polyvinylether and ethylene-vinylacetates, polyvinylchloride and polyolefins. The class of natural and synthetic rubbers include carbon-carbon double bonds in monomers as either isoprene diene, styrene or nitrile. Copolymerisates are known as thermoplastic elastomers (styrene-butadiene-rubber, SBR). Typical formulations are styrene-block-polymers, butyl rubber, or nitrilebutadiene-rubber (NBR). In polyaddition reactions the connection of monomer groups does not involve a splitting of a carbon-carbon double bond but an adsorption of reactive monomer molecules while an active hydrogen migrates from one reactive partner to another. The most important polyaddition adhesives are epoxy resins and polyurethanes. Epoxy resins are thermoset adhesives based on the epoxide group (also known as the oxirane group), a three-membered carbon, carbon, and oxygen ring structure, as shown in Fig. 17.5. The ability H
H C
C
H
O
17.5 Epoxid.
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of this group to undergo a variety of reactions and crosslinking leads to multiple epoxy resins with a wide range of chemical and physical properties and molecular weights and structures. Best known are epoxy adhesives based on bisphenol A and epichlorohydrin. For polyaddition reactions compounds with mobile hydrogens are necessary, e.g. amines, carbonic acids anhydrides or resins such as phenolic resins. The oxirane ring opens while forming a hydroxyl group with the mobile hydrogen, the addition of the molecular group occurs to the free valences (Fig. 17.6). Epoxies are among the most widely used adhesives for both structural and nonstructural applications. They are commercially available as liquids, pastes, films and solids in one-component, heat-curing form or as a two-component adhesive, curing at room or elevated temperatures. Systems curing at elevated temperatures generally have a higher crosslinking density and glass transition temperature, which provides better shear strength and environmental resistance. Epoxies are used in a number of medical or electrical devices for bonding and sealing applications, e.g. fixing of SMD components or for flip chip contacting. In this cases radiation curing epoxies with photo initiators are used [Del07]. The second important group of polyaddition adhesives are polyurethanes. Most commercially available systems are based on polyethers or polyesters with terminating hydroxyl functional groups. The reaction of an alcohol and an isocyanate results in the formation of a urethane (Fig. 17.7). The materials used in most polyurethane systems usually consist of one of several different formulations: di- or polyfunctional alcohols; polyhydroxy compounds (known
H
H C
C
H H
H
X
H C
O
C
HO
X H
X = e.g. –NH–R; –OOC–R; –O–R
17.6 Epoxid curing.
R
N
C
O
R
R
NH2
O
H
N
C
H
O
C
N
OH
R
N
C
H
O
OH
OH
R
NH2
R
R
N
C
N
H
O
H
17.7 Urethane formation.
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as polyols); di- or polyfunctional isocyanates; or low-molecular-weight alcohols or amines. One-component polyurethanes cure in the presence of humidity, as shown in Fig. 17.8. After one-component reactions urethane (coming from prepolymers) (Fig. 17.9), and urea compounds (coming from hydrolysis) form the cured polymer side by side (Fig. 17.10). Two-component polyurethane adhesives are less common. Well known are applications with polyurethane dispersions, especially water dispersed or reactive urethane hotmelts [Hab06, Bro05]. During a polycondensation reaction low molecular weight compounds, e.g. water, alcohol or acids separate and exist alongside the forming polymer. Polycondensation comprises the formation of the basic adhesive polymer (e.g., in polyamide and polyesters) as well as the reaction during adhesive curing (e.g., in silicones or formaldehyde resins). One of the first adhesives used in the aircraft industry was based on formaldehyde resins which are formed by the reaction of formaldehyde and phenol or compounds with amino group (melamine, urea). In microsystem technologies hotmelt adhesives on polyamide and polyester bases are of growing interest. Low activation temperatures and fast setting are goals when bonding thermally sensitive actuators or MEMS in rapid bonding processes [Boe06, Boe05]. Silicons or silicon rubbers cure by the polycondensation mechanism but differ totally from organo-carbon polymers. Typical is the silicon-oxygen structure (siloxane) with carbon-hydrogen (functional) groups. Setting usually occurs at room temperature, curing in the presence of humidity. Applications are as under-filler in electronic devices or basic formulations for PSAs, but the main use is as a sealant in combination with suitable fillers.
R1
N
C
O
R2
OH
R1
17.8 Urethane curing.
R1
N
C
N
H
O
H
R2
17.9 Urethane compound.
R1
N
C
H
O
O
R2
17.10 Urea compound.
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O
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Additives for adhesives In adhesive applications additives and filler materials are commonly used to add defined features to mechanical, physical or chemical performance [Hab06]. In adhesive formulations additives and fillers must be incorporated homogeneously. Their main properties are for modifying chemical configuration and inert behaviour, grading, density, wetting, thermal and electrical conductivity and the coefficient of thermal expansion. They are able to expand the temperature range of end-use or adhesive resistance, they can reduce shrinkage, influence rheology and application behaviour and give defined electrical and thermal conductivity. For bulk applications, inorganic mostly crystalline particles like quartz flour, chalk, glass fibres, mica and metal powder are used. New adhesive features and applications became possible when nano-sized particles became part of adhesive formulations. Isotropic conductive adhesives with an electrical conductivity in the x-, y- and z-axis by contact of the filler particles, are in the majority of cases epoxy resins (one or two component) and to a lesser extent acrylic-, silicon- or polyimide-based. The basic resin formulation typically has a specific electrical resistance of 1012–1015 Ωcm, which can be reduced by addition of metal particles to 10–3–10–4 Ωcm. Additives are silver and gold flakes or discs, or with reduced conductivity, copper, nickel and carbon. Additionally silver-coated glass spheres or gold-coated polystyrene spheres are used. The filler ratio in cured adhesives is about 60–80% by weight with a common particle size between 10 and 50 µm. Nanoscale silver-particles are about 50–150 nm in size and are made by inert gas condensation. Their porosity allows a reduced filler concentration [Hab06]. Anisotropic conductive adhesives have a conductivity only in the z-axis which is generated during the bonding process by pressure and temperature. The filler particles do not contact in the uncured adhesive (so the filling degree is clearly lower compared to isotropic conductive adhesives) and the thermoplastic (polyester) or reactive (epoxy) adhesive functions like an isolator. So in the direction of conductivity an electrical resistance of 1 × 10–4 Ω cm is possible with 1 to 2 × 1014 Ωcm in the cross direction. Fillers are gold, silver, graphite, solder powder, silver-coated nickel particles, noble metal coated polymer particles or glass spheres, in a size of 10–15 µm [Hab06]. For jet dispensing conductive adhesives, developed in the last few years, filler size has to be 5 µm or smaller [Kol05]. Adhesives with thermal conductivity, commonly used in electrical applications, can be filled (naturally with metal fillers with electrical conductivity) with aluminium oxide, aluminium nitride or boron nitride with filler ratio in cured adhesives of 60–75% by weight. Typically epoxy resins are used [Hab06].
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17.2.5 Characteristics of adhesive joints One of the main characteristics of a polymer material is the glass transition temperature [Yac02]. The mechanical properties can vary below and above this temperature range. For fibre-optic and micro-electro-mechanical devices the operating temperature range is important to know. Thermal analysis methods are used to measure this property, e.g. by change in heat capacity, elastic modulus, coefficient of thermal expansion or dielectric constant. The glass transition temperature is not a change in phase, e.g. melting or crystallization of the polymer, but is a kinetic event that depends on the timescale on which a property of the material is observed. Below the glass transition temperature the polymer chains are unable to reorganize within the timescale of the experimental measurement. Each thermal analysis method has its own way of heating the sample, under the influence of different stresses, glass temperature range for the same adhesive can thus differ from one method to another. For microsystem technology the reproducibility of components is important. One major factor is to make sure the degree of cure is homogeneous which means a complete cure without crosslinking reactions following on. As well as methods like destructive and non-destructive testing, a routinely employed technique is the use of differential scanning calorimetry (DSC), which provides measurements of heat capacity and endothermic or exothermic reactions. As a result the glass transition temperature and the degree of cure can be determined. For measurements a sample is heated or cooled at a controlled rate while the flow of energy in the sample is monitored as a function of time or temperature. Typically, the electrical energy supplied to a heating element of a furnace containing the sample is compared to that supplied to an identical furnace used as a reference. Consequently, the heat capacity of the sample and thermodynamic effects like glass transition effects, crystallization, melting, crosslinking reaction, decomposition and outgassing are detectable. The mechanical properties of a material can be measured by using dynamic mechanical analysis (DMA) and thermal mechanical analysis (TMA). In DMA, the viscoelastic response of a material to a dynamic load at a fixed frequency is measured as a function of temperature or time. The well defined test specimen is heated or cooled at a controlled rate under oscillating stress with fixed amplitude. This gives information about the storage and loss properties and thereby tan δ as their ratio. The influence of humidity can degrade the adhesive insulating characteristics, especially in electrical devices. Moisture leads to corrosion of metallic components and a breakdown in performance of the whole system. To protect the joint, adhesive formulators have to consider moisture resistance and outdoor weathering or other environmental factors. In medical applications
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this can lead to extreme environmental effects. Certain types of medicalgrade epoxy adhesives are capable of being sterilized by autoclaving and chemical methods. These epoxy resins can be used in medical devices that require sterilization prior to use.
17.2.6 Adhesive bonding as microjoining technology Microadhesive bonding offers the possibility of creating new innovative technical solutions. For the assembly of microparts, adhesive bonding is often the only way to join dissimilar materials. The multi-functionality of the joint (e.g., mechanical fixing plus conductivity, even in dynamically stressed systems) is a main objective for making processes effective, faster and leads to freedoms of design. Low process temperatures, stability to temperature and humidity and even reduced shrinkage are possible. It has to be mentioned that the properties of adhesives in the microscale are different from those in the macroscopic scale. Adhesives which join two substrates well in the macroscopic scale can be completely unsuitable in microscopic applications (e.g., because of their high viscosity or the size of the filler material). The smallest quantities have to be dispensed and placed precisely. The precise volumetric repeatability of adhesive dispensing is necessary to create products of consistent high quality and performance. Capillary effects cause problems if selective bonding is required. Currently, the assembly line for mobile phone production has a typical movement reproducibility of 20 µm. Special assembly lines are offered with resolutions of about 10 µm. On granite plates at constant temperature these systems achieve resolutions of about 5 µm. Next generation systems will require 1 µm or even better. Thermal expansion in the classical design of meter sized robots and actuators already avoid precise positioning [Nan07].
17.3
Adhesive bonding process: pre-treatment
17.3.1 Requirements of surface treatment Successful adhesive joining normally requires suitable surface treatment of the substrates prior to bonding. The selection and application of an appropriate surface treatment is one of the major factors for achieving good wettability and improved long-term durability in adhesively bonded joints. Inadequate or unsuitable surface treatment is one of the most common causes of degradation and failure. The function of surface treatment includes the removal of contaminants or weak boundary layers and the alteration of surface chemistry, topography, and morphology to enhance adhesion and durability.
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17.3.2 Pre-treatment methods Surface preparation techniques are generally divided into mechanical or chemical methods [Hab06, Bro05, Dun03]. Mechanical methods especially in macroscopic scale include abrasion, grit blasting, and shot blasting or even laser technologies. Chemical methods include degreasing, etching, and anodizing, the use of adhesion promoters and flame, corona, and plasma treatments. The first step of any surface preparation is cleaning with organic solvents such as acetone, isopropanol or methylethylketone. For environmental reasons the use of chlorinated solvents is being phased out. The solvent is used in an ultrasonic bath or by wiping with a commercially available cloth. Most of the organic contaminations are removed in this first step so further pretreatments can work more effectively. In many applications, simple degreasing and abrasion is sufficient to provide good adhesion. Polymers with low surface energy and bondability like those used in medical applications require a more specialized treatment, e.g. plasma or corona treatments, in order to provide better adhesion and joint durability [Tav01]. Electrical corona discharge is a dry surface preparation method consisting of a high-potential (typically 20 kV) stationary electrode, connected to a high-frequency generator (10–20 kHz) and an earthed treater roll, electrically insulated with a dielectric covering as the other electrode. The result is an intense electrical field that ionized the surrounding air. The ionized particles, e.g. free electrons and ions, bombard and penetrate the molecular structure of the surface. By reaction with molecular oxygen from the air, polar chemical groups are formed which deliver new bonding partners for the adhesive. Usually, primers are used to provide a compatible surface for the adhesive and to protect established surfaces from corrosion. For this purpose primers are low viscous systems with a good ability to wet the surfaces. Silanes and epoxy primers are used to improve durability and to provide a barrier to moisture permeation along the adhesive interface when bonding to metals or glass substrates. In case of silane primers the molecules are able to form strong chemical bonds with especially, glass substrates through hydrolyzable alkoxy groups. In addition those molecules are able to polymerize side groups and physically form bonds with the adhesive. To modify the surface chemistry of polymers, chloride-containing primers are applied especially on vulcanized and thermoplastic rubbers. In the past years, new grades of adhesives with the ability to bond low surface energy polyolefins like polyethylene or polypropylene, without specific pre-treatment, have become commercially available [Pro07].
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17.4
Dispensing and mixing of adhesives
Adhesive dispensing is a key process in adhesive bonding. It guarantees correct transfer of adhesive to different locations on the substrates and ensures that the adhesive remains in place until hardened. Adhesive application can be grouped into two main categories: • •
mass dispensing, such as pin-transfer and screen-printing selective dispensing, such as needle-based or jet-based dispensing techniques.
Selective dispensing can be classified into two techniques: • •
contact-based, in which the dispensing device actually touches the board non-contact which avoids physical contact with the board.
17.4.1 Mass dispensing techniques Pin transfer This is one of the fastest methods of applying patterns of adhesive to surfaces (e.g., printed circuit boards). The pin array is dipped into an open tray of adhesive to wet the pins in predictable amounts. The wetted pins are then touched onto the substrate and the adhesive transfers to the board (see Fig. 17.11). Thus, whole boards can be treated in a single operation. Pin transfer tends to be used in high volume applications that have long production runs. This mass dispensing technique has several drawbacks. The tooling must be redone if the pattern is changed and this can be expensive. This technique requires a low-viscous adhesive with good resistance to moisture absorption because of the exposure to ambient conditions in the bath. It is difficult for the pin transfer method to apply the adhesive to package dimensions smaller than 1206 (3.2 mm × 1.6 mm) [Pir00]. Advantages: very fast, consistent; disadvantages: difficult/expensive retooling and adhesive handling. Pin
Substrate
Tray of adhesive
17.11 Principle of pin-transfer.
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Screen/stencil printing This is a rapid method of applying surface mount adhesives. Screen printing uses a stencil or screen with holes that are designed specifically to match the dot pattern required on the substrate. To apply the adhesive, the stencil is aligned on the board and a squeegee or a roller wipes the adhesive across the stencil, forcing the adhesive through the holes and onto the board (Fig. 17.12). The stencil is lifted from the surface and the adhesive is pulled from the apertures to form an array of adhesive dots on the surface. Screen printing is a practical and cost-effective method of rapidly applying adhesives. The size and shape of the dots can be controlled through the geometry of the stencil apertures (principally stencil thickness and aperture diameter) and the rheology of the adhesive. Dot height may be expected to be nearly constant to the stencil thickness but it is known that the aperture diameter plays a role. Screen printing is even more sensitive to viscosity changes of the adhesive through ambient temperature variations and partial curing than the pin transfer method [Cla03]. Stencil printing has some disadvantages as storage space is required to store stencils and tooling when not in use, and cleaning the stencil or tooling after a production run can be time-consuming. New stencils will be required each time the dispensing pattern is changed. Advantages: fast, can have multiple dot sizes; disadvantages: retooling, adhesive handling, cleaning process, storage.
17.4.2 Selective dispensing Needle-based selective dispensing techniques have been the preferred method for most adhesives such as surface mount adhesive (SMA)-dispensing applications (e.g. underfill) over a long time. They offer more flexibility for adapting the dispensing process to different board designs, e.g. a new pinarray or stencil. In needle-based dispensing, precision motion systems move the needle to a certain location. The needle is then positioned above the substrate typically by using a physical standoff mechanism to achieve the
Squeegee Stencil
Substrate
17.12 Screen printing.
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correct dispense height. A controlled amount of adhesive is dispensed using one of several proven techniques: • • • •
time/pressure pump auger pump piston pump jet dispensing (contactless).
Time/pressure dispensing The adhesive is contained in a pressurized syringe and the flow is controlled using air pulsed through a nozzle valve to dispense the desired amount of adhesive (see Fig. 17.13). This method is simple and reliable; rates as high as 40,000 dots per hour can be achieved. As the adhesive level in the syringe gets lower, the column of air in the syringe becomes larger and the change in pressurization time leads to dot inconsistency [Dun03]. The time/pressure method has difficulties in producing consistent dots in 0603-size (1.6 mm × 0.8 mm) components. Time/pressure dispensing is a well proven method of dispensing that has a simple and time efficient clean up stage that involves disposing of the syringe and cleaning or replacing the nozzle. In comparison with all other dispensing methods for surface mount adhesive applications it has the advantage of flexibility in changes to the application. Normally a relatively simple automated control platform-software is the solution. Advantages: flexible, simple operation, easy cleanup; disadvantages: speed affects consistency. Auger pump dispensing This uses a pump with a rotating screw thread to displace the adhesive within the reservoir. Turning the screw’s electric motor on and off dispenses Pressure
Nozzle
Substrate
17.13 Time/pressure-dispenser.
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measured amounts of the adhesive. The motion of the pump will shear a thixotropic adhesive. Once the fluid in the pump contacts the surface, a consistent flow will be maintained. This method works with a wide range of adhesives. The restricting factor on dot dispensing speed is three-dimensional motion control of the robotic machinery. The auger pump will dispense dots in the approximate range of the time/pressure method but with more consistency. Advantages: flexible, handles wide range of adhesives, less sensitive to air in the fluid, better than time pressure; disadvantages: sensitive to viscosity changes, some affect of speed on consistency. Piston pump dispensing This is a positive displacement method of adhesive application where the movement of a piston in a closed fluid-chamber dispenses a volume of adhesive precisely determined by the volume the piston displaces. In normal operation, viscosity changes or needle size have little effect on the piston pump flow rate. In most dispensing applications, the pump is set to a fixed displacement to give a specified shot size. However, continuous extrusion is possible for other applications. As with the auger pump, the nozzle contacts the surface. A disadvantage to the piston pump is the complex cleaning method required on a daily basis, compared to the time/pressure and the auger pump method. Advantages: consistency at higher speeds, capable of large shot sizes; disadvantages: complex cleaning procedures, sensitive to air in the fluid. Jet dispensing This is a contactless method of applying adhesive using a piezo-actuator driven pin to force the adhesive through the nozzle in a rapidly cycled mode. It is similar to the fixed volume piston pump with a high velocity that jets the material. The adhesive is ejected from the nozzle and accelerated onto the surface, forming a dot (see Fig. 17.14). As the nozzle does not move in the z-axis, significant increases in the number of dots per hour over auger and piston pumps are possible. Jet dispensing can give small dot profiles (diameter of dots down to 50 µm), which is especially advantageous for smaller components. A limitation is that the jet can dispense only one dot size without changing equipment settings. Non-contact jetting provides advantages such as easy programming, fast set-up and robust process control in addition to higher dispensing rates [Cla03]. Advantages: non-contact dispensing, less sensitive to board warpage or dispense height variation, faster than other dispensing methods; disadvantages: one dot size per jet, complex clean-up.
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Piezo-actuator
Substrate
17.14 Jet-dispensing.
17.4.3 Mixing The mixing of the adhesive initiates the cure reaction. The time from mixing until the mixed adhesive gels and becomes too viscous to use is known as the work life. Work lives for adhesives depend on the formulation/application and can vary from seconds to hours. Many types of adhesive are supplied as two-component systems, a resin and a hardener, which are mixed in order to initiate the cure reaction to harden the adhesive. The performance of the bonds made from two-part adhesives depends on having the correct ratio of components intimately mixed and applied within the working life of the system. Mix ratios of adhesive components can vary enormously. Ratios of one part resin to one part hardener (1:1) are common. In most applications the adhesive components are mixed before they are applied to the substrates. This can be done from dual pack cartridges or automated dispensing equipment using nozzles containing in-line mixing elements. A static mixing system is used if the materials to be mixed exhibit similar viscosities and have a relatively long pot life. It is usually used within a mixing ratio of 100:100 to 100:20. The static mixing system consists of a pneumatically actuated twin snuffer valve fitted with a static plastic mixing tube. The static mixer, also referred to as a motionless mixer, is a narrow cylindrical tube which contains the mixing elements. The elements are stationary parts that are positioned to force the materials to combine as they travel through the length of the mixer. After stirring through the length of the mixer, the mixed material is forced through the tip of the mixer to the point of application. Another type of mixing is the dynamic method (see Fig. 17.15). Mixers of this type are used if difficult media with widely dissimilar viscosities and a very short pot life have to be mixed. It is used with all mixing ratios. Pneumatic or electric motors are used to drive the rotor in the mixing chamber, then the adhesive is forced to the substrate through a nozzle. Adhesives can also be mixed in bulk. Adhesives must then be decanted from their packaging and
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Hardener DC servo motor
Mixing chamber for different ratios
Substrate
17.15 Principle of dynamic 2K-dispensing.
mixed in bulk before application. A third method of mixing two-component adhesives is through the application of the components (resin and hardener/ activator) onto the different surfaces to be bonded. The mixing and preparation of adhesives should be in accordance with manufacturers’ instructions. Only adhesives from the same batch number should be used in a single joint to prevent uneven properties. Preparing and testing adhesive joint specimens for bond performance can also be done to assess mix quality. Thermal analysis methods such as DSC or DMA can be used to check mix consistency in the final cured material.
17.4.4 Adhesive packages There are many different packaging types for single- and dual-component adhesives, e.g. bottles, cartridges, jars, tubes, bulks, single and dual barrel syringes.
17.5
Curing and setting of adhesives
There are many different types of curing techniques. The curing of adhesives can be achieved using moisture or catalysts in the presence or absence of air (e.g., in the absence of air for anaerobics) at room temperature; thermally, at elevated temperatures; photochemically, using irradiation (e.g., by UV or visible light), induction or microwaves. The basis of these techniques is briefly described in the following sections.
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17.5.1 Irradiation There is a range of materials available that can be cured using radiation sources such as UV and visible light. Acrylated resins can be cured using radiation energy. A typical UV energy of 80–120 mW/cm2 produced from a UV source (300–400 nm wavelength) is usually sufficient to cure a UVcurable adhesive within 10 to 60 seconds. An alternative radiation curing technique is to use visible light (470 nm wavelength). Radiation curing adhesives are often used for joining clear polymers in disposable and nondisposable medical devices. In general, both UV and visible-light curing can be achieved using light boxes or focused beams and light guides. In some cases heat is also employed to encourage the curing process, for the completion of curing, or to cure areas that cannot be reached by radiation energy.
17.5.2 Heat There are three mechanisms by which heat (energy) is transferred into the adhesvies: 1. irradiation (non-contact) 2. conduction (contact-based) 3. convection (non-contact). Heat cured adhesives can be either one-part or two-part products. Heatactivated curing can be initiated by using the following methods: • • • • • • • •
local heat application at the joints (thermal conduction) hot air heating oven (convection) hot-plate heating (conduction) infra-red heating (irradiation) vapour-phase heating liquid-phase heating laser heating (Nd:YAG or CO2-lasers) microwave – particularly variable-frequency microwave (VFM).
17.5.3 Moisture A room temperature cure process can be used for adhesives that are either moisture cure or condensation cure products. After application, the adhesive is simply allowed to cure at ambient room conditions. Condensation cure products will cure in 30 minutes to 24 hours. Moisture cure products will require several hours to cure or need to be left undisturbed overnight as the cure is dependent on moisture from the air permeating through the material.
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17.5.4 Curing hastening processes Induction An effective way of speeding the cure of conventional adhesives has resulted from looking at the heating process itself. Electromagnetic heating provides reliable, repeatable, non-contact and energy efficient heat in a minimal amount of time. Electric power is used to generate heat in conducting materials (e.g., steel or aluminium) placed in close proximity to an inductor coil through which an alternating current is passed. The process is, therefore, limited to applications when at least one substrate is metal or to adhesive materials that are filled with electrically conducting powder. Speed is the most important feature of induction heating. Work piece heating rates greater than 40 °C/second are possible. The heat induced in the work piece will conduct instantly into the adhesive providing the catalyst for curing. If one of the two substrates is not electrically conductive then an adhesive can be used which includes a small percentage of metal oxide particles. The metal oxide particles within the adhesive become heated in an induction field and provide the source of heat to cure the matrix material in the adhesive [Pet04]. Radio frequency and microwave curing Short length, electromagnetic waves also present a fast method for curing structural adhesives. This method of heating is also known as microwave heating. Like induction curing, RF and microwave curing have been investigated as an alternative to conventional processing. Such processing is more energy efficient and has reduced processing time when compared to conventional heating methods. There are basically two forms of dielectric heating: radio frequency and microwave. Radio frequency heating uses a frequency (13–100 MHz) to generate heat in polar materials. Microwave heating uses high-frequency (2–20 GHz) electromagnetic radiation to heat a polar material or a material with a susceptor located at the joint interface. The susceptor material absorbs the microwave energy and converts it into heat. Radio frequency heating is often used on thermoplastics with a susceptor embedded in the material and on water-based adhesives, such as acrylic, starch, and polyvinyl acetate. The RF energy is used to drive off the water quickly and dry the adhesive. Microwave heating is generally used on structural adhesives, such as the epoxies, where heating occurs via rapid oscillation of the polar groups in the field [Sau04].
17.6
Quality control
The best method of quality assurance is to produce quality not to test it. In Section 17.3 it is outlined why it is necessary to treat surfaces, how to WPNL2204
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calculate dosages and mix adhesives and how to cure adhesives quickly. In this section possible defects and their reasons are described and how the adhesive bond process can minimize the danger of bond defects. Further, it will give a survey of which test procedures exist to test the quality of adhesive bonds.
17.6.1 Defects due to adhesive bonding Defects due to adhesive bonding are listed below [Gar01]. •
•
•
•
Unsuitable choice of adhesive: the chosen adhesive does not have the needed properties, e.g. the mechanical properties including strength or elasticity are unsuitable or the adhesive bond process does not fit the right process for the adhesive. Wrong dosage: in micro adhesive bonding, the adhesive bond volume is very low. Often the sizes of the adhesive bonds are in the area of mm2 or the volume in the range of nanolite. So an overdose can contaminate the adherents and influence the function of the component. An underdose can negatively influence the mechanical properties of the component. Bad curing: depending on the chemical properties of the adhesive, the curing of the adhesive can differ considerably. For example, adhesives used in microjoining often cure due to UV light emission or by mixing two components. Here the danger of bad curing is possible because the UV light can be (partly) shaded, the dosage can be too low or the mixture of the two components (ratio or homogeneity) is not adequate. Poor positioning: often in microjoining, the positioning of the adhesive is poor. The demand for accurate positioning increases with decreasing size of the components.
17.6.2 Test methods In principle, test methods can be classified into destructive testing (DT) and non-destructive testing (NDT) methods. For adhesive bonding, DT methods serve for the determination of the bond strength. Such tests can be static, dynamic or highly dynamic (crash). In Section 17.11.1 all relevant test methods are listed. NDT methods for adhesive bonding can be classified into six categories: Mechanical vibrations, electromagnetic, thermal, penetrating irradiation, visual and chemical-electrochemical.
17.7
Principles of adhesive selection and examples of applications
Compared to other joining technologies, adhesive joining has some advantages. It is possible to join different materials with low tension forces at low processing
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temperatures. On the other hand, the working temperature and resistance to temperature alternation are fairly high. Additional functions besides the joining function are possible including heat conduction or protection against environmental influences. As described in the preceding sections, a wide range of adhesives and technologies are available. The purpose of this section is to give guidelines for choosing an appropriate adhesive technology and some ideas of practical applications by discussing some examples.
17.7.1 Choosing the right adhesive technology The selection of the appropriated adhesive and bonding technology is influenced by a wide range of parameters to incorporate. First it is necessary to analyse the gluing components. The second step is to search for a suitable adhesive system. The third is to qualify the best choice. The fourth is to optimize the process and fifth is to choose the right technology, equipment and preparation. Other considerations are protection, biological compatibility, long-term resistance and durability against environmental influences. The process of curing or setting needs a considerable time, depending on the adhesive and the chosen technological parameters. The technological choice is to hold the assembled parts in place or to stitch the parts with another joining technology. The main criteria are accessibility and manageability of the parts, processability, possibility of process control, potential strength of joint, and economic efficiency. An often used substrate material in microjoining is silicon. Typical joining partners are glass, plastic or metals. There are four main goals for the use of adhesives: 1. mechanical fixation and setting of parts from the same and different materials 2. functional connection such as thermal, optical, fluidal, chemical or mechanical interconnection 3. protection from environmental influences such as temperature, humidity, corrosives, vibrations or, electromagnetic radiation 4. customer benefit through high reliability and steadiness at lowest product costs. UV curing of one-component epoxies is very common in electronics. Other typical adhesives for fixation or sealing are polyurethane solution adhesives (PUR) and two-component epoxies. Acrylates are also commonly used, e.g. cyanoacrylates (CA) which are well known as ‘superglue’. To improve the reaction rate it is usually possible to support the reaction by heating or in some cases by addition of an accelerating agent. In practice it is often possible to use the adhesive system straight out of the box. Some systems, e.g. one-component systems, need no extra preparation. Other
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adhesives are prepared by the manufacturer to minimize the effort with integrated mixing devices. In the case of dispersion adhesives it is often necessary to agitate the mixture. In rare cases it is necessary to mix the components in an extra step. The preparation for hot melts is to heat the glue in a tank or cartridge and to keep it hot in the product pipeline including the dispensing unit.
17.7.2 Health and safety A very important aspect of using adhesives is occupational health and safety. Because of the chemicals and substances used there is often a potential hazard to employees and the environment. It is necessary to instruct the users appropriately. It is important to use personal protective equipment, e.g. protective goggles, gloves and clothing, and to comply with the appropriate laws, directives and rules. Used substances require waste management to prevent environmental contamination. Depending on the adhesive and technology, the use of an exhaust or good ventilation may be required. In the case of UV-curing adhesives it is prohibited to discard the uncured resin. During processing the worker has to protect eyes and skin and has to wear UV protective goggles while the resin is exposed to UV. An extra exhaust is usually not necessary. Most adhesives come with a material safety data sheet and a technical data sheet. The material safety data sheet (MSDS) contains information about manufacturer, instructions for safe use, transport, handling and storage, first aid in case of accident or misuse, fire fighting, disposal and other factors in a formal manner. It is an official document and required if chemicals are handled. It helps if the information is required quickly. The technical data sheet gives some information about properties, typical applications and processing instructions for the adhesive. The manufacturer is free to select and to arrange the given information. This document is primarily a marketing tool.
17.8
Microelectronic interconnections and packaging
The demands for adhesives in microjoining are easy processing with common dispensing technologies such as printing, stamping or dosage and well known curing methods. Often adhesives are used for stitching purposes or sealing and for structural joints with added functional benefits.
17.8.1 Chip bonding For bonding of the die to a substrate, the so-called chip-on-board (COB) technology, there are several possibilities such as adhesive bonding or soldering. Often the substrate is a lead frame with metal leads that extend outside the WPNL2204
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housing of the chip package, e.g. dual in line package (DIP). Another type of common substrate for hybrid technology or higher integrated circuits is the printed circuit board (PCB). As described in [Hil06] the major technology for die bonding with 99% is adhesive bonding (Figs 17.16 and 17.17). The main adhesives are one- or two-component epoxies filled up to 80% with silver for thermal conductivity. The film thickness is usually 25 µm applied by pad printing or stamping to the substrate. The chip is then picked and placed by a die bonder machine with an accuracy of around 25 µm. The curing is carried out at higher temperatures.
17.8.2 PCB mounting of SMDs A surface mounted device (SMD) is usually glued onto the PCB to keep the component in place during the process of wave or reflow soldering. In
17.16 Dispensing an adhesive for bonding.
17.17 Bonded die.
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electronics this is the main application for unfilled adhesives (Figs 17.18, 17.19, and 17.20).
17.8.3 Flip chip assembly A common technology for bonding dies or integrated circuit chips (IC chips) to a substrate like a PCB is the flip chip assembly (FCA). With flip chip technology it is possible to create a surface mounted device with a high density of connections per area. The main idea is to have a grid of contact pads, so-called ball grid arrays (BGA), and to connect it to contact pads on the surface of the PCB. This provides an advantage in size, cost and performance. Beside soldering with soldering bumps in a reflow process it is possible to glue the component with a non- or anisotropic conducting adhesive
17.18 Dispensing an adhesive for bonding.
17.19 Surface mounted device bonded to the PCB.
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17.20 Bumps of adhesive and stuck SMDs before soldering (photograph by courtesy of Henkel KGaA).
to the substrate. The main applications are flip chip in package (FCIP) and flip chip on board (FCOB).
17.8.4 Sealants and sealings The technology of sealing is an important field in micro electronics and mechanics. The checklist for choosing the right adhesive includes the protective influence such as moisture, chemicals, light, the temperature range, electrical properties, and other requirements such as electrical, chemical and mechanical compatibility or matching the coefficient of expansion. Suitable adhesive materials can be acrylic, polyurethane, silicone, etc. Processing parameters are to clarify rework of applied material, curing time, type of dosage process, needed equipment and the price of the adhesive.
17.8.5 Glob-top for chip-on-board Glob topping is a conformal coating to protect a device against mechanical damage, exposure to weather and dirt. Usually a bare chip is directly mounted onto and connected to a PCB by wire bonds or flip chip technology, (Fig. 17.21). To protect the COB assembly from environmental impact it is sealed
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17.21 Chip mounted on board and connected.
under an adhesive (Figs 17.22 and 17.23). The challenge is to choose a specially formulated resin with a viscosity high enough to prevent the sealant from spreading but low enough to easily fill all gaps.
17.8.6 Dam-and-fill coating With higher density of chips or extended size of the component, the problem of spreading becomes unsolvable. Here, the dam-and-fill technology is the solution. A dam from an adhesive with a paste-like viscosity surrounds the device (Figs 17.24 and 17.25) and a low viscous adhesive is filled inside the dam (Figs 17.26 and 17.27). The last step is to cure the adhesive.
17.8.7 Underfill of flip chips Flip chip assembly provides an advantage in size, cost and performance. For protection and mechanical enhancement flip chips with BGA normally undergo an underfill process which will cover the sides of the die. The sealing nonconductive adhesive will fill each gap under the component by capillary forces (Fig. 17.28).
17.8.8 Contacts with non-conductive adhesives The simplest method for contacting two contacts is to put them together. To keep them in touch an adhesive is added by underfill technology (see above). While curing, the adhesive between the substrates shrinks and forces the contacts even closer together (Fig. 17.29). Usually, a special form of contact is applied, so-called stud bumps, made from metal or from isotropic conductive adhesives.
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17.22 Dispensing the sealing.
17.23 Glob-top completed.
17.8.9 Contacts with isotropic conductive adhesives An isotropic conductive adhesive (ICA) is one of the oldest functional adhesives used in electronics. The main applications are die bonding on lead frames (see above) and the connection of a surface mounted device, electrically conducted by an adhesive, to the printed circuit board (Figs 17.30, 17.31 and 17.32). The meaning of isotropic is conductivity equal in all directions. Fillers are gold or silver in the form of flakes. The filling rate is at 60–80% weight by weight (w/w). The contact of the flakes to each other results in the conductivity of the adhesive (Fig. 17.33). Isotropically conductive film (ICF) is usually used to provide shielding
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17.24 Building a dam.
17.25 PCB with dam completed (photograph by courtesy of Henkel KGaA).
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17.26 Filling in the sealing adhesive.
17.27 PCB Dam and fill coating completed (photograph by courtesy of Henkel KGaA).
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17.28 FCBGA with underfill (photograph by courtesy of Henkel KGaA).
17.29 Substrates with contact pads joined by an underfill from a non conductive adhesive (sectional drawing).
materials with self-adhesive film. The benefit is an electrical connection through the film to the PCB for example.
17.8.10 Contacts with anisotropic conductive adhesives Advantages of conductive adhesives compared to soldering are lower processing temperatures, suitability for fine pitch applications, high reliability and easier processing in most cases. An anisotropic conductive adhesive (ACA) is often used to connect different conducting materials. One type of conductive adhesive is film with conducting particles. Such anisotropic conductive films (ACF) are used for joining heat and pressure flexible PCBs
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17.30 Dispensing the ICA to the pads.
17.31 SMD glued to the PCB with conductive adhesive.
to rigid PCB or glass with metallised conducting circuits, e.g. displays with fine pitch connection [Due05]. The function of ACA is as follows. An insulating adhesive with conductive particles is, under the influence of pressure and heat, pressed between the contacts. The pressure results in pressing the particles into the pads, giving conduction between the pads. The adhesive keeps the tractive force for sustaining the conduction and mechanical joint after curing (Fig. 17.34).
17.8.11 Thermal conductive adhesives Thermal conductive adhesives (TCA) are for the improvement of the heat transfer by closing insulation gaps caused by surface roughness of the parts
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17.32 Sample of glued contacts of an SMD IC (photograph by courtesy of Polytec PT GmbH).
17.33 A contact joined to an other contact on a substrate with ICA (sectional drawing).
to be joined. They contain heat conducting fillers like ceramic particles or metal flakes in different sizes, to optimize the heat transfer, embedded in an organic polymeric matrix. Due to the adhesive feature, an additional mechanical fixation often becomes redundant.
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17.34 Contacts on substrates contacted by ACA (sectional drawing).
17.9
Other applications of adhesive bonding
In the following section are some sample applications and common technologies associated with adhesive bonding in microjoining.
17.9.1 RFID RFID stands for radio frequency identification device and usually consists of an antenna coil, contacts, a capacitor and an integrated circuit chip. The chip is often a smart device with its own processing capabilities. The antenna serves for information interconnection and to supply the chip with energy. The typical technology for assembly of a RFID label is gluing with anisotropic conductive adhesive or with non-conductive adhesive. Different technologies are used to dispense the adhesive. The bonding procedure needs a flip chip bonder to realise the accuracy, pressure and temperature until the adhesive is cured. This takes approximately 5–20 seconds. Roll-to-roll processing is the best way to create the labels or tags.
17.9.2 Roll-to-roll processing Roll-to-roll processing is the process of creating electronic devices on a roll of flexible plastic or metal foil. Roll-to-roll is a description of the manufacturing process. Because the length of the rolls can be up to 50 metres it is possible to create a continuous process compared to the batch process. Printing, bonding, punching and other technologies can easily be involved in the process. Typical applications are solar cells, RFID labels or RFID tags or smart cards. Possible fields for future applications are functional clothing or large-area flexible displays.
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17.9.3 Batch processing Batch processing can be defined as accomplishing several processing technologies step by step with a batch of parts or devices at the same time. In micro technologies this unit is usually a substrate (or a batch of substrates) with a large amount of the same components, e.g. a silicon wafer in microelectronics to keep down the costs.
17.9.4 Optoelectronics interconnections and packaging Packaging and interconnection techniques in electronics manufacturing are playing a key role for miniaturization of microsystem devices and for increasing the density of functions and components in electronic products. Nowadays, one of the most important factors is the information carrier medium – light. Typical examples for miniaturized optical components are optical interconnects like fibre-space connections via prism or fibre-fibre coupling. In information technology, light is the information transmitter of the century. To join optical components such as optical fibres and lenses, a wide range of specially conditioned adhesives for microjoining purposes are on the market. A stress-free junction with adhesives avoids failures and serves for the success of the product. Other examples are waveguide optics, e.g. planar waveguides, switches and fibre alignment. With special adhesives low-loss, transparent and light conducting bonds can be formed. One example for using adhesive joints in optoelectronic interconnection is a cylindrical endoscope objective. High-performance endoscopes are essential for minimal-invasive surgery. For this optical application different kinds of glasses/lenses (with different surfaces and thermal expansion coefficients) were built up with different diameters in an optical path. The objective is ground by a grinding process to a uniform diameter. This procedure alone places enormous stress on the adhesive bonds. Highly transparent, UV-curing adhesives with adapted refractive indices are used for the assembly of the complex optical components.
17.9.5 Medical applications In the area of electronic components for medical devices there is a wide range of polymeric materials available [Tav01]. They are used variously as attachments, substrates and interconnections and for encapsulation or protection. In recent years, the use of adhesives and encapsulants in medical electronics and implantable devices has increased considerably. This is due to the availability of a wide range of materials that offer different properties, better adhesion, improved durability, suitability for automated dispensing, and rapid curing. Typical examples are the use of epoxies in ultrasound catheters and pacemakers and light curing adhesives in medical electronic packages. One
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example of such a device is based on potting the stainless-steel or titanium access port with an epoxy adhesive or encapsulant. This particular adhesive passed all USP Class VI biological and toxicity tests (i.e., acute systemic and intracutaneous toxicity, implantation tests, and cytotoxicity tests). For minimal invasive surgery the new high-performance endoscopes are indispensable. Lenses with diameters of less than 3 mm are bonded into an optical complex by using highly transparent adhesives with adapted refractive indices. Furthermore, lancets, syringes, injectors, hypodermics, blood collection sets and catheters are assembled using acrylic based adhesives [Tav01].
17.10 Future trends The high degree of miniaturization of MEMS and MOEMS and the wish to reduce their production costs require joining processes that can be used in batch processes, that allow short process cycles and that are able to interconnect the smallest bond areas. A newly developed interconnecting technology with hot melt adhesives fits all these requirements. Instead of the commonly used viscous adhesives this interconnection technology uses hot melt adhesives. An important advantage of hot melts in relation to viscous adhesive systems is the possibility of being able to pre-apply hot melt systems, e.g. as a powder or adhesive ball, in fused form, dispersion or as an adhesive foil. Further on, the joining procedure does not have to take place directly after applying of the hot melt adhesives to the substrate. Hot melts possess no pot life time, after application the bonding process can happen at any time later. The hot melt is melted during the bonding process by a thermal impulse. Due to cooling, the adhesive sets. Different filler materials enable hot melts to interconnect MEMS and MOEMS electrically, optically or thermally, while being conductive with high bonding strength and good ageing behavior. Hot melts usually possess a very small set shrinkage in comparison to chemical curing adhesives. First industrial applications show the extremely high relevance of this interconnecting technology for the MEMS and MOEMS industry. Miniaturization down to nano scale parts require more than a scaling of common adhesive technologies. Side effects such as electrostatic or van der Waals’ forces and the influence of molecular structures gain in importance. One nanometre is the distance of a few atoms, e.g. the diameter of a DNA double helix is only two nanometres. We are at the start of the interesting field of nano research in adhesive technologies. Chemical analyses and biochemical applications with microfluidic systems like ‘lab-on-a-chip’ boosted the needs for resistant and biocompatible materials for bonding, packaging and assembling at low temperatures. Adhesives are the solution for assembly of micro parts made of silicon, glass or polymers. Finding better adhesives is the subject of present research.
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17.10.1 Application methods for hot melts due to MEMS and MOEMS One of the most important advantages of hot melts in comparison to viscous adhesive systems is the possibility of pre-applying hot melt systems, e.g. as powder or adhesive spheres, as dispersion or as an adhesive foil, (Fig. 17.35(a)–(d). In order to use hot melts in spherical form or as foil in microsystem engineering, particles of less than 100 µm must be used. Handling the smallest particles, however, causes problems due to the fact that the forces based on weight do not dominate compared to its behaviour in the macroscopic scale. With smaller parts, the influence of the Van-der-Waals forces, the electrostatic forces and the adhesion forces increase. This influence becomes apparent in the behaviour of different parts. For hot melts the application of the particles can be achieved by heating the substrate’s surface. A vacuum grip with a small capillary (< 100 µm) is used to grip hot melt balls and foils to deposit the particles while thermally supported. By using hot melt powder the pre-application of hot melts can be realized by sintering the powder with a focused laser. The surface, which has been applied homogeneously with hot melt powder, is scanned by a focused laser beam. Because of the thermal impulse the adhesive melts and sets on the joining surface, (Fig. 17.36). A following cleaning process removes the remaining, non-adhesive powder. The stencil or screen printing of hot melt dispersions is one of the interesting processes of pre-applying hot melts as the first industrial applications have shown. The powder fraction is the basis of a dispersion which is printed by fine pitch printer on a substrate, e.g. a wafer. The following process included dehydration of the hot melt dispersion and heating to melt the adhesive (Fig. 17.37). The printing process is reproducible and low budget, because it is very similar to the printing of solder paste in the reflow process.
17.10.2 Bonding process using pre-applied hot melts The adhesive is melted by a thermal impulse only during the bonding process and moistens the surface of the other substrate. The heating can be accomplished directly by heating the adhesive or indirectly by heating the substrate. Due to cooling, the adhesive sets. During suitable heat conduction hot melts set very fast, i.e., a handling strength (usually the ultimate strength) can be achieved in less than one second as experiments have shown.
17.11 Sources of further information 17.11.1 Adhesive bond standards There are many standards for adhesive testing, processing, handling of adhesives, etc. There are no international standards dedicated to adhesives in WPNL2204
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17.35 Methods to pre-apply hot melts (a) laser-sintering of hot melt powder (b) adhesive spheres (c) stencil or screen printing of hot melt dispersions (d) application of adhesives in foil form.
microjoining. Therefore the following list gives only an overview and is not specialized on microassembly. Note, as stated before, in some cases it is not possible to transfer the knowledge and specifications to micro adhesion. The information given applies particularly to standards valid for Germany [Beu07]. Since every nation uses national and international standards, the following list gives an incomplete but exemplary overview only.
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17.11.2 Definitions DIN 8580 – Manufacturing processes – Terms and definitions, division. DIN 8593 – Manufacturing processes joining. EN 26927 – Building construction; jointing products; sealants; vocabulary (ISO 6927); German version DIN EN 26927. EN ISO 10365 – Adhesives – Designation of main failure patterns (ISO 10365); German version DIN EN ISO 10365. WPNL2204
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17.36 Laser melted powder, diameter of structure about 150 µm.
17.37 Fine pitch structures of pre-applied hot melt dispersion.
17.11.3 Stress tests EN ISO 527 – Plastics – Determination of tensile properties (ISO 527); German version DIN EN ISO 527. EN 1465 – Adhesives – Determination of tensile lap-shear strength of rigidto-rigid bonded assemblies (ISO 4587); German version DIN EN 1465.
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DIN 54455 – Testing of adhesives for metals and of bonded metal joints; torsional shear test. EN ISO 10964 – Adhesives – Determination of torque strength of anaerobic adhesives on threaded fasteners (ISO 10964); German version DIN EN ISO 10964. pr EN 3793 – Aerospace series – Anaerobic polymerisable compounds; test method – Determination of static shear strength (draft). DIN 54452 –Testing of adhesives for metals and adhesively bonded metal joints; compression shear test. DIN 54461 – Structural adhesives – Testing of adhesive bonds – Bending peel test. DIN 53293 – Testing of sandwiches; Bending test. EN 1966 – Structural adhesives – Characterization of a surface by measuring adhesion by means of the three point bending method; German version DIN EN 1966. EN 1464 – Adhesives – Determination of peel resistance of high-strength adhesive bonds – Floating roller method (ISO 4578); German version DIN EN 1464. EN 1895 – Adhesives for paper and board, packaging and disposable sanitary products – 180° ‘T’ peel test for a flexible-to-flexible assembly; German version DIN EN 1895. DIN 53295 – Testing of sandwiches; Peel test by means of a drum (LN 53295 – Testing of adhesives for metallic honeycomb sandwiches; climbing drum peel test). DIN 65448 – Aerospace; structural adhesives; wedge test. EN 14444 – Structural adhesives – Qualitative assessment of durability of bonded assemblies – Wedge rupture test (ISO 10354); German version DIN EN 14444. EN 15109 – Adhesives – Determination of the time to rupture of bonded joints under static load (ISO 15109); German version DIN EN 15336. EN ISO 9664 – Adhesives – Test methods for fatigue properties of structural adhesives in tensile shear (ISO 9664); German version DIN EN ISO 9664.
17.11.4 Ageing tests EN ISO 9142 – Adhesives – Guide to the selection of standard laboratory ageing conditions for testing bonded joints (ISO 9142); German version DIN EN ISO 9142. DIN 54456 – Testing of structural adhesives – Test of resistance to climatic conditions. EN 60068 – Environmental testing (IEC 60068); German version DIN EN 60068.
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DIN 50014 – Climates and their technical application; standard atmospheres. VDA 621-415 – corrosion fatigue test (automotive standard). VW P 1200 – climatic alternating test (automotive standard). Besides all these standards there are further standards for qualifying an adhesive joint and there are an unmanageable number of company standards. In commerce, there are safety-related rules with respect to adhesive such as directives of professional organizations or occupational medical guidelines like accident prevention regulations or maximum permissible values, e.g. maximum allowable concentration (MAC) or threshold limit values (TLV). The commission directive 91/155/EEC defining and laying down the detailed arrangements for the system of specific information relating to dangerous preparations in implementation of Article 10 of Directive 88/379/EEC (German version 91/155/EWG) with obligations for identification marking and transmission of the material safety data sheet (MSDS) is very important. The newest European Union law in this context is about the Registration, Evaluation, and Authorization of Chemicals (REACH) from December 2006.
17.12 References [Beu07] Beuth Homepage, Beuth Verlag GmbH, Berlin, 2007, http://www.beuth.de/. [Boe05] S. Böhm, K. Dilger, E. Stammen, J. Hesselbach, J. Wrege: Adhesive Joints for MEMS using Hotmelts; mst news No. 1, 2005. [Boe06] S. Böhm, G. Hemken, E. Stammen, K. Dilger: Micro Bonding using Hot Melt Adhesives, Journal of Adhesion and Interface, Vol. 7, No. 4, 2006. [Bro05] W. Brockmann, P.L. Geiß, J. Klingen, B. Schröder: Klebtechnik; Wiley-VCH Weinheim, 2005. [Bue05] J. Büchner, W. Henning, H. Stepanski, B. Raffel: Lagerstabile latent-reaktive Klebfolien und ihre Einsatzchancen, Adhäsion – Kleben und Dichten 7-8/2005. [Cla03] Clayton, J. et al., Screen Printable Polymers for Wafer Packaging, Polymer Assembly Technology, Billerica, MA, 2003. [Del07] http://www.delo.de. [Due05] Rainer Dübbers, Einsatz von Klebstoff-Filmen in der Elektronik, MSTI-Konferenz, Klebtechnik in der Elektrotechnik, Stuttgart, 2005 [Dun03] B. Duncan, S. Abbott, R. Court, R. Roberts and D. Leatherdale: A Review of Adhesive Bonding – Assembly Processes and Measurement Methods; NPL Report MATC (A) 135, Crown copyright 2003, published online at www.npl.co.uk. [Gar01] Gartner, J.: Qualitätssicherung bei der automatisierten Applikation hochviskoser Dichtungen. TU München, Diss. München: Utz Wissenschaft 2001. (iwb Forschungsberichte 160). [Hab06] G. Habenicht: Kleben: Grundlagen, Technologien, Anwendungen; 5. Auflage, Springer Verlag, Berlin-Heidelberg, 2006. [Hil06] Ulrich Hilleringmann, Mikrosystemtechnik, B.G. Teubner Verlag, Wiesbaden, 2006. [Kol05] J. Kolbe, A. Arp, F. Calderone, E.M. Meyer, W. Meyer, H. Schaefer, M. Stuve:
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[Mit93] [Nan07] [Pet04] [Pir00] [Pro07] [Sau04] [Tav01] [Yac02]
Microjoining and nanojoining Inkjettable conductive adhesive for use in microelectronics and Microsystems technology, IEEE Polytronic 2005 Conference. K.L. Mittal (publ.): Contact Angle, Wettability and Adhesion; VSP BV, UtrechtTokyo, 1993. http://www.nanomotor.de, Klocke Nanotechnik Aachen, Germany Petrie, E. et al., Electromagnetic Curing of Structual Adhesives, Adhesives and Sealents, published online at www.specialchem4adhesives.com, 2004. Piracci, A.-F., Practical Production Applications for Jetting Technology, APEXCongress 2000, Long Beach, 2000. http://products3.3m.com; 3M USA Sauer, H.-M. et al., Beschleunigte Klebstoffhärtung mit Mikrowellen und Nanoferriten, Adhäsion – Kleben und Dichten, Vieweg Verlag, 2004. M. Tavakoli: The Adhesive Bonding of Medical Devices, Medical Device & Diagnostics Industry 2001; http://devicelink.com. B.G. Yacobi, S. Martin, K. Davis, A. Hudson, M. Hubert: Adhesive bonding in microelectronics and photonics; Journal of applied physics, Vol 91, No. 10, p. 6227–6262, 2002.
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18 Introduction to nanojoining S S A H I N, Celal Bayar University, Turkey and M Y AV U Z and Y Z H O U, University of Waterloo, Canada
18.1
Introduction
There is a growing interest in 0- and 1-D nanostructures (nanoparticles and nanowires and tubes, respectively) with functional electronic characteristics, since the assembly of 0- and 1-D materials into nanoscale devices and circuits could enable diverse applications in nanoelectronics and photonics [1]. Most of these applications involve the use of hetero-structures in which materials of different compositions meet at interfaces [2]. For instance, nanoparticles are fundamental building blocks for nano-engineering, with applications in different areas from biosensors to electronic nano-devices. They have a variety of unique spectroscopic, electronic, and chemical properties that originate from their small sizes and high surface/volume ratios [3–6]. To fabricate the nano-devices from the ‘parts’, it is necessary to interconnect the parts to have ohmic nano-contact. Individual semiconducting nanowires have already been configured as field-effect transistors, photo-detectors and bio/chemical sensors [7–9]. Moreover, more sophisticated light-emitting diodes and complementary and diode logic devices have been manufactured using different type of semiconducting nanowires or nanotubes. However, as indicated above, to realize electronic applications, such as quantum wires [10], ballistic conductors [11] microchip interconnects [12] and transistors [13], the need to reproducibly fabricate connections between individual nanotubes or wires and electrodes has been identified as a major impediment [14]. For devices of any complexity, both high repeatability and selectivity are essential. To obtain nano-electronic circuits, a technique is required that enables them to connect with each other and with their periphery [15]. However, for instance, as previous studies show, making an electrically conductive connection between nanotubes is not straightforward. Instead of the desired ohmic contacts, tunnel junctions are often generated. Studies of unmodified crossed singlewalled metallic nanotube junctions show a resistance of approximately 200 kΩ [16]. Such a high resistance has to be interpreted in view of the small contact area of the order of 1 nm2, where crossing tubes touch at one point 545 WPNL2204
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only [15]. Improvements may be expected when a connection by a technique such as welding, soldering, or pasting can be established that joins the tubes with an electrically conductive material over a larger surface area. Conventional techniques such as pasting with conductive epoxy or soldering with liquid metals (like tin) are hardly applicable due to the small dimensions of the arrangements. Until recently, although multiple-way junctions of carbon nanotubes have been synthesized [17], using a simple chemical vapour deposition (CVD) method without use of any template, the formation of these different junctions is still random. In addition, several ideas for post-growth modifications have also been proposed [18, 19], such as welding, but it seems that they did not allow one to make a more complicated network of joint nanotubes with predetermined positions of junctions between nanotubes [20–23]. This chapter aims to give readers a comprehensive summary of the current techniques for the joining of nano-materials rather than growing them in a vacuum by electro-deposition. But it is important to state that some of the nano-joining techniques briefly described here have also been widely used for metallic and nonmetallic nanowires.
18.2
Nanojoining methods
18.2.1 Joining by electron beam nano-welding Mechanisms of irradiation Irradiation of CNTs with energetic particles has successfully been used for creating molecular joining between nanotubes [19, 21, 24–26]. When an energetic particle such as an electron or ion hits the target, different mechanisms of damage creation can work. Depending on the target material and the particle characteristics, the main mechanism can be the kinetic energy transfer, electronic excitations, ionization, etc. For CNTs, the most important mechanism is the knock-on atom displacements due to kinetic energy transfer, both for electrons and ions [27]. To understand the damage mechanisms in CNTs, it is necessary to study the atomic structure of CNTs in detail. The atom networks of typical CNTs are shown in Fig. 18.1. The CNTs may have only one shell (single-walled nanotubes, SWNTs) or many shells (multiwalled nanotubes, MWNTs). The former can bundle up to form a triangular lattice due to attractive van der Waals interactions (see Fig. 18.1(b)). However, all these structures retain graphitic arrangements of carbon atoms. In SWNTs, the collision of an energetic particle with a carbon atom will result in displacement of the atom, i.e., formation of single or multi-vacancies, and a number of primary knock-on atoms. If the energy of these atoms is high, they leave the tube or displace other atoms in the SWNT. If it is low,
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18.1 Molecular models of irradiated carbon nanotubes. A short fragment of a single-walled nanotube with a vacancy (V) and a double-coordinated carbon atom (A) adsorbed onto the outer surface of the tube (a), a bundle of nanotubes (view along the tube axes) after the impact of a 500 eV Ar ion (b), a multi-walled nanotube before and after 300 eV Ar ion irradiation with a dose of 2 × 1016 cm2 (c) [27].
they can adsorb onto the tube walls. These adsorbed atoms (adatoms) play the role of interstitials [28, 29]. It is important to note that due to the quasi one-dimensional morphology, all displaced atoms can be sputtered from the SWNT, so that no interstitial can exist in the system. Defects in CNTs can be created only if the incident particle is energetic enough to displace carbon atoms. For graphitic structures, the threshold energy Td (the minimum energy transferred to the atom required to produce the Frenkel pair which does not spontaneously recombine) is estimated to be about 20 eV (different values from 15 up to 30 eV have been reported, see [30] and references therein). By knowing this value, therefore, the minimum kinetic energy of the incident particle required for the damage production can be estimated. The high resolution transmission electron microscope (HRTEM) has usually been used for these studies since it can not only create the damage in CNTs by energetic (up to 1.25 MeV) electrons but also monitor the irradiationinduced changes. The minimum incident electron energy required to remove
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a carbon atom by a knock-on collision was found to be 86 keV, which correlates with the values of Td given in the literature for graphitic structures [31]. Early experiments have provided evidence that SWNTs exposed to focused electron irradiation were severely locally deformed and they developed necklike features along their bodies due to the removal of carbon atoms by knock-on displacements. Uniform irradiation of SWNTs [20, 32] also resulted in surface reconstruction and drastic dimensional changes, as a corollary of which the apparent diameter of the CNTs shrank from ≈ 1.4 to 0.4 nm. The reason for these transformations was vacancies in the walls of the SWNTs created by energetic electrons. The vacancies proved to be unstable under high beam dose when a large number of atoms are removed very rapidly and at high temperatures, which gave rise to surface reconstructions and diameter reductions. It is important to notice that vacancies are created in CNTs spatially non-uniformly and carbon atoms are most rapidly removed from surfaces lying normal to the electron beam [31]. These vacancy transformations were found to be the driving force for the irradiation-mediated coalescence of SWNTs [33]. Vacancies were found to induce coalescence via a zipper-like mechanism, imposing a continuous reorganization of atoms in individual tube lattices along adjacent tubes. Creation of vacancies by the focused electron beam followed by dangling bond saturation was demonstrated to result in welding of crossed SWNTs to form molecular junctions [21]. Electron irradiation of MWNTs have been found to result in forming vacancies on their walls and eventual amorphization upon high-dose irradiation [30, 34]. In general, MWNTs seem to be more stable than SWNTs because Frenkel pairs created inside the MWNT can easily recombine [30]. Annealing effect of irradiation-induced defects Defects in CNTs can also anneal, but it happens at elevated temperatures, and the annealing mechanisms are somewhat different from those in metals [30]. Experiments on electron irradiation of both SWNTs and MWNTs have indicated that irradiation induced damage in nanotubes can easily be annealed at temperatures higher than 300 °C [30]. Two mechanisms seem to govern defective annealing [29]. The first mechanism is vacancy healing through dangling bond saturation and by forming non-hexagonal rings and single wall (SW) defects. An illustration of this mechanism is given in Fig. 18.2 where the front walls of one and the same SWNT just after ion impact (a) and after annealing (b) are shown. It was seen that during annealing the double vacancy in the middle of the carbon network has been transformed to an agglomeration of non-hexagonal rings. The annealing also gave rise to the transformation of the single vacancy and the nearby carbon adatom in the upper right-hand corner of the network
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7 5
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18.2 Front walls of one and the same SWNT just after ion impact (a) and after annealing (b). During annealing the double vacancy (D) in the middle of the carbon network transformed to an agglomeration of non-hexagonal rings. The single vacancy (S) and the nearby carbon adatom (A) in the upper right-hand corner of the network transformed to a Stone–Wales 5–7 defect [27].
to a SW defect. It is also noticeable that the annealing led to the local diameter reduction. The annealing should eventually result in disappearance of the SW defect, especially if an extra carbon adatom (which works as the catalyst for the transformation thus substantially reducing the defect annihilation barrier) is nearby [35]. Similar vacancy-mending occurs upon electron irradiation [20]. It is important to note that ion irradiation creates more severe local damage than electron irradiation, since in the former case a vacancy clusters may easily be formed by an energetic ion impact, whereas in the latter case the predominant defects are single vacancies. Nevertheless, this mechanism works for ion irradiation as well. The second mechanism of annealing is the migration of carbon interstitials,
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followed by Frenkel pair recombination. The interstitial can migrate over the surface of SWNTs (isolated or bundled-up). Early calculations indicated that the migration energy Em is very low, but more rigorous recent results give higher values (0.5–1 eV; the actual value depends on the tube diameter and chirality), which is in good agreement with experimental values of Em, ~0.8 eV [30, 36–39]. Migration of interstitials in the open spaces between the adjacent shells in MWNTs seems to be similar to that in graphite (Em~0.1 eV for single carbon interstitials in graphite). Further studies are required to determine Em in MWNTs more accurately. A combination of these two mechanisms gives rise to a very efficient insitu defect annealing under irradiation at high (>300 °C) temperatures [30]. Thus, CNTs have surprisingly high ability to heal the irradiation-induced damage, which should facilitate the nano-engineering of CNT-based materials. Annealing of defects is also the driving force for irradiation-mediated CNT welding and coalescence. Electron beam nano-welding Direct joining by using TEM It was reported that crossing single-walled carbon nanotubes were joined by electron beam welding to form molecular junctions [21]. With this method, stable joining of various geometries can be created in-situ in a transmission electron microscope. Before the joining process, SWNTs [40, 41] were dispersed ultrasonically in ethanol and deposited onto holey carbon grids for transmission electron microscopy (TEM) observations. The joining process was carried out in a high-voltage TEM with accelerating voltage of 1.25 MV at 800 °C specimen temperature. Electron beam exposure at high temperatures induces structural defects which promote the joining of tubes via crosslinking of dangling bonds. The tubes were welded together under the influence of electron irradiation and annealing at their contact region. As depicted schematically in Fig. 18.3, during the formation of the present junction, from the random criss-crossing distribution of individual nanotubes and nanotube bundles on the specimen grid, several contact points could be identified where tubes were crossing and ‘touching’ each other. After a few minutes of irradiating two crossing tubes, their merging was observed at the point of contact, resulting in the formation of a junction with an ‘X’ shape. Since the merging of crossing tubes did not occur in the absence of irradiation, it is concluded that electron beam effects are responsible for the formation of the junctions. It is well known that knock-on displacements of carbon atoms, i.e., the formation of vacancies and interstitials, induce rearrangements within graphitic structures under high-energy particle irradiation [32]. It is important to emphasize that electron irradiation at room temperature
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18.3 (a) A SWNT of 2.0 nm diameter crossing with an individual SWNT of 0.9 nm diameter (b) 60 sec of electron irradiation promotes a molecular connection between the thin and the wide tube, forming an ‘X’ junction. Schematics show that this junction is twisted out of the plane. Molecular models of each image are provided [21].
would rapidly lead to heavy radiation damage of the tubes [33]. Thus, the presence of irradiation induced vacancies within the tubes is also responsible for the formation of junctions [21]. Dangling bonds around vacancies at the point of contact of the two tubes can serve as bridges for the merging process. Rearrangement of carbon atoms occurs so as to form heptagonal or octagonal rings at the common surfaces, thus introducing negative curvature where tubes intersect as shown in Figs 18.3–18.4 [42–47]. At the high specimen temperature, carbon interstitials are highly mobile, leading to the annealing of vacancy-interstitial pairs before interstitial agglomerates can form. The ready-formed X junctions can be manipulated in order to create Yand T-like molecular connections (Figs 18.4–18.5). It has been established from earlier work that continuous sputtering of carbon atoms from the nanotube body takes place during irradiation, leading to dimensional changes and surface reconstructions [20]. Using careful conditions of irradiation, one of the ‘arms’ of an X junction was removed in order to create a Y or T junction. This indicates that controlled electron irradiation can tailor the transformation of the junction geometry.
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18.4 High-resolution transmission electron microscopy image and molecular model of a Y junction created following electron irradiation of an ‘X’ structure. One of the arms of the ‘X’ junction vanished due to continuous sputtering under the electron beam, and a three-terminal junction remained. The junction exhibits tubes of different diameters which are molecularly joined [21].
2 nm (a)
(b)
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18.5 (a)–(c) HRTEM images of a ‘T’-like’ junction formed after irradiating a preformed Y junction. This junction may not be a perfect ‘T’ (90° angle between the crossing tubes) because we observe the projection of the object. The sequence shows the motion of this junction and the rotation by 1800 under the electron beam; note the circular cross section in one of the tubes (b). Atomic models of the junctions are depicted below [21].
Direct joining by using SEM The method of generating micro-to-nano-scale (50 to 900 nm) welds using a scanning electron microscope’s (SEM) electron beam was reported by Neely et al. [25]. According to Neely et al. the energies involved in beam-
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to-specimen interactions were high enough to melt or drill into metallic samples of small (micro to nano) dimension. Micro-to-nano adapted electron beam welding (MAEBW) also mitigates the amount of contamination (which can be especially ruinous at this small scale) found with other processes as the welding takes place within a vacuum atmosphere. The electron beam of the SEM can be focused to spot sizes in the nano range (approximately 50– 125 nm) and can melt a variety of materials with melt areas ranging from 100 nm to 50 µm. The ability to focus the beam at high magnifications on the faying surface affords reliability and an unparalleled control at this scale. The electron beam can be controlled such that the beam does not raster a large area but traces a pattern on the faying surface to be joined. This was accomplished by decreasing the raster area (increasing the magnification) and by controlling the beam deflection lens currents (raster signal) such that custom signal generation can direct the beam to trace. The amplitude of the scan area was defined to produce welds commensurate with the part size. Successful MAEBW of such materials as polysilicon, nickel, Alumel, Chromel, and Tophet C was accomplished with proper beam control relative to part size where sample sizes range from 500 nm to 50 µm. Resultant welds were crack free and are not porous. The welds have not shown the spiking or ripple formation typical of conventional macro scale electron beam welding; however, these are not deep partial penetration fusion zones. Heat input was also low enough that distortion of the small parts is avoided and the heataffected zone is minimized. MAEBW can be successfully used to join many materials of varying geometries (MEMS and LIGA components). As a process that provides a clean, consistent method for joining micro-to-nano scale materials, MAEBW is one of the promising new nano-joining tools [25]. Indirect nanojoining by SEM and TEM Modern SEMs with field emitters have a capacity to focus the electron beam onto a spot of less than 1 nm in diameter. These instruments enable us to image the shape of CNTs with almost subnanometer resolution and to manipulate surfaces by depositing contamination selectively. In such a way, a pellet of contamination can be deposited on the junction of two crossing tubes [15]. This is actually a technique of soldering because the tubes do not coalesce; instead, they are held together by a small aggregate of amorphous carbon. Madsen et al. presented an in-situ method for the highly conductive attachment of nanoscale components by the use of a gold-carbon composite soldering material deposited by a focused electron beam. This method does not require electrical contact to the electrodes or the component and allows for the assembly of 3D structures [24, 26]. They used a Philips XL30 ESEMFEG environmental scanning electron microscope, operating at a water vapour
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pressure of 100 Pa. Dimethylacetylacetonate gold (III), which has a vapour pressure of 1 Pa at 25 °C, was placed in a container with a narrow bore tube to control the diffusion of organometallic vapour onto the sample. The electron beam locally decomposed the organometallic compound and thereby deposited a material with metallic content [48]. Using a 2-mm long tube with a diameter of 0.8 mm, they obtained a growth rate of 500 nm/min. Tips with lengths of more than 10 µm could be grown without a significant decrease in the growth rate. All depositions were made at room temperature. A nanomanipulator stage inside the chamber was used to move a silicon chip with two cantilever microelectrodes [49]. The electrodes were connected to a DC voltage source, and the current was monitored continuously (see Fig. 18.6(a)). Samples of free-standing multiwalled carbon nanotubes (MWNTs) were prepared and characterized by transmission electron microscopy (TEM), which revealed the presence of more than 20-µm-long, 80–120-nm-wide MWNTs [50]. In the ESEM, a microelectrode pair was aligned to a nanotube extending from the sample so that both electrodes touched the nanotube. By slowly scanning the beam across the nanotube at the point of contact to the first electrode, two cross-shaped gold-carbon soldering bonds were formed. It was reported by Madsen et al. that the deposition of a set of protective bonds near the edge of the microelectrodes allowed mechanical breaking off of the MWNT parts extending beyond the electrodes without damaging the
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18.6 (a) Illustration of two microelectrodes positioned close to a multiwalled nanotube (1) extending from a catalyst particle on a substrate. Organometallic molecules decomposed by the electron beam (2) are deposited to form a cross-shaped solder bond (3) and a protective bond (4) near the edge of the electrode. (b) ESEM image of a carbon nanotube across two electrodes, connected by soldering bonds and protective bonds (top). When the electrode pair is withdrawn, the nanotube breaks at the protective bonds (bottom) [24].
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soldering bonds (see Fig. 18.6(b)), lower panel) [24]. They consistently observed the MWNTs to break rather than the protective bonds, both for tube-electrode bonds (Fig. 18.6) and for tube-tube bonds, as shown in Fig. 18.7. To avoid unwanted contamination of the nanotubes by soldering material, they avoided imaging the suspended part of the nanotubes at high magnification. For all of the examined bridges, electrical contact was established during the soldering procedure at the second electrode. In the example shown in Fig. 18.8, the current increased in two steps, first from 0 to 60 nA and then abruptly to 300 nA, at a bias voltage of 10 mV. A linear current-voltage curve was measured, indicating metallic conduction (Fig. 18.8, inset), and the resistance was 27 kΩ. Madsen et al. connected four nanotubes and every time obtained reliable ohmic contacts upon soldering to the second electrode, with resistances of 9, 11, 27, and 29 kΩ, with no clear correlation to the length of the MWNT bridge [26]. The resistances of the nanotube bridges were unaffected by the breaking of the nanotube extensions and by the deposition of the protective bonds and were found to be constant in air for days. Resistivity of the soldering material slightly larger than the value of 1.3 × 10–5 Ω cm had been reported by Bietsch et al. [51] for micro-contact printed pure gold nanowires of similar dimensions and two orders of magnitude larger than that of bulk gold. For electron beam deposited nanowires, resistances as small as theirs had been obtained by heating the sample to 80 °C during deposition to increase the relative content of gold [48]. These values were
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18.7 SEM image sequence of tube-to-tube soldering showing the approach of a nanotube soldered to a microelectrode toward a MWNT extending from the substrate (a) and subsequent soldering of the tube ends (b). Withdrawal of the electrode broke the nanotube and not the bond (c) [26].
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18.8 Current vs time during soldering of the MWNT using the electron beam. A fixed voltage of 10 mV is applied across the microelectrodes. During the formation of the second soldering bond, the current suddenly increased to a stable 300 nA. The IV characteristic (inset) was found to be linear [24].
obtained only after annealing at 180 °C, which further reduced the resistivity by 2–3 orders of magnitude. According to Madsen et al., one possible explanation for the high conductivity of their material achieved at room temperature without annealing could be that the presence of H2O in the sample chamber reduces the relative amount of carbon, as had been suggested by Folch and co-workers [52]. This could be clarified by analyzing the chemical composition of the soldering material deposited at different vapour pressures. Madsen et al. also attached MWNTs to microelectrodes by means of nonmetallic carbonaceous material; these devices showed electrical conduction in the mega-ohm range [24]. This strongly indicated that the metal content of the soldering material is necessary for good electrical contact. The soldering bonds were found to be mechanically strong compared to the MWNTs. They anticipate automated electron beam nano-soldering to be useful for quickly connecting complex circuitry consisting of nanoscale components in a way similar to the soldering of electronic components on the macroscale. Banhart et al. joined CNTs mechanically by forming an aggregate of amorphous carbon also using a beam of an SEM [15]. Hydrocarbon molecules from the air and other environments are attracted by CNTs as they are by most other specimens. At room temperature, these molecules are highly mobile on almost all surfaces. Once such molecules migrate into the area that is irradiated with an electron beam of sufficient energy, dissociation under the beam leads to transformation into amorphous carbon, which is
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immobile and remains in the irradiated area [53, 54]. Thus, sustained irradiation of a slightly contaminated surface leads to the accumulation of amorphous carbon. The electrical conductivity of such a deposit has not been studied in detail; however, amorphous and graphitic carbon are known to be conductive, whereas hydrocarbons are insulators. The conductivity of the deposit depends therefore on the completeness of the transformation of hydrocarbons to amorphous or graphitic carbon. Multiwalled carbon nanotubes were obtained from the deposit on the cathode of an arc-discharge apparatus [15]. The deposit containing nanotubes was prepared on metal support grids for transmission electron microscopy (TEM). These specimens were studied and irradiated in an SEM (Hitachi S-5200) that was equipped with a field emission gun and an in-lens specimen stage for ultimate resolution. It was found that the amount of contamination depends on the period during which the specimen was exposed to air; whereas newly prepared specimens show almost no contamination even after longer irradiation [15]. Nanotubes in old specimens are covered by a layer of hydrocarbons that leads to rapid accumulation of contamination in the electron beam. Apparently, CNTs strongly attract hydrocarbon molecules from the air. When two crossing nanotubes were found in the specimen material, the beam of the SEM was focused onto the point of contact until the contamination deposit had reached a measurable size. Figure 18.9 shows junctions before (a) and after (b) the deposition. The soldering effect of the deposit can be clearly seen. Banhart utilized a higher SEM voltage (20–30 kV) to promote the transformation of hydrocarbons into amorphous or graphitic carbon already in the aggregation. To improve the electrical conductivity of the deposited material, graphitization of the initially amorphous deposit was attempted,
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18.9 SEM images of a nanotube junction before (a) and after (b) soldering by deposition of amorphous carbon. Due to a high contamination rate, the deposit at the junction is large. Some contamination is already visible in (a) due to inevitable irradiation during recording of the SEM image (droplets on the tubes) [15].
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either by irradiation with an electron beam of higher energy in a TEM or by annealing in a furnace at high temperature. After the generation of junctions in the SEM, some specimens were transferred to a TEM where the junctions were irradiated with electrons of 80 keV energy for several minutes. As an alternative approach to graphitization, some other specimens were annealed in a furnace in high vacuum (2 × 10–7mbar) at 700 °C for about 1 h [15]. Figure 18.10 shows a TEM image of a junction of two nanotubes before and after deposition of contamination in the SEM. The deposit shows amorphous carbon and some disordered graphitic basal layers. Figure 18.10(b)) shows the result of irradiation with an 80 keV electron beam in the TEM.
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18.10 Transmission electron microscopy images of a nanotube junction with a carbon deposit before (a) and after (b) irradiation induced graphitization. The deposit connects the tubes on the righthand side of the junction [15].
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The graphitization was improved; however, some damage of the tubes appears due to inevitable irradiation with 200 keV electrons during imaging in the TEM. Figure 18.11 shows another junction after thermal annealing. As a result, irradiation with an electron beam of appropriate energy seems to be more effective in graphitization than does annealing at moderate temperatures. Ion beam nano-welding Fabrication of amorphous nanowire junctions from crossed multiwalled carbon nanotubes (MWNT) using carbon ion-beam irradiation was previously reported by Wang et al. [19]. The multiwall carbon nanotubes used in this research were synthesized by chemical vapour deposition (CVD) by catalyst iron particles. Without the condition of in-situ observation for ion beam irradiation, the experiments were performed in three stages. Firstly, purified and un-irradiated nanotubes were transferred into the holey carbon microgrid for TEM (Philips CM200-FEG-EM430 mould) examination with an
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18.11 TEM image of a T-junction of two nanotubes after deposition of amorphous carbon and annealing at 7000°C (the black dots are metal crystals that have condensed on the sample during the annealing procedure in the furnace) [15].
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accelerating voltage of 160 kV, and the images of un-irradiated nanotubes were obtained. Secondly, after TEM analysis, this holey carbon micro-grid filled with un-irradiated nanotubes was manipulated that it could be irradiated perpendicularly by carbon ion beam with an energy of 50 keV in the vacuum chamber of Electromagnetic Isotope Separator at a pressure of about 10–3 Pa. The current and dose of carbon ion beam were about 10 µA and 1 × 1017 cm–2, respectively. The focused area of the incident ion beam was about 10 mm2 larger than that of the micro-grid, which guaranteed that all the un-irradiated nanotubes in micro-grid were irradiated by carbon ions. Finally, these irradiated MWNTs were analyzed by TEM and electron diffraction (ED) again. An electron micrograph of the un-irradiated nanotubes on the carbon micro-grid is shown in Fig. 18.12(a). This structure not only caused the graphite layers to appear in HRTEM, but also the ED ring pattern as in Fig. 18.12(b) inset. Figure 18.12(b) represents a typical high-resolution transmission electron microscopy (HRTEM) image and electron diffraction pattern (EDP). After carbon ion irradiation, the nanotubes evolved into amorphous solid carbon nanowires (see Fig. 18.13). The amorphous nanowires have uniform diameters and a smooth surface along their length. It is also seen that the nanowires have a circular solid cross-section with short-range-order graphitic lattice lines around the amorphous nanowires surface region (as shown in Fig. 18.13(b)), indicating a common characteristic feature of amorphous nanowires [55]. The blurred graphitic and amorphous halo rings in selective area electron diffraction pattern (SAEDP) (inset in Fig. 18.13(b)) of a nanowire demonstrate that the nanowire consists of both amorphous solid and shortrange graphene sheets.
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18.12 (a) TEM image of carbon nanotubes before ion irradiation; (b) HRTEM image of a CVD carbon nanotube [19].
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18.13 (a) TEM image of carbon nanowires obtained by ion irradiation on CVD nanotubes; (b) cross-section image of amorphous carbon nanowires. The inset is a SAEDP recorded perpendicular to the long axis of nanowires [19].
The most interesting observation in these experiments is the web-like nanowires created by ion irradiation on MWNTs (Fig. 18.14). Several stable amorphous nanowire junctions of various geometries (Fig. 18.15) could be found in the web-like nanowires. The contact points or regions could be identified where nanowires were crossing and touching each other. As depicted in Fig. 18.15, welded junctions with ‘T’ and ‘Y’ shape have uniform amorphous microstructures and a smooth surface. This was the first experimental observation of such amorphous nanowire junctions via ion-beam irradiation, although branched tubular carbon fibres had been synthesized earlier [56– 58]. The ion irradiation method is very suitable for production of novel type junction-like structures consisting of amorphous nanowires with uniform diameter of trunk and branches. Whereas it has so far been believed that the synthesis of connections between two or more different nanotubes and nanowires was an important step in the nano-electronic and photonic devices and circuits. The present method, ion-beam weld of crossed multiwall carbon nanotubes, was used to provide a powerful way in which to achieve perfect welding to build various junction-like structures. This ion beam nano-weld method may be helpful for preparing other systems or aligned arrays on substrates by micro-control for potential device applications. It should be noted that irradiation effects in carbon nanotubes are different from those that occur in regular crystals. Generally, large curvature of graphene layers could make linear displacement collision sequences in them very few, but displacement cascades in nanotubes are possible if the structures are
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18.14 TEM image of web-like carbon nanowires with fairly straight shapes an uniform diameters [19].
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18.15 HRTEM images of amorphous nanowire junctions with ‘T’ (a) and ‘Y’ shape [19].
large enough and built up of several walls of carbon nanotubes. It has been found that multiwall carbon nanowires are comparable with graphite under ion-irradiation [30]. The single-vacancy-contained region in them would transform into a disordered region gradually, of which accumulation leads to amorphization [59]. This kind of irradiation-induced amorphization of CVD nanotubes has already been found in experiments [60, 61].
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18.2.2 Joining by resistance nano-welding Direct resistance nano-welding Cox et al. presented a technique to select, manipulate, and produce nanotube joints with good mechanical and electrical integrity, without the need for a solder material, and with significantly reduced time taken per connected joint [62]. The technique was based on controlled heating of the joints formed between nanotubes and substrates, and relies on careful control of applied voltages and subsequent current induced heating. Arc-discharge-produced nanotubes of high crystalline quality were loaded in a scanning electron microscope (SEM) that includes electrical feedthroughs connected to a pair of sharp tungsten tips and the specimen substrates. The tips can be moved independently of the SEM stage and each other by three-axis piezo sliders, with a minimum step size of 20 nm and up to 5 mm total travel (Fig. 18.16(a)). A lab-view program was used to run two Keithley 238 source/current meters connected to the tips and the SEM stage that provide a means of electrically characterizing, cutting, and welding individual nanotubes. To select a nanotube for the manipulation process, one of the tips was brought close to a single nanotube on its substrate, and the gap reduced until contact was made. Often, electrostatic attraction pulls the nanotube onto the end of the tip, relieving the need for accurate positioning. A voltage was then applied across the nanotube to establish an electrical contact. Initially, a poor contact was achieved and the measured current rarely exceeded a few nanoamperes, until a threshold voltage (~3 V) was reached (Fig. 18.16(b)). As the voltage was raised further, the current then rapidly increased by several orders of magnitude up to a predetermined limit of 1 × 10–5 A. Subsequent cycling of the voltage over a typical range 0 to 5 V with a current limit still imposed resulted in a lowering of resistance with each successive cycle. After typically ten cycles, completed in under a minute, stable I-V behaviour was obtained, with no further conditioning. As can clearly be seen in Fig. 18.16(b), initially, the current limit of 1 × 10–5 A was observed at around 4.5 V, but following cycling, the same current was achieved at lower than 2.5 V. The tip-to-nanotube contact obtained was mechanically robust and electrically active. According to Cox et al. this conditioning behaviour was due to ohmic heating taking place at the connection, caused by the large current densities, facilitating alloying at the connection. Cutting nanotubes to a predetermined length was by utilizing the second tip [62]. A nanotube was connected between the first (primary) tip and the substrate, using the welding technique, and then a voltage of typically 1–2 V was applied. On bringing the second (cutting) tip, with a similar voltage, into contact with the nanotube, at the point cutting was required, and then raising the voltage on the cutting tip
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(a) × 106 Initial cycle Second cycle Tenth cycle
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18.16 Secondary electron micrograph of the tungsten tips and substrate in the SEM (scale bar is 10 µm) (a), I-V curves for the nanotube-tip-substrate connection (b) [62].
while limiting the current that can flow into the primary tip to zero, all the current was caused to flow in the lower portion of the nanotube attached to the substrate. This results in a small region of around 100–150 nm vapourizing near to the point of contact of the cutting tip, leaving the desired length of nanotube attached to the first tip. Following the cutting procedure it becomes possible to move the nanotube attached to the tip to a new substrate and connect it as previously shown. It is also possible to connect additional nanotubes to the previously deposited ones. As with attaching to substrates, the nanotubes were brought into contact
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at the desired point on the previously deposited nanotube and cut off from the tip. As a demonstration of the versatility of this new technique, Figure 18.17 shows outlines that depict the letters A, T, and I constructed from carbon nanotubes and deposited onto the edge of a chromium foil. The height of the three letters is approximately 1.5 µm, occupying a length of less than 4 µm. The figure demonstrates the most important requirements necessary for the production of three-dimensional nanotube structures; the ability to pick and place a single nanotube onto a substrate with a specific length (letter ‘I’), the ability to attach a second nanotube to the first (letter ‘T’), and the ability to connect nanotubes together that have previously been attached to a chosen substrate (letter ‘A’). All of the connections are mechanically strong as well as electrically active. According to Cox et al. it is possible to weld nanotubes to metal contacts and to each other with both good electrical and mechanical quality. The (a)
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18.17 Nanotube construction of letters A, T, and I using six nanotubes. Secondary electron micrograph of nanotube construction (scale bar is 300 nm) (a), and its schematic diagram showing order in which the nanotubes were attached and the connections formed (b). The nanotubes were placed in the order shown numerically (1–6) and the connections made in the alphabetical order also shown in the schematic diagram [62].
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procedure is fast enough to produce relatively complex structures and uses only the number of nanotubes required. They also reported that in electrical conduction through the tubes, heat dissipation is crucial if high current density devices are to be fabricated. Any structural or mechanical defect compromises the integrity of the tube, particularly at current densities in excess of 2 × 106 A/cm2. This procedure may be utilized as a technique to construct carbon nanotube-based devices and connections in electronic applications of higher complexity than previously available by other techniques. Indirect resistance nano-welding Dong et al. [63] show that carbon nanotubes filled with copper metal can act as nanoscale spot welders. When electrically heated, the encapsulated metal melts and flows out, being delivered in a highly controllable manner with nanometer positional precision. Dong et al. used this method to weld two nanotubes together. They found that, because the melting point of copper is rather modest when encapsulated (approximately 700 °C) [63], it might be possible to use copper-filled nanotubes as nano-pipettes for delivering molten metal to well-defined locations. They positioned a 50-nanometre-wide carbon nanotube filled with copper inside a nanorobotic manipulator, and ran a small voltage through it to melt the copper. In experiments the researchers positioned the manipulator so the melting metal connected one carbon nanotube to another. The researchers pasted a bundle of single-walled carbon nano tubes SWCNT onto a gold wire with the needle-like tips pointing outwards. They then used a tungsten probe tip, 200 nm across at the end, to make contact with an individual nanotube, while observing them through a field emission electron microscope (FESEM) (Fig. 18.18). They found that by applying a voltage between the probe and a nanotube while the two were in contact, the temperature could be raised past the melting point of copper by resistive heating. This started to happen at a voltage of 1.5 V. At that point, they saw a small vacant region, like a little bubble, open up in the copper held within the first conical bamboo section of the nanotube. This empty space was caused by metal flowing out of the tube end onto the probe. The bubble moved to the top end of the bamboo section, after which the meniscus of molten copper within the tube gradually descended towards the tip as the copper flowed out. They also detached the nanotube from the probe, with a small droplet of metal still stuck to the end. An estimated 5.5 femto-grams (5.5 × 10–15 g) diffused onto the probe due to the capillary action. They calculated that the flow rate at a voltage of 2.5 V was 120 atto-grams (120 × 10–18 g) per second, allowing for very precise control of the amount of material delivered from the tube.
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18.18 (a–d) Copper-filled CNTs. (a) FESEM images of Cu-filled CNTs. Observation shows that all the CNTs have sharp tips filled with metal nano-needles. These CNTs are up to 5 µm long, with outer diameters in a range of 40–80 nm. (b) A typical copper-filled CNT synthesized for 30 min. (c) HRTEM image reveals that the Cu nano-needles are encapsulated in graphite walls approximately 4–6 nm thick. The inset is the corresponding SAED pattern of the Cu nano-needle along the [112] zone axis, showing that the Cu nano-needle is single crystalline. The appearance of a pair of arcs in the SAED pattern indicates some orientation of the (002) planes in the carbon tubes. (d) The magnified image of the rectangular region in (c). The large arrows indicate the growth direction of the CNTs. The interlayer spacing of carbon nanotube is about 0.34 nm, consistent with the (002) plane lattice parameter of graphite. It can also be seen that the graphite layers are not parallel to the tube axis. (e, f) Nanorobotic manipulation system in a TEM. (e) ST1000 STM-TEM holder (Nanofactory Instruments AB). (f) Schematic setup [63].
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18.2.3 Joining by ultrasonic nano-welding A simple ultrasonic nano-welding process was developed to fabricate reliable bonding between single-wall carbon nanotubes (SWCNTs) and metal electrodes by Chen et al. [64]. Contacts formed by the present process were found to have low contact resistance, good long-term stability and mechanical strength. Single-wall carbon nanotubes (SWCNTs) with an average diameter of 1.6 nm synthesized by an arc discharge process were used in that work. SWCNTs were purified and subjected to chain-scission in a 3:1 mixture of concentrated sulfuric acid and nitric acid at 80 °C for 30 min in a reflux system. The undried SWCNTs were filtrated with isopropanol to remove the water in the sample. Then the SWCNTs were ultrasonicated in isopropanol and centrifuged at 15 000 rpm for 2.5 h to further separate large SWCNT bundles from single SWCNTs. The obtained supernatants at concentrations of several µg ml–1 were ultrasonically treated for about 20 h to sufficiently disperse SWCNTs. A droplet of the SWNT suspension was introduced onto the patterned wafers. After the deposition or electric field alignment process, SWNTs were laid over the electrode pair. The number of SWNTs bridging the electrode pair can be readily controlled by varying the concentration of the suspension or the parameters of applied AC electric field such as the bias magnitude and frequency. Ultrasonic nano-welding was carried out in an FB-128 ultrasonic wire bonder. An Al2O3 single crystal with a 50 µm2 pressing surface and an rms roughness of 0.2 nm was mounted onto the bonder to act as the welding head. Figure 18.19 shows a schematic diagram of the ultrasonic nano-welding process. A clamping force of 78.4 mN was applied to press the welding head against the nanotube and electrodes. At the same time an ultrasonic vibration with a frequency of 60 kHz was applied to the welding head through an ultrasonic transducer (Fig. 18.19(a)). The ultrasonic energy was transferred to the bonding interface through the ultrasonic welding head. Thus the ends of SWCNTs and electrodes were welded together under the combined action of the ultrasonic energy and a clamping force (Fig. 18.19(b)). To investigate the effects of ultrasonic energy on the bonding process, a series of different ultrasonic powers was applied. The welding process was carried out at room temperature for 0.2 s. Figures 18.20(a) and (b) show typical scanning electron microscopy (SEM) images of an SWCNT bridging two electrodes before and after welding with an ultrasonic power of 0.07 W. It can be observed that an SWCNT was hanging over the two electrodes before welding. The morphology of the nanotube on the two electrodes can be clearly observed. After welding, the ends of the SWCNT were embedded into the electrodes and the nanotube morphology was now almost invisible on the electrodes. Scratches caused by the friction between the ultrasonic welding head and the electrodes can
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18.19 Schematic diagrams of the ultrasonic nano-welding process: (a) before welding; (b) after welding [64].
also be observed on the electrode surfaces (Fig. 18.20(b)). Figure 18.20(c) shows that multiple nanotubes were welded onto the titanium electrodes under pressing as well as ultrasonic treatment [64]. Ultrasonic nano-welding can also be used for welding bulk quantities of SWCNTs onto metal electrodes. Figure 18.21 shows SEM images of a 50 µm2 electrode surface with nano tubes welded onto it. As shown in Fig. 18.21(a), the ultrasonic nano-welding produces an apparent directional welding zone due to the directional vibration of the welding head. In the non-welded zone, the SWCNTs loosely stay on the surface of the metal electrode (Fig. 18.21(b)). Inside the welding zone, the SWCNTs and metal substrate are bonded together to form a new weld surface. Almost all SWCNTs are embedded and welded into the Ti electrode and only the ends of some short SWCNTs protruding from the welded substrate can be observed (see Fig. 18.21(c)). In short, during ultrasonic nano-welding, the high-frequency ultrasonic energy softens the metal and causes plastic deformation of the metal under the clamping stress because of the acoustic softening effect [65, 66]. Thus, the nano-sized SWNTs with a one-dimensional structure have been shown to be ‘embeddable’ and weldable into the metal electrodes. To examine the nature of the welded bonds, micro-area x-ray photoelectron spectroscopy
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18.20 (a) and (b) SEM images of an individual SWCNT bridging the Ti electrodes before and after the ultrasonic nano-welding with an ultrasonic power of 0.07 W, respectively. (c) SEM SWCNTs welded onto the electrode with an ultrasonic power of 0.16 W; the marked regions show the weld junctions [64].
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18.21 SEM images of the electrode surface after welding of bulk quantities of SWCNTs onto a Ti electrode at an ultrasonic power of 0.21 W. (a) Surface topography of the welded (upside) and nonwelded (downside) zone. The arrow represents the direction of the ultrasonic vibration motion. ((b), (c)) Zoom-in views of a local region (indicated by rectangles) in the un-welded and welded zone, respectively. The arrows in (c) show that some short ends of SWCNTs protrude from the welded substrate [64].
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(XPS) and x-ray diffraction (XRD) analysis were used to study the welded zone. All these results indicate that no detectable chemical reaction occurs between the nanotube and the metal electrode at low ultrasonic power. When the ultrasonic power is higher than a threshold value (>0.14 W) SWCNTs do chemically react with the titanium electrode to form titanium carbide at the welding junction. More comprehensive and in-depth study is needed to substantiate this observation. It was reported that to evaluate the electrical performance of the contact, two-terminal (2t) resistance tests were carried out on the nanotube welded across the electrode pairs [59]. Before the ultrasonic nano-welding, the 2tresistances of individual metallic SWCNTs were of the order of tens of mega-ohms at room temperature. After ultrasonic nano-welding with a power of 0.16–0.19 W, the 2t-resistances decreased to a narrow range of 8–24 kΩ. The lowest resistance obtained was as low as 7.94 kΩ, which approaches the theoretical minimum resistance (6.45 kΩ) of a ballistic conducting metallic SWCNT with perfect contacts, although the resistance was measured at room temperature. Besides the electrical tests, the mechanical integrity of the welded bonds has also been tested by mechanically pushing the hanging middle segment of a welded nanotube with an AFM tip. It was found that when SWCNTs were not welded to the electrodes, SWCNTs were easily shifted and could no longer bridge both electrodes (Figs 18.22(a) and (b)) [64]. This observation qualitatively shows that the van der Waals’ forces between the CNTs and substrate underneath are not large enough to resist the movement of the CNTs. In contrast, for the nanowelded samples, the two ends of the 95% SWCNTs remained bonded to the electrodes even after the nanotube had been pushed to fracture (Fig. 18.22(c)). This suggests that the bonds bear reasonable mechanical strength. According to Chen et al., this technique does not depend on the specific kind of nano-component or metal electrodes. Moreover, it is hopeful that the ultrasonic welding area can be scaled up so that multiple bondings on a substrate can be achieved in one single step, under one single press vibrating at a specific ultrasonic frequency.
18.2.4 Joining by laser nano-welding Nano-particles have a variety of unique spectroscopic, electronic, and chemical properties that originate from their small sizes and high surface/volume ratios [3–6]. The assembly of nano-particles into a symmetrically and spatially well-defined organization is important to achieve properties that are desired for nano-devices [67–70]. Whereas metallic nano-particles do not have any ohmic contact, their networks have mutual contact. To construct the nanodevices from the ‘parts’, it is necessary to interconnect the parts together to have ohmic nano-contact.
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18.22 Characterizations of the mechanical stability of the ultrasonic nanowelding by AFM manipulation. (a) Three-dimensional AFM image of an un-welded SWCNT bridging the electrodes. (b) AFM image of the SWCNT in (a) shifted laterally by the AFM tip. (c) AFM image of a nano-welded SWCNT pushed to fracture by the AFM tip [64].
Kim and Jang showed that gold nano-particles can be welded on a carboncoated copper TEM grid in a programmed manner to have ohmic nanoconjunction using a pico-second laser. They considered that their method of laser-induced nano-conjunction would be greatly beneficial to the fabrication and manipulation of nanostructured materials and devices [71]. Pico-second laser pulses were employed to adjoin, to hold closely, and to weld gold nanoparticles on carbon-coated copper grids [71]. They recruited pico-second laser pulses to adjoin, to hold closely, and to weld gold nano-particles present in applied drops of aqueous gold colloids on carbon-coated copper grids of transmission electron microscopy (TEM). They showed that the laser-induced nano-welding of metallic nano-materials can be achieved in a controlled manner to yield ohmic nano-contact having a single phase by considering
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the size-dependent thermodynamic, dynamic, and kinetic properties of nanostructured materials. According to the report of Kim and Jang, surfactant-free gold nano-spheres of 9.3 nM dispersed in water, having an average diameter of 13 nm and absorption maximum of the surface-plasmon resonances at 520 nm, were prepared by the citrate reduction of HAuCl4. Gold nanoparticles were applied to carbon-coated copper TEM grids by dropping many small drops of aqueous gold colloids slowly and successively using a micro-pipet. TEM grids loaded with gold nanoparticles were irradiated to weld by using the second harmonic pulses of 532 nm having a duration of 30 ps from a mode-locked Nd:YAG laser (Quantel, YG701) run at 10 Hz. The spot diameter of the irradiation beam was 4 mm. Energy-dispersive x-ray (EDX) profiles and high-resolution transmission electron microscopy HRTEMd images were measured with a JEOL JEM-3000F and TEM images with a JEOL JEM-2000. Figure 18.23 shows that gold nanoparticles were welded together in a programmed manner using a pico-second laser. Kim and Jang had found from magnified HRTEM images that gold nanoparticles were well contacted in a single phase to show the typical crystalline fcc structure of gold. Laser pulses had been utilized to join, to
20 nm
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18.23 HRTEM images of gold nanoparticles held together and nanowelded by irradiating laser pulses of 532 nm and 0.2 mJ for 10 min on a carbon-coated copper TEM grid, indicating that gold nanoparticles are welded in a single phase to have ohmic nanocontact [71].
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hold closely, and to weld gold nanoparticles on TEM grids. They tried to find the optimum conditions for the nano-welding of gold nanoparticles on carboncoated copper TEM grids. Nanoparticles should be adjoined and held closely during laser welding. They had found that laser irradiation induces the association of gold nanoparticles on a wet TEM grid to be withheld in contact during welding as well as the nano-welding of associated nanoparticles (Figs 18.24(a) and (b)). Nano-welding was found to proceed linearly with irradiation time as expected. On the other hand, the control of irradiation power was found to be the most critical for optimum nano-welding because welding as well as melting proceeds nonlinearly fast with power increase (Figs 18.24(a) and (c)) [72].
(a)
20 nm (b)
(c)
18.24 TEM images of gold nanoparticles irradiated with 532 nm laser pulses of 0.2 sad, sbd and 0.8 mJ scd for 10 min. The nanoparticles were irradiated right after sad, scd and a day after sbd applying aqueous gold colloids on carbon-coated copper TEM grids. The comparison of sad and (b) suggests that laser irradiation, especially on the wet grid, induces the association of gold nanoparticles to be held closely during welding. On the other hand, the comparison of (a) and (c) indicates that the control of irradiation power is very critical to achieve optimal nano-welding [71].
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Summary and future trends
Carbon nanotubes (CNTs) have great potential as nanoscale building blocks for future nanoelectronics due to their unique one-dimensional nanostructure and properties [73]. To explore their potential in various domains, an essential prerequisite is to build reliable interconnections between the CNTs and the external electrical circuits and mechanical systems. To address this need, various chemical and physical processes have been explored to build such interconnections. For example, Burghard et al. reported controlled adsorption of CNTs on chemically modified electrodes for interconnection of CNTs [74, 75]. However, a stronger bonding instead of a weak chemical adsorption is mandatory for constructing reliable nano-devices. Ruoff et al. showed that a focused electron beam in a scanning electron microscope (SEM) can be used to deposit a small amount of hydrocarbon contamination so as to attach nanotubes on an AFM tip [76, 77]. Such a spot welding technique has also been used for connecting CNTs and a polysilicon surface electrically and mechanically [73]. Madsen et al. presented an in-situ method for highly conductive attachment of multiwall carbon nanotubes (MWCNTs) onto microelectrodes by depositing a gold–carbon composite using a focused electron beam system [24]. Dong et al. reported that multiple joinings or junctions for nanoelectronics could be made by running voltage through a copper-filled nanotube lying across two electrodes by robotic spot welding. That would be easier and take less energy than having to pattern extra electrodes on top but only a very small number of labs in the world have access to nanorobotic manipulators. Though robust contacts can be obtained by the above methods, limited access to a focused electron beam system and the small-scale spot-treatment nature prevent their large-scale industrial applications. To meet the needs of future large-scale applications, simpler, less capital intensive and more scalable processes are highly desirable.
18.4
References
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9. Y. Cui, Q. Q. Wei, H. K. Park and C. M. Lieber, Science, 293, 1289–1292 (2001). 10. S. J. Tans, M. H. Devoret, H. J. Dai, A. Thess, R. E. Smalley, L. J. Geerligs and C. Dekker, Nature (London) 386, 474 (1997). 11. C. T. White and T. N. Todorov, Nature (London) 393, 240 (1998). 12. J. Li, Q. Ye, A. Cassell, H. T. Ng, R. Stevens, J. Han and M. Meyyappan, Appl. Phys. Lett., 82, 2491 (2003). 13. R. Martel, Nat. Mater. 1, 203 (2003). 14. R. H. Baughman, A. A. Zakhidov and W. A. de Heer, Science, 297, 787 (2002). 15. F. Banhart, Nano Lett. 1, 329 (2001). 16. M. S. Fuhrer, J. Nygård, L. Shih, M. Forero, Y.-G. Yoon, M. S. C. Mazzoni, H. J. Choi, J. Ihm, S. G. Louie, A. Zettl and P. L. McEuen, Science, 288, 494–497 (2000). 17. J.-M. Ting, C.-C. Chang, Appl. Phys. Lett. 80, 324 (2002). 18. P. G. Collins, H. Bando, A. Zettl, Nanotechnology, 9, 153 (1998). 19. Z. Wang, L. Yu, W. Zhang, Y. Ding, Y. Li, J. Han, Z. Zhu, H. Xu, G. He, Y. Chen and G. Hu, Physics Letters, A 324, 321–325 (2004). 20. P. M. Ajayan, V. Ravikumar and J.-C. Charlier, Phys. Rev. Lett. 81, 1437 (1998). 21. M. Terrones, F. Banhart, N. Grobert, J.-C. Charlier, H. Terrones and P. M. Ajayan, Phys. Rev. Lett. 89, 075505 (2002). 22. A. V. Krasheninnikov, K. Nordlund, J. Keinonen and F. Banhart, Phys. Rev. B 66, 245403 (2002). 23. Y. H. Tang, N. Wang, Y. F. Zhang, C. S. Lee, I. Bello and S. T. Lee, Appl. Phys. Lett. 75, 29221 (1999). 24. D. Madsen, K. Nørgard, R. Mølhave, A. M. Mateiu, M. Rasmussen, C. J. Brorson, H. Jacobsen and P. Bøggild, Nano Lett. 3, 47 (2003). 25. B. Nowak-Neely, D. MacCallum and G. Knorovsky, AWC (2004). 26. D. N. Madsen, K. Mølhave, R. Mateiu and P. Bøggild, IEEE, 335 (2003). 27. A. V. Krasheninnikov and K. Nordlund, Nuclear Instruments and Methods in Physics Research B 216, 355 (2004). 28. A. V. Krasheninnikov, K. Nordlund, M. Sirvio, E. Salonen and J. Keinonen, Phys. Rev. B 63, 245405 (2001). 29. A. V. Krasheninnikov, K. Nordlund and J. Keinonen, Phys. Rev. B 65 (2002) 165423. 30. F. Banhart, Rep. Prog. Phys. 62, 1181 (1999). 31. B. W. Smith and D. E. Luzzi, J. Appl. Phys. 90, 3509 (2001). 32. C.-H. Kiang, W. A. Goddard, R. Beyers and D. S. Bethune, J. Phys. Chem. 100, 3749 (1996). 33. M. Terrones, H. Terrones, F. Banhart, J.-C. Charlier and P. M. Ajayan, Science, 288, 1226 (2000). 34. J. P. Salvetat, J. M. Bonard, N. H. Thomson, A. J. Kulik, L. Forro, W. Benoit and L. Zuppiroli, Appl. Phys. A (Mater. Sci. Process.) 69, 255 (1999). 35. C. P. Ewels, M. I. Heggie and P. R. Briddon, Chem. Phys. Lett. 351, 178 (2002). 36. Y. H. Lee, S. G. Kim, D. Tomanek, Phys. Rev. Lett. 78, 2393 (1997). 37. M. Heggie, B. R. Eggen, C. P. Ewels, P. Leary, S. Ali, G. Jungnickel, R. Jones and P. R. Briddon, Electrochem. Soc. Proc. 98, 60 (1998). 38. P. O. Lehtinen, A. S. Foster, A. Ayuela, A. Krasheninnikov, K. Nordlund and R. M. Nieminen, Phys. Rev. Lett. 91, 017202 (2003). 39. A. V. Krasheninnikov, K. Nordlund, P. O. Lehtinen, A. S. Foster, A. Ayuela and R. M. Nieminen, Phys. Rev. B 69 (2004). 40. C. Journet et al., Nature (London) 388, 756 (1997). 41. P. Nikolaev et al., Chem. Phys. Lett. 313, 91 (1999).
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19 Joining of high temperature superconductors G Z O U , Tsinghua University, P.R. China
19.1
Introduction
Among the more than 100 kinds of high temperature superconducting (HTS) materials, also called high temperature superconductors, Y-based (YBCO, or Y123) and Bi-based (BSCCO, or Bi-2223) cuprates have the most promising applications. They will be widely used in electric powder systems such as for cable, electromotor, generator, fault current limiter, flying-wheel-based energy storage devices, etc., in transportation and national defense fields such as for superconducting magnetic levitation (MagLev) vehicles, fighterclass lasers, etc., as well as in other fields such as medicine, electronics and superconducting quantum interference devices (SQUID). While HTS technology is still in the research and development stage, researchers are now demonstrating full-scale prototypes of electric powder cables, motors, transformers, and other heavy electrical gear made with HTS wires. These prototype systems waste much less energy than the existing technologies. In addition, HTS systems are usually smaller, lighter and safer, and, in some cases, more environmentally benign. In this chapter, HTS materials and their processing technologies are introduced first, followed by an explanation of the needs for developing joining techniques for HTS materials. Secondly, several joining technologies of BSCCO HTS materials are described with emphasis on the conventional diffusion bonding method and direct diffusion bonding method with high temperature pressing. Thirdly, the joining technologies, especially the soldering process, of YBCO bulks are expounded. Finally, conclusions and future trends are discussed.
19.2
Superconducting materials
The electrical resistivity of some materials vanishes at the so-called critical transition temperature (Tc). This phenomenon, known as superconductivity, was first discovered by H. Kamerlingh Onnes and his collaborators in 1911 583 WPNL2204
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and these materials are called superconducting materials or superconductors. Note that the superconductors always change from a superconducting state to normal as the current density through them exceeds a certain value called critical density (Jc). In addition, a superconductor at superconducting status is completely antimagnetic. On the other hand, a superconductor generally changes from a superconducting state to normal as long as the external magnetic field is over a certain value called the critical magnetic field (Hc).1,2,3 Up to now, more than 10,000 superconducting materials have been discovered and synthesized. According to the Tc value, they can be divided into low temperature superconducting (LTS) materials and HTS materials, which are generally used at the boiling point of liquid helium temperature of 4.2 K and boiling point of liquid nitrogen temperature of 77 K, respectively. On the other hand, according to the materials’ constitutions, superconducting materials consist of pure metals such as Hg, Pb, Nb, etc., binary alloys such as Nb-Ti, Nb-Zr, Pb-Bi, etc., superconducting compounds such as MgB2, Nb3Sn, NbN, V3Si, Nb3Ge, etc., superconducting oxides mainly including La-based, Y-based, Tl-based, Bi-based, Hg-based cuprates, and a few organic superconductors such as K3C60. In general, HTS materials are regarded to be equal to copper oxide superconductors or cuprates superconductors.2,3,4 However, some researchers also regarded the MgB2 compound as HTS material.5,6 The number of conducting CuO2 planes of typical cuprates, superconducting phases, and their Tc are summarized in Table 19.1.1,2,3,4,7,8 Table 19.1 The Tc and CuO2 planes of main cuprate superconducting phases1–4,7–8 Cuprates
Number of CuO2 layers
Tc (K)
Abbreviation
La2–χSrχCuO4 YBa2Cu3O7–χ LaBa2Cu3O7–χ NdBa2Cu3O7–χ SmBa2Cu3O7–χ TmBa2Cu3O7–χ YbBa2Cu3O7–χ ErBa2Cu3O7–χ Bi2Sr2CuO6 Bi2Sr2CaCu2O8 Bi2Sr2Ca2Cu3O10 Tl2Ba2Ca2Cu3O10 HgBa2CaCu2O8+χ HgBa2Ca2Cu3O10+χ HgBa2Ca2Cu3O10+χ
1 2 2 2 2 2 2 2 1 2 3 3 2 3 3
38 93 90 88 90 92 91 93 ~12 95 110 125 128 135 164*
LSCO YBCO LBCO NBCO SBCO TBCO YBCO ErBCO Bi-2201 Bi-2212 Bi-2223 Tl-2223 Hg-1212 Hg-1223 Hg-1223
*at high pressure of 30 GPa
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Cuprates are variations of the perovskite crystal structure type. In general, the high-Tc cuprates superconductors are basically orthorhombic, when crystal structures are sandwiched and highly anisotropic. All cuprates have one or more CuO2 planes called conducting planes, and superconductivity occurs in them. The CuO2 planes are always separated by the layers of other atoms such as Bi, O, Y, Ba, La, etc., which provide the charge carriers into the CuO2 planes, often called charge reservoirs. In the CuO2 planes, each copper ion is strongly bonded to four oxygen ions separated by a distance of approximately 1.9 Å. Up to now, more than one hundred cuprates with the CuO2 planes characteristic have been synthesized (see Table 19.1). In general, the number of CuO2 layers per unit cell in different cuprates, Ne, is dissimilar. Tc is related to the number of CuO2 planes and reaches the maximum at Ne = 3 at a fixed element-doping level. Furthermore, Tc is also significantly dependent on the other parameters such as the hole concentration in the planes, the orthorhombic lattice constant a, b or c, the buckling angle of CuO2 planes, etc. In summary, the maximum Tc value can be achieved only when all the necessary parameters have their ideal values.
19.3
Processing technologies of high temperature superconductors
So far, the shape styles of HTS materials include bulk or block, film, wire (including tape-shaped wire). In the last 20 years since Bednorz and Müller discovered the La-Ba-Cu-O cuprate superconductor (Tc ≈ 30 K) in 1986, tremendous efforts have been directed to improve the high current carrying capability of HTS materials. Efforts were directed first on the first generation (1G) HTS materials called BSCCO, and now on the YBCO-based coated tape-shaped conductors called second generation (2G) HTS materials as well as YBCO bulks. As a consequence, a variety of processes and technologies have been developed.
19.3.1 BSCCO tapes At present, the only commercially produced HTS wire is Ag- or Ag alloysheathed Bi-2223 and (Bi, Pb)-2223 (totally called Bi-2223/Ag) tape-shaped wire. The 1G HTS Bi-2223/Ag tape is fabricated by a standard powder-intube (PIT) process.6,9,10,11 A PIT process is described as follows. Firstly, a pure Ag or Ag-rich alloy (Ag-Cu, Ag-Sb, etc.) tube is filled with precursors (oxide powders) generally with a nominal composition of Bi1.8Pb0.33Sr1.87Ca2Cu3Oχ, which mainly consists of the primary phase Bi2212 and the second phases such as Bi-2201, Ca2PbO4, 3221, (Sr, Ca)14Cu24O41, CuO, etc., and, as a result, the powders packing density is ~40%. Secondly, the filled tube is drawn into round wire of 2 mm diameter and cut into 37~91
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filaments followed by repacking into a Ag or Ag-rich alloy (Ag-Cu, Ag-Sb, Ag-Mn, Ag-Li, Ag-Mg, Ag-Mg-Ni, etc.) tube, and then the repacked composite is subjected to the same drawing process to a final diameter of about 1.5 mm. Thirdly, the composite wire is rolled to a flat tape called a green one. Fourthly, the green tape is first heat-treated at 820~845 °C for 100~200 h accompanied by a cooling rate of 0.01~10 °C/min, using a flowing mixed atmosphere (approximately 8% O2-balance N2). Finally, further heat-treatments similar to the first is given after an intermediate rolling step. A final tape (reacted tape) is approximately 0.2 mm thick, 4 mm wide with different lengths. So far, thousands of meters of multi-filamentary Bi-2223/Ag tapes have been fabricated successfully, and maximum Jc and critical current (Ic) of tapes with 200 m long length have reached 145 A and 15 kA/cm–2 (77 k, 0T), respectively.12 For comparison, the Jc and Ic of a 1000 m length 61-filamentary tape have reached over 90 A and 9 kA/cm–2 (77 k, 0 T), respectively.11 Summarily, while the PIT process is scalable and increasingly mature for the commercial production of Bi-2223/Ag tapes, several crucial problems including reducing of alterative current (AC) loss and thermal conductivity, increasing mechanical strength and insulation properties, etc., must be solved for facilitating the large-scale practical engineering applications and other special project research.11,13
19.3.2 YBCO bulks Bulk HTS superconductors include BSCCO and MeBaCuO, or briefly notated MeBCO (Me = Y, Nd, Sm, Gd, Eu, Dy, etc.). MeBaCuO bulks, especially YBa2Cu3O7 (YBCO, or Y123) one with higher Jc and higher irreversibility field compared with BSCCO ones, will be widely used in practical applications, working at 77 K.14–19 These bulk superconductors can be fabricated with both the traditional solid sintering process and the melt-textured growth (MTG) one. Conventional solid sintering processing of YBCO bulk consists of the following steps: (i) preparation of precursor powders with a nominal composition of Y1.8Ba2.4Cu3.4O7–X containing the phases of Y2O3, BaCO3, CuO, using a solid-state reaction method, and then fully grinding; (ii) 2–3 times pre-sintering of the powders in air at a temperature range of 920 °C~950 °C, lower than the YBCO peritectic temperature of approximately 1008 °C~1015 °C, for several tens of hours, accompanied by intermediate grinding and mold-pressing; (iii) final sintering in air at about 965 °C for 20 h and then oxygenation in flowing oxygen at about 500 °C for about 20 h. Although solid processing is relatively simple, the Jc of bulk is usually far less than 10 kA/cm–2 and is severely degraded in the magnetic field because of high-angle grain boundaries, impurity layers such as carbonates, presence of porosity due to insufficient densification and micro-cracks, and
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other variations in chemistry of crystal structure at grain boundaries, etc.20 For improving Jc, doping treatment to YBCO matrix with powders of Ag, Pt, ZrO2, or Ca has been adopted.18–19 The MTG process for fabricating YBCO bulk was first used by Jin and his collaborators.20 In a typical MTG processing method, YBCO precursor bulk, which is prepared as in the above traditional method, is heated to a temperature such as 1050 °C~1200 °C above its peritectic point where it melts incongruently into Y2BaCuO5 (Y-211 solid phase) and Ba-and Cu-rich liquid (BaCuO3 + CuO). Then the semi-solid melt is directionally solidified with cooling in a temperature gradient of at least 50 °C/cm to ~900 °C, followed by a long oxygenation with a heating rate of ~10 °C/h to 400 °C so as to compensate for the slow kinetics of oxygen diffusion in dense Y-Ba-Cu-O bulk. In comparison with solid-state sintered bulk, the Jc of the bulk with the MTG method is significantly improved, from 500 A/cm–2 to 1,7000 A/cm–2, because of (i) the crystallographic alignment along the preferred superconducting direction resulting in eliminating the antisotropy-originated weak link at high-angle grain boundaries, where a redistribution is required; (ii) the formation of dense structure with enhanced connectivity, and (iii) the formation of fewer and cleaner grain boundaries. In order to increase the Jc, the top seeded MTG (TSMTG) method has been developed and has become the most common technique. In this method, the single crystal top seeds such as NdBCO, SmBCO, whose melting point, or peritectic temperature, is about 40 °C~65 °C higher than that of the matrix bulk such as YBCO, play the role of initiating grain growth. In the presence or absence of a favorable temperature gradient, the seed not only ensures a single nucleation site but also permits controlled growth orientation of the grains. The seeding technique enables growth of matrix domains such as YBCO as large as that of the original bulk size, and a single domain can be obtained. The procedure for the TSMTG method is basically similar to that of the MTG method mentioned above. The differences are as follows. The seed must be positioned on the upper surface center of precursor pellet or bulk (cylindrical or hexagonal or square) prior to melt processing. Bulk is first heated to a certain temperature such as 1045 °C~1050 °C when using SmBCO as seeds, lower than the peritectic point of the seeds, for 20~90 min, and then cooled rapidly to the bulk peritectic point followed by a slow cooling to 975 °C~990 °C at a rate of 0.2~1.0 °C/h. Finally, the bulk is cooled down to room temperature at a normal rate of about 100~200 °C/h. To the above final bulk, a post-annealing (oxygenation heat-treatment) for restoration of oxygen content must be performed in flowing oxygen or in high pressure oxygen at 400~700 °C for 40~200 h to change the YBCO structure from tetragonal phase to an orthorhombic one, which is really superconducting phase.14,21–23 Note that many parameters such as the maximum heating temperature and
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its holding time, especially the rate of slow cooling, as well as oxygenation temperature and its holding time and the oxygen pressure, affect the Jc of TSMTG bulk. Up to now, the Jc of TSMTG bulk with a 45 mm diameter under optimal conditions reaches 100 kA/cm–2,24 and the levitation force of the bulk with a 30 mm diameter at zero distance is over 12.6 N/cm–2.22 For increasing the yield and bulk size, multi-seeding MTG techniques have been investigated.25,26 In the meantime, doping with the addition (0.2~1.0 wt%) of Pt, PtO2, or CeO2 powders has been utilized to aid in dispersing the Y-211 particles in the YBCO for enhancing the flux-pinning ability.23,26
19.3.3 YBCO-based coated tapes As is well known, the 2G coated tape-shaped superconducting conductor structure is a sandwich one, which is a YBCO superconductor layer (current carrier)/buffer layers (as a texture base)/substrate (as a carrier). Many traditional techniques such as liquid phase epitaxial-growth (LPE), vapor deposition (VD), and liquid deposition have been introduced both in buffer layer developments and YBCO layer fabrications. In order to obtain high in-plane and out-of-plane textures on the metallic substrate tapes, and on the buffer layer itself, two well established methods known as rolling assisted biaxially textures (RABiTS) and ion beam assisted deposition (IBAD) were developed. The procedure of making RABiTS consists of cool rolling and recrystallization. RABiTS was first used on textured Ag tape and then on textured Ni and its alloy tapes. With RABiTS, textured metallic tapes can be fabricated quickly and subsequent deposition of buffer layer is a simple process of substrate texture extension to a buffer layer. The architectures of buffer layers have been developed from YSZ (Y-stabilized zirconia) single layer to CeO2-YSZCeO2 multi-layers and so on. With this buffer structure, high Jc of >1 M/cm2 (77 K, 0 T) has been achieved in small samples.5,6,27 IBAD is a technique for obtaining a biaxially textured film by simultaneously utilizing an ion beam bombarding the film surface during its growth. This well-textured substrate by IBAT can solve the ‘weak-link’ problem of the YBCO film layer and achieve a high performance, and the Jc has reached up to 1MA/ cm2 (77 K, 0 T) in short sample by the IBAD-YSZ/PLD-YBCO process with a optimized parameters.5,6,28 In order to improve the deposit rate on the substrates, MgO and gadolinium zirconate (GZO) buffer layers have been used to replace the YSZ one because of the extremely slow deposition rate of YSZ.5,6,29
19.4
Needs for joining high temperature superconductors
Nowadays, the length of commercial Bi-2223/Ag tape fabricated with the PIT technique has reached up to 1000 m. However, the longer the tape is, the
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lower its average Jc is, as well as the more inhomogeneous the Jc. Therefore, superconductive jointed tapes are useful for making long superconducting tapes with high Jc by fabricating medium length tape and joining them together. In addition, a superconductive joint technique is needed to interconnect various parts of a superconducting apparatus such as subunits, busbars, as well as for making a reliable current switch for the application of magnetic resonance image (MRI) and nuclear magnetic resonance (NMR). TSMTG technology is the promising method for fabricating MeBCO bulks with high current capability up to 100 A/cm–2. However, it imposes limits on the materials size and shape because of both the much slower growth of the Me-123 such as Y-123 phase and the high viscosity of the melt during texturing. At present, the dimensions of the bulk that can trap magnetic fields high enough for the efficient use in cryogenics are limited to 40~60 mm, while the maximum size of the bulk reaches up to 93 mm.24 In addition, the farther the materials are located from the seed, the more its structure and composition deviate from the optimal one, and as the bulk becomes larger, the yield of the single domain falls off rapidly. However, in practical applications, the construction of many HTS devices made of bulks, such as flywheel energy storage systems, electric motors, maglev transport equipment, needs large and complex-shaped superconducting components. If single domain bulks could be reliably joined with high Jc superconducting joints, the size and shape requirement of superconductive components could be inexpensively resolved. Differently from the joining of structural materials, both electrical and mechanical properties of the joints of HTS material are considered to be the most important factors because they are usually reduced in the joined region. Additionally, the superconductive properties of the original HTS materials should not be obviously degraded by any means while making a joint. So far, HTS materials concerned with joining in practical applications are mainly BSCCO and MeBCO, especially Bi-2223 tapes and YBCO-based bulks. The developments of their joining technologies and the mechanisms will be expounded in the following section.
19.5
Joining of BSCCO bulks
19.5.1 Fusion welding At present, with respect to fusion joining, the flame-melting method and microwave-melting joining have been investigated.30–33 In the case of the flame-melting method, a liquid natural gas-oxygen flame is utilized to heat the edge portions (surfaces) of the two rod- or bulk-shaped specimens to be joined. When the portions are partially melted, the two specimens are joined by quickly putting the molten faces against each other with a relatively low
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force since the viscosity of the melt at the faces is very low, followed by a natural cooling to room temperature. Research results showed that the sequential heat-treatment on the flame melted portion in the unified specimen, i.e., joint, significantly affected both mechanical and superconducting properties of the joints. Some cracks usually occurred around the interface between the flame melted portion and the original one because of the formation of plentiful glass phases during rapid cooling.30,31 Kasuga, et al. 31 have made a comprehensive study of the flame-melting joining of Bi-2212 rods. The primary results are as follows. No serious gaps or cracks were found around the joined portion after the flame-melted joint was annealed with a soft flame at 900~950 °C for ~5 min. In addition, although both the directly flame-melted joint and annealed one were broken in superconductivity, the post-heated joint, under the treatment conditions of temperature of ~830 °C for 50 h to annealed joint, came to be a superconductor through the transformation of a large amount of 2201 phases to 2212 phases. The postheated joint showed Tc = 89 K and Jc = 155 A/cm2, both of which were almost the same as that of the original Bi-2212 rod. Microwave with a frequency of 2.45 GHz is used as a heating resource in microwave-melting joining. During the joining, the joining region is located in the maximum electric field strength zone of the wave system, and an axial compressive load is applied to the end of each specimen to be joined. In addition, the joining is generally carried out at a temperature range of 900~1000 °C for about 10 min, and the joined sample must be finally postheated at about 855 °C for various times. Cai et al.32,33 have investigated the microwave joining of the Bi1.6Pb0.4Sr2Ca2Cu3Oχ bar fabricated by the codecomposing method, with a Tc of 107 K. The results indicated that the superconducting connection was not established in the joined region because no zero resistance was found due to the formation of some low-Tc 2212 phase in the joining region. In contrast, the Tc of the joint after post-heating at 855 °C for 48 h was restored to 104 K. Further, the joint after post-heating at 855 °C for 60 h possessed the approximate room-temperature resistance and Tc as those of the original Bi,Pb-2223 bar due to complete transformation of the 2212 phase to 2223 phase, called oxygen content restoration, during the long post-heating in air. Additionally, the microwave joined region has higher strength than that of the original part.
19.5.2 Diffusion bonding Diffusion bonding of BSCCO bulks or rods includes two types of methods either with interlayers or without interlayers. Earlier in 1990s, Suzumura et al.34 reported the diffusion bonding investigation of BSCCO pellets with the sizes of 14 mm or 12 mm in diameter and 2 mm in thickness, which were fabricated by a solid state sintering method at 840 °C for 20 h in air, with the
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oxide mixture of In2O3 and Ag2O in a same percent mol value as bonding interlayers. The In2O3 and Ag2O were mixed using organic solvent and then printed on the joining surfaces of both pellets using a screen of 300 mesh. The joining was performed at a temperature range of 830~870 °C for 5~180 min with a pressure of 2.4 Pa in air. The results indicated that the joint shear strength increased up to about 6.3 MPa as the bonding temperature and holding time increased, while the joint superconductivity appeared until the holding time reached up to 90 min at 850 °C. The Tc changed in terms of the distance from the bonding interface, and rose to 90 K at the joint interface. During heating bonding, the liberated Ag due to the thermal decomposition of the Ag2O oxide easily diffused into the superconducting pellets, while In 2O 3 oxide reacted with pellets to form Bi 3 In 5 compound. Some minor elements such as In and/or Ag could improve the superconductivity itself. In the same period, Mutou et al.35 studied the diffusion bonding of BSCCO tablets with dimensions of 8 mm in diameter and 7 mm in thickness, which were fabricated by a melt-pump-up method in a silica glass tube and subsequent annealing at 840 °C for 250 h in air. The joints involved three types, i.e. type I representing the case for bonding of as-pumped-up tablets, type II for annealed tablets, type III for annealed tablets with an as-pumped-up insert layer with thickness of 2 mm. The mean roughness of all the bonding surfaces, polished with finer grades of emery papers, was less than 0.4 µm. Two topon-top aligned tablets to be bonded, with or without insert layer, were hotpressed at 780~850 °C and 0.01~1.0 MPa for 2 h in air where heating and cooling rates were controlled at 10 K/min and 4 K/min, respectively. The diffusion-bonded joints were successively annealed at 840 °C for 250 h in air for obtaining superconductivity. The results were as follows. The diffusion bonding of as-pumped-up materials is possible within a wide range of bonding parameters with the joints’ Tc (105 K) and Jc (247 A/cm2) almost similar to the Tc (100~105 K) and Jc (150~200 A/cm2) of the original tablets, respectively. However, for the direct diffusion bonding of annealed tablets, the Jc (25 A/cm2) of the joint with a Tc (105 K) is much less than that of the original annealed tablets because of the formation of large defects in the bonding interface. Inserting as-pumped-up layer as bonding materials can obviously improve the diffusion bonding joints’ Jc (81 A/cm2) of annealed tablets. Additionally, the shear strengths of all three types of joints are sufficiently high at about 12MPa, almost same as that of the original pellets. As a comparison, Tseluevskii, et al.36 reported the diffusion bonding of Bi-2212-type superconductor tube using a mixture as an interlayer, which composition with respect to the cations was the same as that employed in preparing the tubes. The bonding procedure consists of two stages, i.e., first pre-melting for forming joining at a temperature somewhat lower than the 2212 phase melting point of 880~900 °C and then annealing at
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780~820 °C for 72 h. The results indicate that the Tc of the joint is 79 K lower than Tc (86 K) of the starting tubes. For fabricating large-sized superconductors, Haseyama et al.37 also bonded four Bi-2223 bulks with dimensions of 20 mm × 10 mm × 2 mm, using superconductor paste as bonding materials. During joint forming, the superconductor paste, prepared by mixing the calcined powders and organic solvent, was first sprayed several times onto the bonding surfaces of the bulks and then the four sprayed bulks to be bonded were stuck together. Subsequently, the assembled sample was sintered at 850 °C and 0.04 MPa for 50 h in air, followed by cold pressing and three times sintering. The joints with optimal parameters possess the superconducting properties of Tc = 104 K and Jc = 7000 A/cm2 almost identical to those of the original sintered bulks, while the joint Jc in the magnetic field of 20 mT perpendicular to the assembled sample surface is degraded to 40% Jc of the sintered bulks. Similarly, for making superconducting prisms, Vipulanandan et al.38 researched the bonding of un-reacted or green monolithic BPSCCO bulks and composite BPSCCO ones containing the addition of 25 wt% Ag. Two kinds of joint configurations, i.e., butt and scarf with the angle of 45 °C, were adopted, as well as two kinds of joining technologies, i.e., direct joining and hot pressing joining. For direct joining, Ag paste was used as an interlayer kept as thin as possible to adhere the bulks, and coated on the bulks for forming a butt joint. The butt joint was obtained by pressing the two buttassembled Ag-coated bulks, followed by a sintering at 840 °C for 70 h in air after the Ag coat was dried. In contrast, for hot-pressing methods to form both butt and scarf joints, two processing stages were studied. One stage included a 90 h sintering at 840 °C and one round of cold pressing, or a 160 h sintering at 840 °C and two rounds of cold pressing. The other was successive annealing at 840 °C for 70 h in a 7% O2 balanced argon atmosphere for restoring the oxygen content of Bi-2223 phase. The Ic of joints reached 142~169 A, corresponding to 93~104% superconductivity efficiency.
19.6
Soldering of BSCCO tapes
With respect to the joining of BSCCO tapes, soldering is considered a resistive-joint method compared with the superconducting-joint one. For soldering, the low melting point filler materials such as Bi-Pb-Sn-Cd, Sn-Pb, Sn-Cu-Ag, Sn-Pb-Ag and Sn-Pb-Sb alloys, pure In, Ag-butyl acetate (Ag paste), etc., are utilized and sandwiched with the two tapes to be joined through different overlapped lengths (10 ~ 200 mm). The overlapped sample is generally heated to 200~300 °C and held for several minutes, accompanied by controlling the soldering zone thickness of about 20 ~ 40 µm. The soldering qualities are usually characterized with the contact resistance and Ic or CCR (critical current ratio, i.e., the ratio of the critical current across the whole
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joint to that in the unjoined region), which are dependent on many parameters, especially filler materials features including types, soldering zone width and uniformity between the two tapes, overlapped length as well as the overlapping structure.39–42 Kim et al.39 reported that the contact resistance of the joined mono-filamentary Bi-2223/Ag tapes with overlapped length of 10 mm is 7.1 × 10 –5 Ω, 5.2 × 10–5 Ω, 3.3 × 10–5 Ω, 1.3 × 10–5 Ω, respectively, when using 50Bi-25Pb12.5Sn-12.5Cd, 63Sn-37Pb, 99.9In and Ag-paste. As a comparison, the resistivity of the fillers materials at 77 K is 1.39 × 10–7 Ω·m, 5.0 × 10–8 Ω·m, 1.3 × 10– 8 Ω·m, 2.9 × 10–9 Ω·m, respectively. The above consistent trend indicates that the contact resistance of the joined tape decreases with decreasing of the filler materials receptivity. The contact resistance value order of the joined tapes decides the value order of both Ic and CCR of the joints. In other words, the smaller the contact resistance value is, the higher the Ic and CCR are. In fact, the Ic and CCR of the joints varies from 1.5 A, 3.8 A, 6.5 A to 11.1A and from 10%, 26%, 44% to 77% in the order of Bi-Pb-Sn-Cd, Sn-Pb, In to Ag-paste soldered joints due to the contact resistance value trend mentioned above. Further, for the resistive-joint method, the obvious degrading of the joined tape Ic, compared with that of the unjoined region undergoing the same soldering thermal cycle, is considered to result from the fact that the superconducting current flows through the Ag sheath and the soldering zone with a width of about 20 ~ 25 µm, leading to a smooth transition from a superconducting state to a normal one with increasing current. In contrast, after systematically investigating the soldering of 61-filamentary Bi-2223/Ag tapes with 60Sn-40Pb alloy, using rosin as the flux both numerically and experimentally, Gu et al.42 pointed out that the CCR of joints is significantly dependent on the overlapped length, i.e., the CCR increases with increasing overlapped length when the soldering zone width is 40 µm constant, reaching 100% at 45 mm length. Only in terms of the Ic criterion, the results indicate that the influence of the joint region can be ignored. In other words, the tape is well soldered and can be regarded as a tape without joint. However, in fact, other criteria such as decay behavior of induced magnetic field from the closed loops made by the resistive-joints technology are also needed. The results in the literature39 indicate that the induced magnetic field from the loop soldered with Sn-Pb filler metal would decay to 0 gauss in less than 0.1 s, more quickly than in the loop fabricated by filament to filament superconducting solid-state diffusion bonding. Note that soldering is nowadays mainly applied in making HTS magnets such as double-pancake coils. This joining method with low temperature thermal cycle can obviously reduce the degradation of the original tape and is very simple. Additionally, the contact resistance of soldered joints can meet the demands of common HTS magnet design.
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19.7
Microjoining and nanojoining
Diffusion bonding of BSCCO tapes
19.7.1 Conventional diffusion bonding without an interlayer Design of joint structure and bonding technology Conventional diffusion bonding (hereafter denoted CDB) technology, which consists first of cold-weld under a high pressure at room temperature to form a joint and then high temperature reaction-annealing under no pressure with or without 1~3 times intermediate pressings at room temperature to improve the superconductivity of the above joint, is based on the mechanism of the PIT-fabricating method for BSCCO tape itself. The CDB method is divided into the CDB without an interlayer and the CDB with an interlayer. In other words, both for obtaining high-content superconducting phases Bi-2223 in the bonding zone and obtaining a high crystalline-orientation, a high density of oxide superconducting filaments (or cores) and good connections between the filaments, it is imperative both to impose enough high pressure for the cold-weld and perform a long-time reaction-sintering with several times intermediate pressing to the cold-welded tapes. CDB without an interlayer is one of superconducting-joint methods. The original tapes to be bonded consist of two types. One is the no- or partially sintering-reacted tape with primary phases of Bi-2212, while the other is a completely sintering-reacted tape with primary phases of Bi-2223 transformed from Bi-2212 phases. Similarly, the BSCCO tapes are also classed into mono-filamentary tape and multifilamentary tape. CDB technology without an interlayer includes the following stages in details. Firstly, for achieving superconducting joining, a suitable ‘window’ so as to expose the superconducting cores is necessarily prepared, which generally consists of three types: type A (located near the end of the tape, with keeping a ~2 mm wide Ag- or Ag-alloy tab retaining), type B and type C (both located at the end of the tape, which difference exists at the core edge, i.e., retaining the core edge for type B and mechanically flaking off the core edge for type C), shown in Fig. 19.1(a), (b) and (c), respectively. The Ag- or Ag alloy-sheath on one side of the two tapes is selectively removed by a mechanical or chemical method or a combination of them, and generally, both a razor and tweezers are utilized for the mechanical method, while a 50% HNO3 + 50% H2O2 solution is used as an etchant for the chemical method. Especially for multi-filamentary tape joining, modifying the window by repeating the etching and mechanical removal above mentioned is important in order to ensure more cores are connected. As a result, the window (contacting surface) becomes stepped, and the number of steps consists of 0, 1, 2, 3, 4, etc., correspondingly accompanied by specimens with no step, 1step, 2 steps, 3 steps, 4 steps, etc., as typically indicated in Fig. 19.2.
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Core Ag
Core Ag
(a) Type A
595
Core Ag
(b) Type B
(c) Type C
19.1 Schematic drawings of joining ‘window’ on one side of each mono-filamentary tape to be joined. There are three types: type A, type B and type C.
Secondly, the exposed superconducting cores of the two tapes are carefully brought into contact with each other in a certain type (such as Type I, Type II, Type III, Type IV), as shown in Fig. 19.3, sometimes with the joining region wrapped in an Ag foil. Subsequently, only the joining region or total tape to be joined is uniaxially pressed with a high pressure to cold weld the Ag or Ag-alloy near the window and ensure the cores in the windows of the two tapes contact each other. Thirdly, the cold-welded tape is annealed at 800~850 °C in air. Fourthly, the above joined tape is sequentially pressed at room temperature with a same pressure and annealed at the same temperature 0~3 times. The schematic drawing of the joined tape morphology and the locations of the voltage probes on the joined tape for critical current measurement with a standard four-probe technique are indicated in Fig. 19.4. A joint is usually divided into three regions: unjoined region (a-b and e-f), transient region (b-c and d-e) and joined region (c-d). CCR is defined as the ratio of Ic (a-f region)/Ic (a-b or e-f region). Both the Ic (a-f region) (briefly called Ic) and the CCR of jointed tape are generally used to characterize joining quality. Effects of bonding parameters on the joining quality CDB parameters mainly include uniaxial pressure, bonding (also called annealing or sintering) temperature, (total) holding time, ambient atmosphere, number of window steps, times of pressing-annealing thermo-mechanical cycles, etc. The effects of bonding parameters on joining quality are typically shown in Table 19.2.38,39,43–57 Tkaczyk et al.44 first fabricated the no- and partially sintering-reacted tape joint using CDB method. The results indicate that annealing is beneficial to repair the damage such as micro-cracks associated with the removal of the Ag and the room temperature pressing, and approximate 50% current capacity of the partially sintering-reacted tapes themselves is obtained. Generally, the Ic across the whole joint is determined by the Ic in the transition region and degraded because of the constriction of the superconducting core, Ag intrusion into the cores and irregular microstructure in this region, consistent with the
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Ag alloy sheath
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Specimen with no step Core
Specimen with 1 step
Specimen with 1 step
Specimen with 2 steps
Specimen with 2 steps
Specimen with 4 steps
Specimen with 4 steps
(a) (b)
19.2 Schematic drawings representing the typical lapping ‘window’ (a) and contacting surfaces (b) of 61-filamentary tapes for forming cold weld joints by a high pressure at room temperature. WPNL2204
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Core
Ag Ag
Core
Core
(a) Type I
(b) Type II
Ag Ag
Ag Ag
Core
Core
(a) Type III
(d) Type IV
19.3 Schematic drawings representing the four types of ‘window’ contacting surface of the mono-filamentary tapes for forming cold weld joints by a high pressure at room temperature.
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Ag Ag
597
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a
b
c
d
e
f
19.4 Schematic drawing of the joined tape morphology and the locations of the voltage probes on the joined tape for the critical current measurement with a standard four-probe technique.
results in the literature,39,45,47,51,52,58 as shown in Fig. 19.5. Besides, transverse micro-cracks in the transition region, which is induced during high pressure pressing by the shear stress due to different thickness and reduction between the joined region and the unjoined region, are also limiting factors.51,55 In order to decrease the Ag intrusion and irregular microstructure, the type III of contacting surface is often adopted.51,55 For no- and partially sintering-reacted tape joining, the parameters of bonding temperature, holding time, times of pressing and then annealing, are critical. Only suitable technology factors such as enough high bonding temperature, long sintering time, multi-cycle of pressing-annealing can achieve a large amount of Bi-2223 superconducting phases and improve the texture degree and the density of the cores, followed by reducing the formation of non-superconducting phases and micro-cracks. Yoo et al.47 expounded that annealing temperature of 845 °C compared to that of 835 °C and 840 °C could significantly enhance the amount of Bi-2223 phases. Additionally, an incomplete formation of joint was found after single pressing and annealing at 835 °C for 150 h. In comparison, after multiple pressing and annealing at 845 °C for 50 h, 50 h, 150 h, respectively, the joint was well formed and texture degree was very strong with the larger grains of Bi-2223. Similarly, Huang et al.55 reported that Ic of the joint and regular tape reached a peak value at ~250 h using a temperature range of 800~850 °C. However, longer annealing would result in the decomposition of Bi-2223 phases and decrease the Ic. The research of Sha et al.48 showed that the Ic of a joint made from completely reacted tape is higher than that of a joint made of no-sintering reacted tape (also called green tape). This phenomenon results from the excellent superconducting zone near the Ag/core interface being introduced into the joint center, which would lead to a larger enhancement of the higher Ic superconductors in the case of utilizing completely reacted tapes than that using no-sintering reacted tapes. In the case of joining completely reacted BSCCO tape, Kim et al.39,51,52 emphasized the effects of uniaxial pressure, type of window shape and contacting surface on the Ic and CCR of the joints. For both type I and type II contacting surfaces, CCR values are as low as 40~50% at a low pressure
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Table 19.2 The bonding parameters and the best properties of BSCCO jointed tapes by conventional diffusion bonding without an interlayer Pressure of room temperature pressing × annealing temperature × holding time
Times of pressing Max. Ic of and then annealing lap-joined thermomechanical cycle, tapes ambient atmosphere
Max. CCR of References lap-joined tapes
햲 햳 햴 햵 햲
C, C, B, C, B,
1, air 1, air 1, air 1, PO2 = 0.076 atm 1, air
14.1 A 16.4 A 25.0 A
67% 65%
44 44
80%
45
2.7 A 2.2 A
75% 73%
46 47
햲
B, 12 mm × 2.5 mm
3, air
9.4 A
63%
47
햲 햴 햴 햴 햴 햴 햴 햶 햲
C, 12 mm × 3 mm C, 12 mm × 3 mm B, 2 mm × 3~4 mm B, 2 mm × 3~4 mm C, 10 mm × 1.8 mm B, 10 mm × 1.8 mm A, 10 mm × 1.8 mm B, 10 mm × 4.2 mm C, 20 mm × 2 mm
984 MPa × 830 °C × 48 h 984 MPa × 830 °C × 48 h 1250~3750 MPa × 840 °C × 100 h 444 MPa × 830 °C × 100 h 835~845 °C × 100~150 h No pressure value provided 835~845 °C × (50 h + 50 h + 100~150 h) No pressure value provided 835~845 °C × (50 h + 50 h +100 h) No pressure value provided 1000~1900 MPa × 835~840 °C × 70 h 0~5500MPa × 835~840 °C × 70 h 140~4000 MPa × 835~845 °C × 50h
3, air
17.0 A 25.0 A 12.5 A — —
85 % 89%
48 48
99%
49
90% 95%
49 39
83% 88 % 65% 30% 80% (butt joint)
51
20 mm × 2 mm 20 mm × 2 mm 5 mm × 8 mm 18 mm × 2.5 mm 12 mm × 2.5 mm
1000~2500 MPa × 840 °C × 50 h 1000 MPa × 840 °C × (50 h + 50 h + 50 h + 100 h)
2, air 1, air 1, air
1, air 4, air
— 20.0 A 52.0 A (butt joint)
52 53, 54 56 56
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햲: no-sintering-reacted mono-filamentary tape. 햳: partially sintering-reacted mono-filamentary tape. 햴: completely sintering-reacted monofilamentary tape. 햵: completely sintering-reacted 9-filamentary tape, with 3 steps in window for joining. 햶: completely sintering-reacted 37-filamentary tape, with 0, 1, 2, 4 steps in window for joining.
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Original Window type, tapes length × width
600
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Unjoined region
Transition region
Joined region Ag
Ag
Core Ag
150 µm
19.5 SEM micrograph of a longitudinal cross section for lap-joined mono-filamentary tape fabricated with a pressing at room temperature under 1600 MPa in type II contacting surface style, showing the three regions.
of 140 MPa, which increase with pressure and reach up to 80~90% at 1600 MPa, and then decrease with a larger pressure. An unbonded interface in the joined region can be found at low pressure, agreeing with the result in the literature,58 as indicated in Fig. 19.6. On the other hand, an appropriate increase of the unaxial pressure can reduce and even completely avoid the unbonding, promote the core connection along the a-b plane of Bi-2223 grain, and lead to improving of the core density and Bi-2223 grain texture degree, resulting in the increasing of Ic and CCR. However, an over-pressure would induce more micro-cracks, a more irregular interface and more severe Ag-intrusion, accompanied with a degradation of Ic and CCR. In order to restrict Ag intrusion into the core, type C window and type III contacting surface style are designed. In this case, an Ag sheath in the transition region can deform more uniformly, resulting in a more straight and a more deformed interface, and therefore, the CCR can reach up to 90~95%. In addition, the results indicated that the tensile strengths of the joint made of type A window and type B window are 86 MPa and 77 MPa, respectively, both somewhat lower than that of the unjoined tape. Besides that, the Icirreversible tensile strain of the joint made of type B window is only about 0.1%, while the irreversible tensile strain of the unjoined tape is about 0.3%.51 In the case of multi-filamentary tape joining, it is imperative to prepare the stepped windows and carefully bring the windows of two tapes into contact with each other, as shown in Fig. 19.2. Kim et al.53,54 investigated the CCR as a function of the number of window steps by joining 37multifilamentary BSCCO tape. It was proved that the CCR is dependent on the step number, and is steadily enhanced from 25% to 58% as the step number increases from 0 to 4. The enhancement is considered to be due to the fact that more superconducting filaments from the two tapes are connected to each other in the joined region. In comparison, Lee et al. 46 also demonstrated the necessity of a multi-step window for joining 9-multifilamentary tapes.
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Ag sheath
Superconducting core Unbonded interface
Superconducting core Ag sheath 30 µm
19.6 SEM micrograph of a longitudinal cross section for monofilamentary lap-joined tapes with a room temperature pressing under 800 MPa, showing the unbonded interface in the joined region after the annealing.
19.7.2 Conventional diffusion bonding with an interlayer In order to achieve superconducting joining by using CDB with an interlayer, it is essential to select appropriate superconducting materials as an interlayer. So far, only Wang et al.59 have systematically studied the joining of both noand completely sintering-reacted 37-multifilamentary BSCCO tapes, with CDB using both BSCCO precursor powders and precursor powders added by 10% Ag. The two tapes were first prepared into type B windows with a length of 40 mm and a width of 4 mm, and then carefully contacted with each other to form a lap joint. Meanwhile, proportional powder interlayers were filled between the two exposed cores. Secondly, the assembly was pressed together under 100 MPa to form a cold-joined tape. Finally, for nosintering reacted joined tapes, a series of thermo-mechanical treatments were performed, i.e., initially sintering at 835 °C for 40 h in air, secondly uniaxially pressing under 1000 MPa and sintering at 835 °C for 60 h in air, thirdly pressing under same pressure and sintering at 835 °C for 35 h in air, finally pressing and reaction-sintering again at 834 °C for 65 h so as to crystallize the high-Tc Bi-2223 grains and heal the cracks associated with a uniaxially mechanical pressing. For completely sintering-reacted joined tapes, however, only sintering at 835 °C for 50 h in air was performed. For no-sintering reacted joined tape, the maximum CCR value is only 11.7% under the conditions of utilizing BSCCO precursor powder as an interlayer and sintering at 835 °C for 100 h in air with an intermediate pressing as above. In contrast, completely sintering-reacted joined tape, the maximum CCR reaches 47%. In comparison, Wang et al.59 also researched the joining of the same tape with direct method, i.e., that no window was
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fabricated, using three kinds of contacting surfaces including direct contacting without interlayer, with BSCCO precursor powders and 10% Ag added precursor powders. For no-sintering reacted tape, the maximum CCR of direct joints is 33% as no interlayer is inserted, while the maximum CCR of direct joints for completely sintering-reacted tapes increases to 77% as using 10% Ag added precursor powder. Further, analysis of the microstructures indicated that there lie some damage and more irregular microstructures, both induced during the pressing, in the transition region of the lapped joined tapes, which result in the degradation of CCR. On the other hand, for direct joints, the interfaces between the cores and sheath are smooth and straight, and the microstructures both from longitudinal and traverse sections are more regular, which are beneficial to the improvement of CCR. Additionally, the retaining Ic of the joints is decreased more quickly than that of the unjoined tapes with the tensile stress increasing due to the irregular interfaces and Ag intrusions into the cores in the transition region. Guo et al.60 also investigated the CDB with an interlayer by a two-step sintering process, first at 920 °C for 0.5 h and then 840 °C for 120 h, on the joining of PIT-fabricated mono-filamentary BSCCO tapes. The BSCCO precursor powder was used as an interlayer inserted in the two type B windows of the two tapes to be joined, and the two-step technology of first premelting the interlayer after wrapping the contacted tapes in a Ag foil and then sintering at 840 °C for given hours was adopted. The results showed that the Ic of the joint and unjoined region was 28 A and 25 A, respectively, degraded compared with Ic of 40 A of the original tape. Additionally, the strength of the joined region is 25.6 MPa, only 61% of the unjoined region.
19.7.3 Direct diffusion bonding without an interlayer Direct diffusion bonding method As explained above, CDB needs a long time, up to 50~200 h, to form a superconducting joint, accompanied by obvious degradation of Ic of the original completely sintering-reacted tapes, resulting in the critical current capacity decrease of the entire joined tape.40,41,60 Zou et al.61 also reported a similar phenomenon when 61-filamentary BSCCO tape underwent annealing at high temperature of 780 ~ 835 °C for more than 5 h. It is considered that the degradation of Ic is due to the decomposition of high superconducting phases Bi-2223 into low superconducting phases Bi-2212 and other nonsuperconducting phases such as Ca2CuO3, CuO, etc. In addition, high temperature pressing would improve the density and texture degree of superconducting cores (filaments), leading to decreasing degradation of Ic, consistent with the effects of high temperature pressing on the Ic of the PITfabricated tape.62–64
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For improving the joining efficiency and avoiding Ic degradation of the original reacted BSCCO tapes, direct diffusion bonding (hereafter denoted DDB) technology with a high temperature pressing has been proposed.65,66 In this case, two multi-filamentary tapes with core-exposed windows are carefully contacted and uniaxially pressed together, under a much lower pressure compared with the CDB method, and then annealed under the above low pressure at high temperature for several hours in air. The typical windowshape and surface-contacting style of 61-filamentary tape, similar to that of CDB, are shown in Fig. 19.2. Zou et al.43,65–68 systematically investigated the DDB without an interlayer of 61-filamentary tape with a width of ~4 mm and a thickness of ~0.2 mm, which was fabricated by the standard PIT method and with a scanning electron microscopy (SEM) micrograph from a cross-section is shown in Fig. 19.7. The affecting factors of the joint properties usually characterized by the interface bonding, microstructures, Ic and CCRo, mainly involve bonding pressure (PB), holding time (tB), bonding temperature (TB), lapped-window length (L), number of the window (n), etc. Here, the CCRo in terms of the common definition in welding science and technology is defined as the critical current ratio of Ic (the joined tape critical current) to Ico (the original tape critical current), which is different from and more strict than the CCR used in CDB joining method as above mentioned. As is well known, Ag alloy sheath has a high ductility. To avoid bonding between the pressure-transferring Al2O3 ceramic column and the Ag alloy sheaths during DDB, ZrO2-based powders or isinglass foils have been selected to be used as baffle fluxes. Effects of bonding parameters on the joint properties The effects of PB, tB, TB and L on the CCRo are shown in Figs 19.8–19.11. A SEM micrograph of a longitudinal section for a typical 61-filamentary BSCCO joined tape with DDB is displayed in Fig. 19.12, indicating the three regions: unjoined region, transition region and joined region. The results indicate that CCRo is first increased with PB and reaches the peak value 89% at 3 MPa, and then falls down. The laws of CCRo dependence on the tB and TB are similar, and the optimal values are 2 h, 800 °C, respectively. As we know, the main phases Bi-2223 of the tape cores are chemically inactive. Therefore, when using the CDB method, the annealing needs a long time both to form a firm interface-metallurgical bonding and heal up the cracks induced during cold weld. In contrast, high temperature direct pressing in DDB method is beneficial to interfacial close contact, promoting the interdiffusion of the interfacial Bi-2223 grains. Figure 19.13 shows that a lower pressure pressing would lead to a loose bonding interface, and inversely, a tight bonding interface can be achieved under a higher pressure. Additionally, the density of the cores and the texture degree of Bi-2223 grains can be
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100 µm
Ag-alloy sheath
Superconducting core
19.7 SEM micrograph of the central position from cross section of 61-filamentary completely reacted BSCCO tape (from Beijing Innova Superconductor Technology Co., Ltd).
Critical current ratio, CCRo/%
100 80 60 40 20 0
0
2 4 6 8 Bonding pressure, PB /MPa
10
19.8 CCRo of BSCCO joined tape vs. bonding pressure. Here, TB = 800 °C, tB = 2 h, L = 20 mm, n = 0, Ico = 85 A. 100 Critical current ratio, CCRo/%
604
80 60 40 20 0
0
2
4 6 Holding time, tB /h
8
10
19.9 CCRo of BSCCO joined tape vs. holding time. Here, TB = 800 °C, PB = 3 MPa, L = 20 mm, n = 0, Ico = 85 A.
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Critical current ratio, CCRo/%
100 80 60 40 20 0
780
790 800 810 820 Bonding temperature, TB /°C
830
Critical current ratio, CCRo /%
19.10 CCRo of BSCCO joined tape vs. bonding temperature. Here, PB = 3 MPa, tB = 2 h, L = 20 mm, n = 0, Ico = 85 A.
100 80 60 40 20 0
5
10 15 20 25 Lapped-window length, L /mm
30
19.11 CCRo of BSCCO joined tape vs. lapped-window length. Here, PB = 3 MPa, tB = 2 h, TB = 800 °C, n = 0, Ico = 85 A.
improved. But an overpressure would result in more Ag alloy intrusion into the cores in the transition region and more irregular interface between the cores and the Ag alloy sheath, leading to the degradation of the CCRo, consistent with that of CDB method. Similarly, inter-diffusion for interface formation demands an enough time, which would be significantly shortened because of the utilizing of a high temperature pressing. However, too long bonding would debase the superconductivity of the tapes due to the decomposition of the 2223 phases. In addition, Bi-2223 phase is more stable at about 800 °C,61 resulting in the highest CCRo value. Therefore, there are both optimal tB and optimal TB with respect to the CCRo of the joints.
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Bonding interface
200 µm
Transition region
Joined region
Unjoined region
19.12 SEM micrograph of a longitudinal section for a typical 61filamentary BSCCO joined tape fabricated using DDB method without an interlayer, showing the three regions of the joints: unjoined region, transition region and joined region. Here, TB = 800 °C, tB = 2 h, PB = 7 MPa, L = 20 mm, n = 0, Ic = 61.8 A.
Note that, from Fig. 19.11, the CCRo increases with L when other bonding parameters are constant. It is considered that increasing L would enhance the superconducting core contacting area between the two tapes, beneficial to the current transfer and improvement of CCRo. In particular, the CCRo of the joined tape with a L longer than 25 mm should be defined as the ratio of critical current in the joined region to that of the original tape, i.e., the ratio of Ic (c-d)/Ico (the original tape critical current) as shown in Fig. 19.4, just because all the probes in the standard four probe technique are positioned within the long joined region. Thus, CCRo is more than 100%, resulting from the higher critical current capacity of the joined region with a doublethickness superconducting cores compared with the original tape. For the joining of 61-filamentary BSCCO tapes with the DDB method, no obvious effects of the lapped-window step number on the joint superconducting properties was found, differing with the results of joining 37-multifilamentary BSCCO tapes with CDB method.53,54 It is considered that the distance between the two neighbor cores is much less in a 61-filamentary BSCCO tape, sometimes resulting in connecting to each other, as shown in Figs. 19.7 and 19.13. Therefore, increasing the number of the lapped-window steps could not enhance the CCRo since the effective connecting area between the cores is approximate among all the joints with various steps. In summary, adopting DDB method without an interlayer can indeed save time and improve the joint superconductivity significantly, i.e. the joining efficiency of BSCCO tapes.
19.7.4 Direct diffusion bonding with precursors For comparison, recently, Zou et al. also studied the joining of 61-filamentary BSCCO tapes with DDB method using BSCCO precursor powders as the
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Bonding interface
20 µm (a)
Bonding interface
20 µm (b)
19.13 SEM micrographs of a longitudinal section for 61-filamentary BSCCO joined tapes with DDB, showing different microstructures and bonding interfaces under lower pressure of 0.2 MPa (a) and higher pressure of 1 MPa (b). Here, TB = 800 °C, tB = 2 h, L = 20 mm, n = 0.
interlayer. The maximum CCRo only reached 49% because most of the Bi2212 phase in the precursor powders still remained in the bonding zone after a relatively short time reaction-annealing, as shown in Fig. 19.14. In other words, no jointed tape with a Bi-2223 phase of high content is formed,
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Bonding zone
Ag-alloy 20 µm
19.14 SEM micrograph of a longitudinal section for 61-filamentary BSCCO joined tapes fabricated with DDB method using BSCCO precursor powders as interlayer, showing most of Bi-2212 phase remaining after short time annealing. Here, TB = 820 °C, tB = 5 h, PB = 3 MPa, L = 20 mm, n = 0, Ic = 44.1 A, Ico = 90 A.
resulting in low superconductivity. Further investigation is still needed to be performed.
19.8
Joining of YBCO bulks
19.8.1 Diffusion bonding Early in 1992, Salama et al.69 reported diffusion bonding without an interlayer for YBCO bulks with microstructures of well-aligned grains in the form of plates oriented along the a-b plane and stacked along the c-axis. The bulks were fabricated using MTG processing method, and were 1~2 mm by 1~2 mm in cross-section and 8~10 mm in length. Firstly, the bonding surfaces of the samples to be bonded were polished to yield a roughness of less than 3 µm. Secondly, the two samples were carefully contacted on the a-c or b-c plane with an area of 2~5 mm2 and the a-b planes of each sample aligned within 5°. Thirdly, the assembly was imposed with a pressure of 2~6 MPa and rapidly heated to 910 °C with a holding for 18 h in air, followed by an annealing of 930 °C for 12 h in air, and then subsequently furnace-cooled to room temperature. Microstructures analysis showed that intimate bonding was formed without secondary phases, and several micro-cracks were formed possibly due to stress relief during cooling. The maximum Ic and Jc of the joint reached 150 A and 6300 A/cm–2 in a zero magnetic field at 77 K, respectively, and were lower than those of the original bulks. The reasons were considered to involve the small bonding area, the misalignment of the
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a-b planes of the two samples and the nature of the interface itself. In comparison, Guo et al.70 also studied the diffusion bonding of MTG-YBCO bars in vacuum, under the conditions of bonding temperature 930~950 °C for 1 h using a pressure of 3 MPa, followed by an oxygen restoration heattreatment at 450 °C for 120 h. The results indicated that the bonding zone mainly comprises non-superconducting tetragonal-structure Y-123 phases after only vacuum processing. The oxygen restoration heat-treatment promotes the transformation of tetragonal Y-123 phases into superconducting orthorhombic Y-123 phases, achieving a joint Tc of 90.5 K identical to that of the original bulks. Ogawa et al.71 researched diffusion bonding with Ag2O or Ag2O-50 mol% PbO pastes as interlayers for the YBCO pellets with 12 mm or 14 mm diameter and 2 mm thickness, which were prepared using solid reaction sintering method under the conditions of 100 MPa hot-pressing at 890 °C for 14 h in air. For the Ag2O paste, the condition of 970 °C bonding temperature for various holding times and 2.1 KPa pressure was utilized. The microstructure analysis showed that Ag2O is thermo-chemically easily decomposed, accompanied with Ag and O diffusing into the YBCO base part to form a multiple structure reaction layer, i.e., YBa2Cu3O7–X and Y2BaCuO5 excluding Ag. On the other hand, for the Ag2O-50 mol% PbO paste, BaPbO3 formation was found at over 930 °C for the same holding time and bonding pressure. The reaction layer near to the YBCO base part is a multiple structure including Ag, while the outer reaction layer is also a multiple structure excluding Ag. The Tc and strength of a joint are determined by the bonding conditions. The maximum Tc and shear strength of the joints with Ag2O paste are ~90 K and 15 MPa, respectively, while the maximum Tc and shear strength of the joints with Ag2O-50 mol% PbO paste are ~88 K and 10 MPa, respectively.
19.8.2 Soldering In soldering of MTG-MeBCO bulks, the joining is accomplished by using the two MeBCO bulks themselves to seed the melting solder, which has a lower decomposition temperature and similar crystal structure as compared with the MeBCO base materials to be joined. Up to now, soldering has been the most promising joining method for MTG-MeBCO bulks and has been systematically studied mainly on YBCO-based bulks or rings.72–89 In the joining operation, firstly, the two YBCO based-bulks or rings to be joined are carefully assembled with the a-c or b-c planes parallel to each other and the solder inserted between the two joining surfaces. Secondly, the soldering is performed in air with a special heat-treatment schedule (temperature vs. time), as shown in Fig. 19.15, and explained as follows. The assembly is heated to a temperature of T1, which is intermediate between the solder melting point temperature of Tm2 and the bulk decomposition temperature
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Temperature
T1
V1 V2
T2
V0
T3 Oxygenation T4
V3 t1
t2
t3
t4 Time
19.15 A typical curve of temperature vs. time used for joining YBCObased bulks with a solder such as TmBCO, Ag, Ag-doped YBCO, YbBCO, ErBCO, etc. The first stage indicates the soldering heattreatment schedule in air, while the second one presents the oxygen restoration heat-treatment schedule in flowing oxygen.
of Tm1. The solder, such as TmBCO, Ag-doped YBCO, YbBCO, ErBCO, etc., decomposes into a two-phase mixture of the Ba-rich and Cu-rich liquid and the dispersed solid green phase of R-211 (R = Tm, Y, Yb, Er), while the two YBCO bulks or rings remain solid. The assembly usually stays at the highest temperature (T1) plateau for a few hours when the liquid can wet the YBCO surfaces to be joined. This is followed by a rapid cooling with a rate v1 of about 100 °C/h to T2 somewhat lower than Tm1 with several hours dwelling time of t2. Subsequently, a very slow cooling to T3, which is 20~80 °C lower than T2 with a rate v2 of about 0.2~5 °C/h, is critical for seeded textured-growth of the melting solder to form a joint. Note that the slow cooling step is very important for the formation of textured-seam structurally similar to that of MTG-fabricated YBCO bulks themselves. After that, a cooling rate of about 100~200 °C/h is adopted from T3 to room temperature. As is well known, the high temperature of the soldering procedure usually reduces the oxygen stoichiometry, resulting in non-superconducting of the soldered bulk. Therefore, subsequent to the soldering, the bulk must be heat treated in a flowing oxygen atmosphere at a temperature T4 about half of T1 for a duration that can vary with the size and shape of the joint. This step can fix the oxygen stoichiometry in the YBa2Cu3Oχ to a level near χ = 6.9 or higher, where Tc ≈ 92K. The larger the size of the bulks, the longer the duration for good oxygen homogeneity throughout the soldered bulk, varying from tens of hours to hundreds of hours. So far, a great effort has been devoted to investigating the effects of soldering parameters on the superconducting properties of joints, and high quality soldering joints with Jc and Fo (levitation force) comparable to those of the original bulks have been obtained, as partially shown in Table 19.3.
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Table 19.3 Typical soldering parameters and best properties of RBCO joined bulks with a RBCO-based solder Solders, Tm2, Soldering parameters: Vo, T1, t1, V1, T2, V2, thickness T3, V3, PB, ambient atmosphere
YBCO, 1015 °C
TmBCO, 995 °C, 1 mm YBCO/Ag, 970 °C, 0.7 mm YBCO/Ag, 955 °C, 1 mm YBCO/Ag, 955 °C, 0.7 mm YBCO/Ag
YBCO, 1010 °C
YBCO(0.5%CeO2 + 0.24%SnO2), 1015 °C YBCO(0.02%Pt), 1015 °C YBCO YBCO, 1010 °C YBCO, 1015 °C
YBCO YBCO, 1026°C
99.9% Ag, 960 °C, 50µm YbBCO, 925 °C YbBCO (0.5%PtO2) TmBCO, 983~987 °C, ≤0.3 mm
Vo = 100 °C/h, T1 = 1010 °C, t1 = 2 h, V1 = 70 °C/h, T2 = 980 °C, V2 = 0.5 °C/h, T3 = 955 °C, V3 = 100 °C/h, air T1 = T2 = 970 °C, 990 °C, 1010 °C, V1 = V2 = 0.5 °C/h, slight pressure by the weight of the bulk itself, air T1 = T2 = 990 °C, V1 = V2 (no mention, but slow cooling to room temperature), air Vo = 50 °C/h, T1 = 995 °C, t1: short time, V1 = 50 °C/h, T2 = 973 °C, V2 = 0.2 °C/h, T3 = 940 °C, V3: slow cooling, air Vo = 332 °C/h, T1 = T2 = 995 °C, t1 = 10 h, V1 = V2 = 0.5 °C/h,T3 = 945 °C, air T1 = 1007 °C, t1 = 3 h, V1 (rapid cooling), T2 = 990 °C, V2 = 0.6 °C/h, T3 = 955 °C, air T1 = T2 = 960 °C, t1 = 1~2 h, V1 = V2 = 1~5 °C/h, T3 = 900 °C, V3 = 100 °C/h, air T1 = T2 = 980 °C, V1 = V2 = 1°C/h, T3 = 940 °C, V3 = 100 °C/h, PB = 20 kPa, air Vo (first 1000 °C/h and then 500°C/h), T1 = 990~1050 °C, t1 = 6 min, V1, V2 (cooling with the furnace, average value of 110 °C/h), PB = 1 MPa, pure O2 with a pressure of 105Pa,
Oxygenation parameters: T4, t4, ambient atmosphere
Properties of the joints
References
440~480 °C, 12~336 h, flowing O2 No mention
Jc ≈ Jco Fo ≈ Foo
72
Jc ≈ Jco microcracks
73
Jc ≈ Jco
77 78 80
430 °C, 250 h, flowing O2 380°C, 10 h, flowing O2 520°C, 150 h, flowing O2 No mention 400~470 °C, 192~240 h, flowing O2 450°C, 100 h, flowing O2 first 550 °C and then 450 °C, total 132 h, flowing O2
Jc = 9 × 103 A/cm2 Jc = 2.5 × 104 A/cm2 Jc ≈ Jco 1.2 × 104 A/cm2 Fo = 92% Foo
83 84 74 75 76
Jc = 1.2 × 104 A/cm2 Jc = 3.4 × 104 A/cm2
81 82 85, 86, 87, 88, 89
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Note: Tm1: the decomposition temperature of melt-textured YBCO bulks; Tm2: the melting point temperature of RBCO-based solders or Agmetal solder; Jc: the critical current density of soldered bulks; Jco: the critical current density of original bulks; Fo: the levitation force of soldered bulks; Foo: the levitation force of the original bulks. YBCO/Ag: Ag doped YBCO composite, PB: joining pressure.
Joining of high temperature superconductors
YBCO-based bulks, Tm1
612
Microjoining and nanojoining
According to the results in the literature72–89 for achieving good soldered bulks or rings, optimization must be performed to reduce drawbacks such as non-superconducting phase segregation at the interface where growth fronts of melting-solder merge, voids commonly appearing in the joining region or seam, the mismatch between lattice parameters of the bulks and solder which can lead to cracking. Note that the above mentioned soldering technology without joining pressure in air needs a long time to form a joint. In order relatively to shorten the soldering period, Prihna et al.85–89 proposed a new soldering technology based on the above soldering schedule. In this technology, during the soldering stage, a pure oxygen atmosphere (Po2 = 1 × 105 Pa) and a pressure of about 1 MPa are utilized. Therefore, the vo can be increased up to 500~1000 °C/h and a direct cooling with furnace with a rate of about 110 °C/h from T1 to room temperature can be used. The results support the fact that during the soldering process no degradation of the bulks or rings occurred and that the Jc in the joining region was not lower than 3.4 × 104 A/cm2 at 77 K in zero field. The high quality joining was considered to be due to redistribution over the MT-YBCO matrix phase during the repeated heat treatment and oxygenation.
19.8.3 Other joining processes As is well known, the microstructure of the top-seeded MT-YBCO based bulks generally consists of a pseudo-crystalline YBa2Cu3O7 (Y-123) platelet matrix containing secondary phases, mostly Y2BaCuO5 (Y-211) dispersed in the matrix, Ba-Cu-O phases trapped in platelet boundaries and other defects such as cracks, twins and stacking faults. According to the above microstructure, an ‘interface liquid phase assisted joining’ method has been proposed by Vanderbemden et al.90 Both joining procedure and mechanism are described as follows. Firstly, the two bulks to be joined, such as with a composition of YBa2Cu3O7–X + 30% Y2BaCuO5+0.1% Pt, are arranged with their a-b planes mis-orientated to each other in an angle from 0° to 90°. Secondly, the arrangement is heated to a joining temperature between 920 °C to 980 °C under a pressure of 0.1~0.5 MPa in air. Between 920 °C and the joining temperature, the trapped Ba-Cu-O phases become liquid and this liquid migrates into the surfaces to be joined by the surface tension effects. Once there, the liquid dissolves Y-123 phase which later recrystallizes during cooling, driving the impurity phases out of the boundary and forming a good joint as the bulk directly solidifies. After that, the joint must be oxygenated as in the soldering method. References 90, 92, 93 reported this joining method and the effects of misorientation angle on the superconducting properties of the joints were investigated. The results indicate that 0° angle joining of two a-b planes can produce a joint with a temperature dependence of the irreversibility field
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similar to that of the bulk itself, and with a Jc exceeding 2000 A/cm2 at 77 K in fields up to 4 T. Another special joining method is defined as ‘surface superheating joining’. In this method, the surfaces to be joined are heated to a temperature somewhat higher than the decomposition temperature under a sample self-weight pressure in flowing pure oxygen atmosphere (Po2 = 1 × 105 Pa) to avoid nitrogen trapping, accompanied with a dwelling of a few hours. The joint is formed as the melting surfaces recrystallize during slow cooling, also followed by a oxygenation, similar to the soldering method. Chen et al.94 reported this method, and pointed that a rejoining of two cut samples from the same bulk can produce a joint with a Jc of about 1 × 104 A/cm2, while a joining of two independent bulks can produce a joint with a Jc of about 5000 A/cm2, under the conditions of joining temperature 1020 °C for 0.5~1 h, slow cooling rate 2 °C/h from joining temperature to 940 °C, and oxygenation temperature 450 °C for 168~240 h as joining the YBCO materials. Note that the two special joining methods require a very tight control of temperature gradients in the furnace and the joining quality is highly dependent on the starting micro-structural quality of the bulks. Thus a scale-up of the process might not be very appropriate.
19.9
Conclusion and future trends
Mature joining techniques are requisite to the extensive applications of HTS materials. As mentioned above, much effort has been devoted to develop joining methods and optimizing the joining parameters for tapes and bulks of BSCCO materials as well as YBCO-based bulks. For multi-filamentary BSCCO tapes with a high Jc, direct diffusion bonding with a high temperature uniaxial pressing is a promising method which can improve both the joining quality and efficiency compared with conventional diffusion bonding. However, bonding technologies are needed to be optimized further, and the mechanism involving the interface formation, the changes of both texture degree and density of Bi-2223 phase, the roles of uniaxial pressing, etc., also need to be profoundly studied. Additionally, the sensitivity of joints’ superconducting properties to deformation, magnetic field and thermal-cycle, etc., must be systematically tested in the future. In fact, for obtaining high quality superconducting joints, this diffusion bonding also belongs to the micro-joining technique. Therefore, more research must be performed with an emphasis on superconducting-property-affecting factors including the fine core to core structure of the joint, stress distribution in the cores and Ag-alloy sheath during the joining, etc. For comparison, soldering technology is the most promising joining method for MeBCO-based, specially YBCO-based, bulks or rings. Further research emphases must be concentrated both on eliminating defects in joining the
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region such as non-superconducting phase segregation, voids, micro-cracks and shortening the processing period. In addition, as is well known, worldwide activities are being focused on developing processing technologies and applications for both YBCO-based coated conductors (tapes) and superconducting MgB2/Fe wire fabricated by PIT. So far, both of them with high Jc are still limited in length because of the current immature preparing technique. Joining research is also urgently needed for their practical applications.
19.10 Acknowledgements This work was financially supported by the National Natural Science Foundation of China under grant No. 50575114 and No. 50635050, and by Beijing Natural Science Foundation under grant No. 3052010. The author is very grateful to Prof. Jialie Ren, Prof. Aiping Wu, Prof. Zhenghe Han and Ph.D Hanping Yi with Tsinghua University for discussions. The author also thanks Naijun Hu, Yanjun Wang, Wei Guo, Hailin Bai, Qing Wang, Mingya Li, Chen Gu, Kai Shi with Tsinghua University, as well as Xiuhua Song, Jun Zong, Hongjie Zhang, Suli Wang with Beijing Innova Superconductor Technology Co., Ltd for support of this work.
19.11 References 1. T Tsuneto, Superconductivity and superfluidity (translated by M Nakahara), Cambridge, Cambridge University Press, 1998. 2. A Mourachkine, High-temperature supercondutivity in cuprates, London, Kluwer Academic Publishers, 2002. 3. H Zeng, Essentials of advanced materials for high technology, Beijing, China Science and Technology Press, 1993 (in Chinese). 4. C W Chu, Y Y Xue, Y S Wang, et al., ‘High temperature superconducting materials: Present status, future challenges and one recent example-the superconducting ferromagnet’, Physica C, 2000 341–348 25–30. 5. Y L Xu, D Shi, ‘A review of coated conductor development’, Tsinghua Science and Technology, 2003 8(3) 343–369. 6. Y W Ma, L Y Xiao, ‘Second YBCO coated conductors: A review’, Chinese Science Bulletin, 2004 49(23) 2435–2439. 7. Z Li, Z Wei, ‘Crystal structures and physical properties of copper oxide hightemperature superconductors’, Journal of Gansu University of Technology, 2000 26(4) 104–109 (in Chinese). 8. W Z Zhou, W Y Linag, Fundamental researches of high temperature superconductors, Shanghai, Shanghai Science and Technology Press, 1999 (in Chinese). 9. N Ayai, K Hayashi, K Yasuda, ‘Development of Bi-2223 superconducting wires for AC applications’, IEEE Trans Appl Superconductivity, 2005 15(2) 2510–2513. 10. T J Arndt, A Aubele, H Krauth, et al., ‘Progress in the prepration of technical HTS– tapes of type Bi-2223/Ag-Alloy of industrial lengths’, IEEE Trans Appl Superconductivity, 2005 15(2) 2503–2506.
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11. H P Yi, X H Song, L Liu, et al., ‘Development of HTS BSCCO wire for powder applications’, IEEE Trans Appl Superconductivity, 2005 15(2) 2507–2509. 12. A P Malozemoff, D T Verebelyi, S Fleshler, ‘HTS wire: status and prospects’, Physica C, 2003 386 424–788. 13. P Vase, R Flükiger, M Leghissa, et al., ‘Topical review: Current status of high-Tc wire’, Supercond. Sci. Technol., 2000 13 R71–R83. 14. S Sengupta, J Corpus, M Agarwal, et al., ‘Feasibility of manufacturing large domain YBCO levitatorsTM by using melt processing techniques’, Materials Science and Engineering, 1998 B53 62–65. 15. L Xiao, H T Ren, Study and progress of REBaCuO single domain superconducting bulks, 2003 25 309–314 (in Chinese). 16. Y X Chen, L Xiao, H T Ren, et al., ‘Effects of procees parameters on YBCO bulks prepared by TSMTG’, Journal of the Chinese Rare Earth Society, 2003 21(5) 526– 529 (in Chinese). 17. H T Ren, L Xiao, Y L Jiao, et al., ‘Batch progress of single domain YBCO bulk superconductor’, Chinese Journal of Low Temperature Physics, 2003 25(1) 11–15 (in Chinese). 18. H Qian, G S Yuan, ‘Developments of preparation of high Tc superconducting materials’, Chinese Journal of Rare Metals, 1998 22(2) 132–137 (in Chinese). 19. H M Fu, S X Liu, K X Xu, ‘Effect of Ca doping on superconducting performance of YBCO bulk’, Cryogenics and Supercondutivity, 2004 32(2) 56–59 (in Chinese). 20. S Jin, T H Tiefei, R C Sherwood, et al., ‘High critical current density in Y-Ba-CuO superconductors’, Appl. Phys. Lett., 1988 52(24) 2074–2076. 21. H T Ren, L Xiao, Y L Jiao, et al., ‘Large domain YBCO bulk superconductor’, Cryogenics and Superconductivity, 2000 28(3) 28–33 (in Chinese). 22. L Xiao, H T Ren, Y L Jiao, et al., ‘Effects of cooling rate on single domain growth and the superconducting properties for YBCO bulk’, Physica C, 2003 386 262–265. 23. M H Zheng, L Xiao, H T Ren, et al., ‘Study of oxygenation process during the preparation of single domain YBCO bulk superconductors’, Physica C, 2003 386 258–261. 24. R Tournier, E Beaugnon, O Belmont, et al., ‘Processing of large Y1Ba2Cu3O7–X single domain for current-limiting applications’, Supercond. Sci. Technol., 2000 13 886–895. 25. C J Kim, H J Kim, Y A Lee, et al., ‘Multiseeding with (100)/(100) grain junctions in top-seeded melt growth processed YBCO superconductors’, Physica C, 2003 338 205–212. 26. M P Delamare, B Bringmann, Ch Jooss, et al., ‘Influence of the seed distance on the microstructure and the superconducting properties of grain boundaries in a multiseeded melt growth monolith’, Supercond. Sci. Technol., 2002 15 16–22. 27. A Goyal, D P Norton, D J Budai, et al., ‘High critical current density superconduting tapes by epitaxial deposition of YBa2Cu3O7–X thick film on the biaxially textured metals’, Appl. Phys. Lett., 1996 69 1795–1797. 28. X D Wu, S R Foltyn, P N Arendt, et al., ‘Properties of YBa2Cu3O7–X thick film on the flexible buffered metalic substrates’, Appl. Phys. Lett., 1995 67 2397–2399. 29. K Kakimoto, Y Iijima, T Saitoh, ‘Fabrication of Y-123 coated conductors by combination of IBAT and PLD’, Physica C, 2003 392 783–789. 30. B Li, X Gao, W Chen, ‘Weldbility of Bi-based superconductor and flame welding method’, J. Less-Common Met, 1990 164 & 165 660–662. 31. T Kasuga, K Nakamura, E Inukai, et al., ‘Direct joining of BSCCO superconducting
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20 Joining of shape memory alloys K U E N I S H I a n d K F K O B AYA S H I , Osaka University, Japan
20.1
Introduction
Shape memory alloys (SMAs) have become popular as a smart material in the engineering world. However, it has taken a long time for shape memory alloys to be practically applied to various fields since the shape memory effect was discovered. The key to practical application of SMAs lay in the following two factors. One is the discovery of a Ni-Ti based alloy called nithinol with high machinability. Another is the maturity of various manufacturing processes of TiNi-based SMAs, such as rolling, drawing and joining processes. Along with the development of these aspects, SMAs have become a popular smart material for applications including, clothing, glasses, electronic products and medical products. For the further application of SMA alloys to the more advanced products in future, it is essential to develop manufacturing processes in the micro scale. In this chapter, recent research on the microjoining of SMA alloys will be briefly reviewed.
20.2
Basics of shape memory alloys
The shape memory effect (SME) means the phenomenon that materials return to their predetermined shapes by heating. The first report of shape memory transformation that Chang and Read observed was a reversible phase transformation in gold–cadmium (AuCd) in 1932 [1]. After the same phenomenon was reported for the In-Tl [2, 3] alloy, Buechler et al. discovered the SME in nickel–titanium alloy system called nithinol [4]. Essentially nithinol is an alloy containing approximately 50 at.% nickel and 50 at.% titanium. Small compositional changes around this 50:50 ratio make drastic changes in the operating characteristics of the alloy. There are two crystal phases in SMAs: the stronger austenite phase, stable in high temperature, and the weaker martensite phase, stable in low temperature. These two phases differ in their crystal structures. The austenite has a body-centered cubic crystal structure, while the martensite has a parallelogram structure (which 620 WPNL2204
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is asymmetric), with up to 24 variations. When SMAs in martensite are subjected to external stress, they deform through a so-called detwining mechanism, which transforms different martensite variations to the particular one variation that can accommodate the maximum elongation. Due to its parallelogram structure, the martensite phase is weak and can be easily deformed. The austenite phase has only one possible orientation and shows relatively strong resistance to external stress. Many types of shape memory alloys have been discovered. Among them, Nithinol possesses superior thermomechanical and thermoelectrical properties and is the most commonly used SMA [5]. Nithinol SMAs have two unique properties, shape memory effect and pseudo-elasticity. The pseudo-elasticity refers to the phenomenon that SMAs can undergo a large amount of inelastic deformations and recover their shapes after unloading. Figure 20.1 shows the typical stress/strain curve when nithinol was subjected to the loading and unloading test. Pseudo-elastic nithinol exhibits comparatively large yet fully recoverable strains. If processed correctly the nithinol may exhibit recoverable strains as high as 8.0%. Although the term ‘elastic’ is used to describe the effect, it is in fact not elastic in the true metallurgical definition of the word. Rather it is a martensitic phase transformation that results in spontaneously recoverable macro-deformations. These unique properties are the result of reversible phase transformations of SMAs. The stress over which this transformation occurs is approximately constant and therefore a plateau is observed in the stress/strain curve. Figure 20.2 shows the phase transformation behavior of NT-E4 shape memory alloy confirmed by differential scanning calorimetry (DSC) measurement during heating and cooling. The exothermal peak during cooling and endothermic peak during heating corresponds to martensite and reverse transformation, respectively.
Tensile strength (MPa)
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4% 8%
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20.1 Typical stress/strain curve when nithinol was subjected to the loading and unloading test.
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Endo
Heating
15.0 °C
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Heat flow (arb.)
–9.3 °C
–20
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Heat flow (arb.)
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20.2 DSC curves obtained by (a) heating and (b) cooling of NT-E4 shape memory alloy.
Other than the shape memory effect and pseudo-elasticity, Ni-Ti based SMAs possess many desirable properties such as high strength to weight ratio, high damping capacity, high power density, solid state actuation, durability, fatigue resistance, good chemical resistance and biocompatibility.
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Application of SMAs
The above mentioned properties of nithinol have attracted much attention as smart (or intelligent) and functional materials. Nowadays, nithinol has become the most commonly used SMA. Equiatomic Ti50Ni50 shape memory alloy (SMA) is well known for its superior shape memory effect and pseudoelasticity. Early commercialization activities were fueled by applications such as rivets, heat engines, couplings, circuit breakers and automobile actuators. [6] Smart systems for civil structures are described as systems that can automatically adjust structural characteristics in response to external disturbances and/or unexpected severe loading toward structural safety, extension of the structure’s life time, and serviceability [7]. One key technology toward this goal is the development and implementation of smart materials, which can be integrated into structures and provide functions such as sensing, actuation and information processes essential to monitoring, self-adapting and healing of structures. Compared with other smart materials such as piezoceramics, magneto-rheological (MR) fluids, or electrorheological (ER) fluids, SMAs possess excellent properties. Though most of the research activities on SMAs’ applications in civil structures are still at the laboratory stage, a few have been implemented for field applications and have been found to be effective [8]. The vibration suppression of civil structures to external dynamic loading can be pursued by using active control, semiactive control, and passive control. The passive structural control using SMAs takes advantage of the SMA’s damping property to reduce the response and consequent plastic deformation of the structures subjected to severe loadings via two mechanisms, the ground isolation system and the energy dissipation system [9]. The reported SMA isolation systems include SMA bars for highway bridges [10], SMA wire re-centering devices for civil buildings [11]. The SMA energy dissipation devices have been seen in the forms of braces for framed structures [12], dampers for cable-stayed bridges [13] or simply supported bridges [14], connection elements for columns [15] and retrofitting devices for historic buildings. The unique mechanical behavior of nithinol and its apparent biocompatibility has resulted in a number of interesting and often unique medical applications [16–19]. In the field of medical applications, particularly implants, attention has turned towards pseudo-elasticity instead of the more complicated shape memory effect. Slightly nickel rich alloys result in the effect known as ‘pseudo-elasticity’ and it is this phenomenon that is utilized in the vast majority of medical applications. A recent trend in the medical industry has been a drive towards less and less invasive medical procedures. This in turn has created a demand for new medical devices that cannot be met with conventional materials. Ti-Ni SMA wires were widely used for medical devices, such as stents, catheter guide wires, or coil anchors, owing to their
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superior machinability, corrosion resistance in a biological environment and unique pseudo-elasticity. The total global market for medical devices was more than US$ 130 billion in 2002 [20]. More recently, thin film SMA has been recognized as a promising and highperformance material in the field of micro-electro-mechanical system (MEMS) applications, since it can be patterned with standard lithography techniques and fabricated in batch processes [21]. Figure 20.3 shows an example of SMA alloys applied to MEMS devices; Si micromirror structure with the arms fabricated with a TiNi/Si beam structure [22]. Thin film SMA has only a small amount of thermal mass to heat or cool, thus the cycle (response) time can be reduced substantially and the speed of operation may be increased significantly. The work output per volume of thin film SMA exceeds that of other micro-actuation mechanisms. Application of SMA films in MEMS also facilitate simplification of mechanisms with flexibility of design and creation of clean, friction free and non-vibration movement. The main advantages of TiNi thin films include high power density, large displacement and actuation force, low operation voltage, etc. Since TiNi films can provide large forces for actuation and large displacement, most applications of TiNi films in MEMS are focused on micro-actuators, such as micropumps, microvalves, microgrippers, springs, microspacers, micropositioners, and microrappers, micro-sensors, microswitches and microrelays, etc.
20.4
Background on welding and laser processing of SMA
The maturity of various manufacturing processes plays an important role in the applications of TiNi-based SMAs, and rolling, drawing as well as joining
Si
TiNi
20.3 Example of SMA alloys applied to MEMS devices; Si micromirror structure with the arms fabricated with TiNi/Si beam structure.22
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have been extensively studied [23–28]. About welding of SMAs, too, many researches were performed by using arc, laser or electron beam welding. It is the first report that Buechler et al. investigated the TIG arc weldability of nithinol in 1961 [29]. They successfully prepared the sound joints of 1/8 inch thin plate without any porosities or cracks. Although the fusion zone exhibited a fine dendritic structure, formation of some inclusions induced by the existence of H, O or N was observed especially near the surface. From these results, the importance for the control of welding atmosphere was confirmed in order to form sound joints with high mechanical properties. In response to the results, Jackson et al. successfully welded the nithinol by spot welding in an Ar atmosphere [30]. Since this research was performed before the shape memory effect was discovered in the Ni-Ti system, functional properties such as shape memory properties of SMA joints were not investigated. Nishikawa systematically reported the shape memory properties of SMA joints [31]. Since the fusion zone formed during the conventional welding process drastically reduces the shape memory effect of joints due to inclusions or other intermetallic phase formation, he pointed out the importance of developing new welding processes where the fusion zone hardly leaves traces in the joints. In this sence, there has been much research to investigate the applicability of resistance welding to the joining of SMAs[32]. Yasunori et al. reported that Ti-50.9at%Ni alloy wires with a diameter of 1 mm were resistance welded to exhibit fracture joint stress more than 1GPa [33]. The friction stir welding process is considered to be one of the promising newer joining processes of SMAs [34]. Even for the liquid phase joining process, Hirose et al. successfully reported the applicability of the laser welding process by minimizing and controling the fusion zone [35, 36]. Thus the welding of SMAs has made progress with the development of welding process such as laser or electron beam welding. Meanwhile, the products to which SMAs are applied have recently been changing. The most interesting emerging products are medical devices. TiNi shape memory alloy wires are widely used for medical devices, such as stents, catheter guide wires, or coil anchors, owing to their superior machinability, corrosion resistance in a biological environment and unique pseudo-elasticity. These micro medical devices have played an important role in the recent progress on medical treatments. The recent demands for these medical devices lie in the development of finer and more complicated devices. For instance, stents and coil anchors, which are implanted into blood vessels, were required to be applied to finer or more complicated parts of the human body. Most stents today are 316L stainless steel, and are expanded against the vessel wall by plastic deformation caused by the inflation of a balloon placed inside the stent. Nithinol stents, on the other hand, are self-expanding – they are shape-set to the open configuration, compressed into a catheter, then
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pushed out of the catheter and allowed to expand against the vessel wall. Nithinol stents may be constructed in a number of different ways. They may be made from wire, coiled strip, laser cut tubing or laser cut or etched sheet. The majority of stents are made from laser cut tubing, which is difficult to fabricate finer or more complicated medical devices. Consequently, establishment of micro welding technology is considered to be one of the available methods for the fabrication of micromedical devices. Apart from medical devices, the development of microjoining processes is required for the fabrication of MEMs devices. To respond to the miniaturization and density growth of electronics products, establishment of the fabrication process for MEMs is indispensable. In order to achieve the welding of SMAs on a micro scale, development of welding materials as well as welding processes based on the design of the interfacial structure is essential. The motivation for development of dissimilar welding techniques with other materials has been increasing. The stiffness of TiNi SMA wires is small compared to stainless steel, one of the conventional biomaterials. When orthodontic wires and arches are made from monolithic nithinol wire with low stiffness, loss of anchorage occurs. In contrast, the use of stiff stainless wires gives excessive orthodontic forces, which will prevent the treated teeth from moving caused by the absorption of alveolar bone [37]. If TiNi SMA and SS wires are bonded together and used in orthodontic treatment, stainless and nithinol wires can be used as anchorage parts and treated parts respectively, which would greatly shorten the period of orthodontic treatment and improve the quality of orthodontic treatment. This situation is similar to the case of guide wires, of which handling and tip parts are required to be made from stiff stainless wire and from elastic nithinol wires, respectively.
20.5
Brazing of nithinol
Brazing is an economic fabrication method for complex assemblies compared to other joining methods [38, 14]. Since shape memory and pseudo-elasticity have a strong relationship with the microstructure, such as composing phases or grain size, the microstructural evolution caused by joining processes have to be minimized in order to form SMA joints without degradation of their characteristics. Differently from the welding process, the base metal is essentially not melted during the brazing process. For the brazing of nithinol, the selection of filler metal plays a crucial role. The wettability of the braze alloy, the mechanical properties of the brazed joint and the formation of intermetallic phases at the joint must all be considered in the selection of the filler metal. The reaction of Ti with other elements at elevated temperatures may reduce many featured properties of
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TiNi-based SMAs. Consequently, the investigation of brazing TiNi SMA is very limited. The application of Ag-Cu eutectic braze alloy (BAg-8), which is well known for its use in brazing copper alloys, stainless steels and nickel-base alloys, is reported to be effective for the brazing of SMAs [39]. Since the melting temperature of BAg-8 is 1053 K, brazing is performed at about 1153 K. The brazed joint microstructure can be divided into that of filler metal and interfacial reaction layer. The filler metal is mainly comprised of Ag-rich and Cu-rich parts. The intermetallic phase composing interfacial reaction layer changes from Ti(Ni, Cu) to TiNiCu, Cu2Ti with increasing the brazing temperature and time. During brazing at about 1173 K, mainly reactive Ti from nithinol base metal dissolves into Ag-Cu fillers. In the ternary Ag-CuTi alloy system, there is a large miscibility gap in the liquid state and molten filler separates into two liquids. One is rich in Ag with almost no solubility in NiTi [40], and the other is rich in Cu with rather high solubility with NiTi [41]. The Ag-rich liquid does not react with Ti50Ni50 substrate. In contrast, the copper content is depleted from the matrix of brazed joint due to the formation of interfacial CuNiTi or TiCu2 phase. Although the interfacial TiCu2 layer is less detrimental to shape memory behavior than the interfacial CuNiTi layer, both CuNiTi and TiCu2 interfacial layers are detrimental to the bonding strength of brazed Ti50Ni50 joints. Shear strength of brazed joints exceeds 200 MPa for both braze alloys if the brazing time exceeds 180 s. However, thick interfacial CuNiTi and TiCu2 layers can reduce the shear strength. Hence, it is highly preferable for brazing that minimum heating is introduced during brazing in order to avoid deterioration of the base metal as well as to avoid the growth of the reaction layer. It is reported that the wettability of Ag-Cu eutectic braze can be significantly improved with the addition of 1–5 wt% titanium [12, 13, 42, 43]. Accordingly, it is possible to decrease the brazing temperature and time by using the Ti added Ag-Cu based braze alloys. It is also reported that addition of Ni to AgCu eutectic alloy enhances the reactivity between filler and base metal, resulting in the increased tensile strength [44]. Although there is no melting of the base metal during brazing, heating can still affect the properties of the materials being joined [37]. It will be highly preferred that minimum heating is used during brazing in order to avoid deterioration of the base metal. The rapid infra-red joining technique was originally developed at the University of Cincinnati for high temperature materials [45]. Infra-red brazing is a novel technique with a rapid thermal cycle in comparison to traditional furnace brazing [46]. In order to minimize the thermal effect on base metals during soldering, application of novel brazing processes have also been tried. Compared with
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the traditional furnace brazing, infra-red brazing is advantageous for realizing the rapid heating and cooling process. The rapid heating rate is reported to reach as high as 3000 deg C/min. [47]. As a rapid brazing process that does not heat the base metal, laser brazing has been successfully applied. There are extremely limited research findings about the dissimilar joining of nithinol to other materials. Application of Ag-Cu BAg-8 based alloys as a filler is effective for the similar joining of both nithinol and stainless steel. However, when this alloy is applied to the dissimilar joining of nithinol to stainless steel, a brittle Fe-Ti based intermetallic reaction readily forms at the interface between filler and both materials, resulting in an extremely low joint strength [48]. Namely, since Fe and Ti have a high affinity with each other, both elements quickly diffuses within the filler to form intermetallic compounds. Thus, in order to form sound dissimilar joints, the filler metal should be designed to suppress the diffusion of these elements. The authors have confirmed the effectiveness of Ni addition to the BAg8 filler. Although Ni in filler cannot suppress the diffusion of Fe and Ti completely, Ni plays a role in hindering the dissolution of Ti from the nithinol base metal and the joints obtained exhibited a joint tensile strength of 400 MPa [48]. For the complete suppression of diffusion, application of BAg-8/Cu/BAg-8 layered filler is effective and the joints obtained exhibited 500 MPa, which is the same as that of similar stainless steel joint [49]. Lin et al. reported favorable biocompatibility of dissimilar joints formed by laser brazing using Ag based Ag-Cu-Zn-Sn alloys as a filler [25]. For the dissimilar joining with Ti, diffusion of reactive Ti and the formation of Ti-Cu or Ti2Ni intermetallic reaction layer becomes a problem. Dissimilar joint prepared by using BAg-8 exhibited a joint tensile strength of about 300 MPa, which is 80% of that for Ti base metal. Application of Cu-Ti based amorphous MBF5004 alloy slightly increases the maximum joint, but the optimum conditions are subtle [50].
20.6
Laser microwelding
Recent progress with lasers such as YAG laser or single mode fiber laser has opened the possibility of laser microwelding of shape memory alloys with extremely small heat effects to the base metal [51]. In recent years the laser has been successfully introduced as a suitable joining technique for NiTi wires or tapes. Compared with SUS304 stainless wire, in order to achieve fully penetrated favorable melting, a smaller heat input was required to penetrate the NT-E4 wires than SUS304 due to the higher laser absorption ratio and lower melting temperature of NT-E4 wires. However, the favorable heat input range became less than that for SUS304 wires, especially for the shorter pulse duration. It should be noted that melting of shape memory alloy wires needed more
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precise control of laser conditions although it needed a smaller power input compared with SUS304 wires [52] The melted metal exhibited a B2 structured dendrite microstructure that grew from the fusion boundary to the center of the melted metal in entire laser conditions. Due to the rapid solidification effect by laser irradiation, the grain size was about 5 µm, which is still larger than that of the base metal. In the heat affected region, fine equiaxed grains were slightly enlarged to about 3 µm by heat effects during laser irradiation. The spot melted wires exhibited the tensile strength 900 Mpa, which is about 70% strength of the base metal. While the base metal exhibited a ductile fracture surface after tensile movement, all of the spot melted samples underwent brittle fracture in the melted region. This difference in fracture mode is attributed to the difference in the grain size between base metal and melted region. In order to obtain sound pseudo-elasticity, it is considered that the grain size of B2 phase should be smaller than 1 µm. In this sense, melted region or HAZ region with a larger grain size does not have a pseudo-elastic function and will be a stress riser during tensile testing [53] By the measurement of transformation temperature of spot melted regions by differential scanning calorimetry it was confirmed that the spot melted region still exhibits martensite to B2 transformation behavior at almost the same temperature although the transformation became slightly less sensitive. By loading and unloading 8% of stress to investigate pseudo-elastic behavior, spot melted NT-E4 wires showed typical stress-strain curves suggesting the occurrence of pseudo-elasticity, but the loaded stress did not recover completely because melted area and heat affected area did not show pseudo-elasticity. To exhibit sensitive pseudo-elasticity, it is reported that the grain size of B2 should be less than 1 mm, the degradation of pseudo-elasticity of spot melted wires is considered to be due to the increased grain size in the fusion and heat affected zone. This is also clearly confirmed by the fact that there is a proportional relationship between the width of the melted region and the residual stress for the spot melted NT-E4 wires. Thus it is highly desirable to develop a welding technique to minimize the fusion and heat affected zone of joints. The microstructure of the weld metal was almost the same as that of spot melted samples. When welding was performed by optimized laser conditions, weld joints with a higher joint fracture load than the transformation induced load were obtained. Tuissi et al. reported similar results for the laser microwelding of 0.5 mm thick Ni–49.6at.%Ti plate [54]. In order to investigate the biocompatibility of spot melted NT-E4 wires, samples were immersed in 1.0% lactic acid at 310 K for up to 240 days [55]. Ti and Ni ions from samples dissolves into the lactic acid. By the practical measurement of the ions by ICP, the dissolved Ti and Ni were evaluated to be 0.6 g/m2 and 0.7 g/m2, respectively. Although the concentration of lactic
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acid used in this research was proved to be larger than that in the actual biological environment for the enhancement of difference in corrosion resistance, there were no significant differences in dissolved metallic ions between base metal and spot melted NT-E4 wires. Similarly, there were no significant differences in the anodic polarization curve of the samples dipped in 0.9% NaCl or H2SO4 between base metal and spot melted NT-E4 wires. Thus, it can be concluded that spot melted shape memory alloy NT-E4 wires retains bio-chemical properties comparable with those of base material [55].
20.7
Future trends
SMAs are promising materials for use as smart materials and are now planned to have new functional properties. However, although most of the research activities on the applications of SMAs are still in the laboratory stage, innovations in the manufacturing processes for SMAs will open the way to practical applications. Research on the microjoining of SMAs has just started and data about the reliability of joints such as biocompatibility should be determined. With the development of welding process such as laser welding, development of SMAs themselves that are suitable for welding may be necessary in the future. By reviewing the trends for future products, it is estimated that products manufactured from monolithic SMAs will be proportionally less. In this sense, development of dissimilar joining such as with Fe, Ti or Si will become more important.
20.8 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14.
References
Chang LC, Read TA. Trans AIME 1951; 189: 47. Burkart MW, Read TA. Trans AIME 1953; 197: 1516. Basinski ZS, Christian JW. Acta Metall 1954; 2: 101. Buehler WJ, Gilfrich JW, Wiley RC. J Appl Phys 1963; 34: 1473. Duerig TW et al. Engineering aspects of shape memory alloys. London: ButterworthHeinemann; 1990. Brook GB, Materials and Design, Vol. 4 1983, pp 835–40, 1983. Otani S, Hiraishi H, Midorikawa M. ‘Development of smart systems for building structures’. Proceedings of SPIE 2000; 3988: 2–9. Indirli M. Proceedings of SPIE; 2001; 4330: 262–72. Saadat S, Salichs J, Noori M, Hou Z, Davoodi H, Bar-on I, Suzuki Y, Masuda A.: Smart Materials and Structures 2002; 11: 218–29. Wilde K, Gardoni P, Fujino Y. Engineering Structures 2000; 22: 222–9. Dolce M, Cardone D, Marnetto R. Proceedings of SPIE 2001; 4330: 238–49. Saadat S, Noori M, Davoodi H, Zhou Z, Suzuki Y, Masuda A. Smart Materials and Structures 2001; 10: 695–704. Leon RT, DesRoches R, Ocel J, Hess G. Proceedings of SPIE 2001; 4330: 227–37. Casciati F, Faravelli L, Petrini L. Computer-Aided Civil and Infrastructure Engineering 1998; 13: 433–42.
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15. Tamai H, Kitagawa Y. Computational Materials Science 2002; 25: 218–27. 16. Duerig TW, Melton KN, Stouckel D, Waguian CM (eds). Engineering Aspects of Shape Memory Alloys, Butterworth-Heinemann, Boston, 1990. 17. Speck K, Fraker A, Dent J. Res. 1980; 59: 1590. 18. Duerig T, Pelton A, Stockel D. Materials Science and Engineering A 1999; 273– 275; 149. 19. Morgan NB. Materials Science and Engineering A 2004; 378; 16. 20. Frost and Sullivan, Opportunities Unfold for Medical Devices Companies, frost.com, March 2003. 21. Krulevitch P, Lee AP, Ramsey PB, Trevino JC, Hamilton J, Northrup MA. J. MEMS 1996; 5; 270. 22. Fu Y, Du H, Huang W, Zhang S, Hu M. Sensors and Actuators A 112 (2004) 395–408. 23. Miyazaki S, Imai T, Igo Y, Otsuka K. Metall Trans 1986; 17A: 115. 24. Liu Y, McCormick PG. Acta Metall Mater 1990; 38: 1321. 25. Lin HC, Wu SK. Acta Metall Mater 1994; 42: 1623. 26. Wu SK, Lin HC, Yen YC. Mater Sci Eng 1996; 215A: 113. 27. Gale WF, Guan Y. J Mater Sci 1997; 32: 357. 28. Hsu YT, Wang YR, Wu SK, Chen C. Metall Mater Trans 2001; 32A: 569. 29. Buehler W.J, Wiley RC. United State Naval Ordnance Laboratory, NOLTR 1961; 61–75; 68. 30. Jackson CM, Wagner HJ, Wasilewski RJ. NASA-SP-5110 1972; 22. 31. Nishikawa M. Kinzoku 1983; 53; 36 (Japanese). 32. Hall PC. Proc. Shape Memory and Superelastic Technologies, Pacific Grove, California, 2000, pp 67–81, 2000. 33. Yasunori M, Kim JI, Sakuma T, Murata K, Miyazaki S. Proc JSME annual meeting, 1; 2001; 81–82. 34. London B, Pelton A, Fuller C, Mahoney M. Proc. TMS Annual Meeting, San Francisco, California, 2005; 67–74. 35. Hirose A, Uchihara M, Araki T, Honda K, Kondo M., 5th International Symposium on Advanced Technology in Welding, Material Processing and Evaluation, Tokyo, Japan, 1990; 687–692. 36. Hirose A, Uchihara M, Araki T, Honda K, Kondo M. J. Jpn Inst Met. 54 (1990) 262– 9. (Japanese). 37. Murrel EF, Yen EHK, Johnson RB. Am. J. Orthod. 1996; 110; 280. 38. Schwartz M. Brazing for the engineering technologist. New York: Chapman and Hall; 1995; 1. 39. Yang TY, Shiue RK, Wu SK. Intermetallics 2004; 12: 1285–92. 40. Wu SK, Wayman CM. Unpublished work, University of Illinois at Urbana-Champaign; 1984. 41. Mercier O, Melton KN. Metall Trans 1979; 10A: 387–9. 42. Li R, Pan W, Chen J, Lian J. Mater Sci Eng A 2002; A335: 21–5. 43. Paulasto M, van FJJ, Kivilahti JK. J Alloy Comp 1995; 220: 136–41. 44. Watanabe T. Quarterly Journal of the Japan Welding Society 1992; 10: 55. 45. Blue CA, Lin RY. Joining and adhesion of advanced inorganic materials. MRS Symp Proc 1993; 314: 143. 46. Shiue RK, Wu SK, O JM, Wang JY. Metall Mater Trans A 2000; 31A: 25. 47. Shiue RK, Wu SK, Chen SY. Acta Mater 2003; 51: 1991–2004. 48. Seki M, Yamamoto Y, Uenishi K, Kobayashi KF, Journal of Japan Institute of Metals 2000; 64; 632–40. (Japanese).
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49. Uenishi K, Kobayashi KF, Welding Technology 2004; 6; 75 (Japanese). 50. Kunimasa T, Seki M, Yamamoto H, Nojiri M, Uenishi K, Kobayashi KF, Journal of Society of Materials Science 2001; 50: 1218 (Japanese). 51. McGeough JA. Advanced Methods of Machining, Chapman and Hall, London, 1988. 52. Uenishi K, Seki M, Kunimasa T, Takatsugu M, Kobayashi KF, Ikeda T, Tuboi A. Proc. of 3rd Inter. Sympo. on Laser Precision Microfabrication, 2003; 4830; 57. 53. Ogata Y, Takatsugu M, Kunimasa T, Uenishi K, Kobayashi KF. Materials Transactions 2004; 45; 1070. 54. Tuissi A, Besseghini S, Ranucci T, Squatrito F, Pozzi M. Materials Science and Engineering A 1999; 273–5: 813. 55. Uenishi K, Takatsugu M, Kobayashi KF. Proceedings of Fourth International Symposium on Laser Precision Microfabrication, Munich Germany, SPIE 2004; 5063: 282.
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21 Wafer bonding J W E I and Z S U N , Singapore Institute of Manufacturing Technology, Singapore
21.1
Introduction
Wafer bonding is an important process and plays a critical role in many applications such as microelectromechanical systems (MEMS) packaging. MEMS packaging technologies have been derived from conventional integrated circuits and hybrid electronic packaging. However, it presents a unique and difficult challenge compared with common microelectronic packaging because of diverse applications of MEMS devices and their frequent interactions with the environment. Besides, MEMS packaging can represent a significant portion of the total device cost, and thus developing packaging technologies that satisfy all of the necessary performance requirements, while keeping the costs of assembly at a minimum becomes crucial. Protection of MEMS devices from diverse environmental factors that can affect durability and performance must be taken into account at the earliest stages of packaging design. To meet this challenge, many MEMS manufacturers have developed wafer-level packaging schemes that provide a low-cost, protective environment for the sensor device during all stages of assembly and testing [1–5]. The prevalent use of wafer-level packaging is unique to the manufacture of MEMS components. Wafer bonding technology raises an important concept of wafer-level packaging for MEMS devices, which makes tremendous overall savings in cost possible, since packaging of a multitude of sensors or actuators can be performed simultaneously [6]. Various bonding techniques have been developed, such as Si-to-Si direct bonding, Si-to-glass anodic bonding, and bonding via intermediate layers, such as sodium silicate [7], polymethylmethacrylate [8] and solders, etc. The adhesion of solids to each other has been observed for centuries. In 1792, Desagulier showed that two spheres of lead, when pressed together, strongly adhered to each other [9]. In this case strong plastic deformation is involved, which is not a desirable option for typical, highly brittle semiconductor wafers. The phenomenon that optically polished bulk pieces 633 WPNL2204
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of metal stick to each other was observed, and an analogous phenomenon was also found for optically polished glasses. The first systematic investigation of room temperature adhesion between two optically polished glass plates was performed in 1936 by Lord Rayleigh [10] who was the first to determine that the interaction energy per area was of the order of 100 mJ/m2. The bonding of silicon wafers to sodium-containing glass wafers at elevated temperatures (around 500 °C) under the influence of an applied electric field was pioneered by Wallis and Pommerantz in 1969 [11]. Ever since, it has been used in the area of the encapsulation of sensors. This type of bonding, termed ‘anodic bonding’, is restricted to ionically conducting glasses with a thermal expansion coefficient close to that of the semiconductor wafer to be bonded and the presence of mobile ions, such as sodium, which is not desirable for most advanced microelectronic devices. The transfer of a thin gallium arsenide layer to a glass substrate was first performed by ‘wafer fusion’ at elevated temperatures by Antypas and Edgecumber in 1975 [12]. They used an epitaxial, aluminum-rich AlGaAs etch-stop layer to etch off the GaAs substrate but leave a thin GaAs layer to be transferred. The etch-stop concept became essential for many wafer bonding applications. Widespread interest in modern wafer bonding techniques was generated by reports on silicon-silicon wafer bonding by two groups [13–15] in 1985– 1986 in which the initial bonding of the two wafers at room temperature was followed by a high-temperature bond-strengthening heat treatment. Shimbo and co-workers from Toshiba [13] investigated the bonding of silicon wafers without any thermal oxide between the wafers. Wafer bonding served as a substitute process for the growth of thick epitaxial layers of single-crystalline silicon on silicon for potential applications in power devices. In contrast, Lasky from IBM [14, 15] bonded two silicon wafers, with one or both of the wafers being covered by a thermally grown, electrically insulating, silicon dioxide layer. Parallel to the electronics-related work on silicon wafer bonding at Toshiba and IBM, in 1988 Petersen and Barth [16] pioneered the use of silicon wafer bonding for pressure sensors, in which at least one of the silicon wafers was structured containing holes or cavities. The process was termed ‘fusion bonding’ and actually went into production. However, the thorough investigation of the science behind wafer bonding was actually carried out during the last decade. In addition, a range of other possible applications has opened up; for example, those used in micromechanics [17–19] and in optoelectronics [20– 22] involving the bonding of III-V compound wafers at elevated temperatures in hydrogen are emerging. In order to bond the wafers together, various techniques have been developed. These techniques can be categorized as direct bonding, anodic bonding and intermediate layer bonding. In the following sections, brief reviews of the various wafer bonding
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processes, such as direct wafer bonding, anodic wafer bonding, wafer bonding via intermediate layers and bonding of dissimilar materials, will be given. In each section, bonding mechanisms, process advantages and limitations, and their current and potential applications will be introduced. Furthermore, some of the challenges in wafer bonding will also be highlighted.
21.2
Direct wafer bonding
21.2.1 Background The direct wafer bonding method relies on forces that naturally attract surfaces together when they are very smooth and flat. A range of mechanisms have been proposed to explain this initial contacting attractive force. It is well known that smooth metal surfaces, if atomically clean, will bond together. This process is often referred to as ‘cold welding’ and is typically achieved by cleaning and contacting the metal surfaces in vacuum to maintain cleanliness. This bond usually relies on plastic deformation of the metal to bring the atoms in close contact. In the case of most direct wafer bonding techniques that have been performed for microstructures, surface treatment (e.g., hydration, oxygen plasma exposure) is conducted prior to the contacting to promote the surface attraction and bonding process. This is sometimes assisted by a modest pressure to expel air between the wafers and to initiate the contact. The bond is usually followed by a thermal cycle, which increases the strength of the bond.
21.2.2 Silicon direct bonding process The silicon wafer bonding process consists of three basic steps: surface preparation, contacting, and annealing. The starting wafers must be smooth and flat. There have been studies of the necessary surface quality for wafer bonding [23, 24], and in general, it has been experimentally observed that the wafers should have a roughness of no greater than about 10 Å, and a flatness of less than 5 µm. Also, protrusions from the surface (resulting from previous processing) of greater than 10 Å can produce problems in the bonding, which implies that this bonding technique puts stringent requirements on surface quality. Therefore, all of the process steps should be conducted in a clean room environment, although Gosele has proposed a powerful ‘micro clean room’ concept that does not require a clean room [25]. The surface preparation step involves cleaning the mirror-smooth, flat surfaces of two wafers to form a hydrated surface. There have been studies of the differences between hydrophobic and hydrophilic surfaces on the final bond interface [26]. Following preparation, the wafers are contacted in a clean environment by gently pressing the two surfaces together at one central
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point. The surfaces come into contact at this point and are bound by a surface attraction of the two hydrated surfaces. A contact wave is initiated at this point and sweeps across the wafer surfaces, bringing them into intimate contact over the entire surface. The contacting process is critical to prevent trapping of particulates or air between the surfaces. The exact origin of the attractive force that promotes the contact wave is not universally agreed upon [26] and depends to a certain extent on whether the bond is Si-Si or SiSiO2. The most common assumption is that a bond is formed between –OH groups on the opposing surfaces. The final step in the bonding process is an elevated temperature annealing of the contacted pair at temperatures anywhere from room temperature to 1200 °C. While the room temperature contacted samples are well adhered, this annealing generally increases the bond strength by more than an order of magnitude at a higher temperature (800–1200 °C). Direct wafer bonding has attracted significant attention in the fields of microelectronics and MEMS. Traditional direct wafer bonding usually requires high temperature annealing above 800–1000 °C to achieve strong bonding between wafers. This high temperature treatment will generate many problems such as thermally induced mechanical stress of the bonded materials due to their different thermal expansion coefficients, undesirable changes and reactions in the case of materials and structures which are sensitive to high temperature, and so on [27]. Therefore, it is necessary to find a bonding process that can result in strong bonding at temperatures as low as possible (e.g., below 400 °C). There are many methods to achieve high bond strength at low temperature [28]. Vacuum wafer bonding is one of the promising approaches. Variation of vacuum wafer bonding at low temperature includes ultra-high vacuum (UHV) wafer bonding [29], low vacuum wafer bonding [30] and high vacuum wafer bonding after argon beam etching [31]. Experiments have shown a significant increase in the bond strength of vacuum wafer bonding as compared to that in conventional wafer bonding. Many bonding process parameters have been shown to have significant effects on bond quality, such as cleaning process (cleaning solution ratio, cleaning time and cleaning temperature) [32, 33], annealing temperature, annealing time and so on [34]. However, the effect of applied load was rarely investigated, except that Kissinger et al. found that the applied load can improve bond strength in conventional wafer bonding at low temperature and it had no effect on the bond strength in plasma activated wafer bonding [35]. Their results were later on confirmed by Takagi et al. [31]. Furthermore, Si wafers can also be bonded in vacuum at low temperatures around 400 °C [36, 37]. Based on the models of conventional direct wafer bonding proposed by Stengl et al. [38] and Q.-Y. Tong et al. [34], bonding sites between two
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wafers play a very important role in enhancing bond strength. More bonding sites will result in higher bond strength. When conventional wafer bonding is performed at low temperature (<600 °C), a lack of bonding sites causes incomplete bonding, and an applied load can push more bonding sites to be connected between two wafer surfaces. Therefore, a high applied load yields high bond strength, as shown in Fig. 21.1 by the lower curve [35]. If, on the other hand, the bonding surfaces are fully hydrated through treatment such as the plasma activated wafer bonding, the distance between the two wafers is closed by van der Waals forces between bonding sites on two surfaces until the entire wafer area is bonded. Therefore, the wafers are able to overcome most of the local waviness and an applied load is ineffective as shown by the work presented in refs [31, 35]. The middle straight line in Fig. 21.1 shows this effect experimentally. However, the study on low temperature vacuum wafer bonding shows that the applied load does have an effect on the bond strength as shown in Fig. 21.1 (the upper curve) [36]. The higher the applied load, the lower the bond strength, and the bonded pairs have highest bond strength when no load is applied. In all cases of vacuum wafer bonding, the bond strengths are above 19 MPa, which is high enough for practical applications, such as dicing and device fabrication. Moreover, in the case of zero applied load, the cracking occurs inside the silicon wafer during a pull test rather than at the bonded interface. The fracture surface is shown in Fig. 21.2. This means that the bond strength at the interface is higher than the silicon per se, and the good bonding quality is also proven by the TEM observation, as shown in 30
Conventional [35] Low temperature [36] Plasma activated [31]
Bond strength (MPa)
25
20
15
10
5 0.0
0.5
1.0 1.5 2.0 Applied load (kN)
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3.0
21.1 Comparison of bond strength under different loads from different studies; data from [31, 35, 36].
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10 mm × 10
(a)
100 µm
× 40 (b)
21.2 Optical micrographs of fracture surface of bonded pair at bonding temperature of 400 °C, applied load 0 kN, annealing time of 2 hours and vacuum of 10–3 mbar. (a) Fracture surface on the SiO2 wafer (b) zoomed picture of the dashed area in (a) (white area Si and black area SiO2); adapted from [36].
Fig. 21.3. It can be seen from the TEM image that perfect bonding has been formed between the two wafers.
21.2.3 Direct wafer bonding mechanism For direct wafer bonding, the wafer surfaces are normally treated to make them hydrophilic and covered with water molecules. The bonding at room
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21.3 TEM micrograph of bonded pair at bonding temperature of 400 °C, applied load of 0 kN, annealing time of 2 hours and vacuum of 10–3 mbar; adapted from [36].
temperature is caused by hydrogen bonds between the chemisorbed water molecules located on the opposing wafer surfaces. After stacking, the –OH starts to dehydrate at a bonding temperature around 200 °C, the hydroxyl groups on two wafer surfaces form a strong electrical field which attracts two wafers to each other [38, 39]. To break the hydrogen bonds between two wafer surfaces to allow silanol groups to react with each other and subsequently form stronger Si-O-Si bonds, high temperature annealing is always needed to enhance the diffusion of absorbed and produced water [34] as seen in Fig. 21.4. For low temperature vacuum direct wafer bonding [36, 37], the bonding mechanism could be different from the conventional and plasma treated low temperature bonding. They are composed of two different dominant mechanisms, namely the electric field model in low temperature wafer bonding and the enhanced diffusion effect of impurities in vacuum. It is known that a vacuum acts as a powerful pump which can suck the trapped gas, including water, out of the interface [40]. The reduction of trapped air and water at the bonding interface under vacuum significantly increases the bonding sites and allows covalent bond (siloxane) to be developed through the polymerization of silanol bonds on the bonding surfaces, i.e., through the reverse reaction of eqn 21.1. Si-O-Si + H2O ⇔ 2Si-OH
(1)
Due to the nano-roughness of the surfaces before bonding, free volume is enclosed at the interface, and thus nano-sized voids are produced in the interface layer, which was observed by Reznicek et al. [41]. Therefore, impurities can be trapped in these nano-sized voids in the interlayer, which results in lower bond strength. The presence of water molecules renders less siloxane to be produced, thus lower bond strength results. This is in contrast to traditional high
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Si 1.63Å
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O H
H 2.76Å 11.54Å 2.76Å
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Si
Si
Si
Si
O
O
O
O
Si
Si O
O
O
Si
Si O
O
(c)
O
O
3.18Å
Si O
O
O
O
(d)
21.4 Proposed bonding model for hydrophilic Si wafer bonding at different temperatures. (a) Room temperature –110 °C: formation of stable hydrogen bonding between molecular water across the bonding interface. (b) 110–150 °C: removal of interface molecular water and formation of siloxane bonds. (c) 150–800 °C: stable siloxane bonds. (d) >800 °C: viscous flow of interface oxide; adapted from [34].
temperature wafer bonding as a reaction (eqn 21.1) is no longer reversible at T > 425 °C. Thus, the presence of water, if any, will not affect the production of the siloxane, and also water can be easily diffused into bulk and react with Si via the reaction shown below
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641
(2)
For a low temperature wafer bonding model [42] based on the oxidation model of Cabrera-Mott [43], when the applied load results in trapped gas or water, these impurities will block ionized oxygen or hydroxyl groups to sweep to the opposite surface to react and form more siloxane groups, resulting in low bond strength.
21.3
Anodic wafer bonding
21.3.1 Introduction Anodic bonding was invented by Wallis and Pomerantz in 1969, and is the most common wafer-to-wafer bonding technique used in the fabrication of silicon sensors [44]. This technique involves bonding a silicon wafer to a glass wafer at an elevated temperature, using the assistance of an electrostatic field. Hermetically sealed sensors and actuators can be packaged using this method [45]. In parallel with its application in sealing Si and glass-based microfluidic devices such as miniaturized chemical reactors [46], ‘Genechip’ for DNA application and analysis was also produced [47]. Compared with other bonding techniques, which include direct, and intermediate layer bonding, an anodic bonding process with the assistance of an electric field offers the advantage of high bond strength and low operating temperature. Anodic bonding is widely used for bonding glass wafer to other conductive materials due to the good bond quality. It serves for the establishment of a hermetic and mechanically solid connection between glass and metal wafers or a connection between glass and semiconductor-wafers [48–51]. In anodic bonding, the substrates are heated to a typical temperature of 400–450 °C. A typical voltage of 400 V to 1200 V is applied to the wafer pair to be bonded. Electrostatic force and the migration of ions lead finally to an irreversible chemical bond at the boundary layer between the individual wafers. By reducing the bonding temperature, degradation or damage to prefabricated devices and integrated circuitry can be avoided. It can also minimize or eliminate bonding-induced stress problems and warping after cooling. The stress problem often causes failure in reliability. In addition, the wafers with a large difference in thermal expansion coefficient may also be bonded. Currently, the mainstream of wafer bonding focuses on achieving strong wafer bonding at temperatures as low as possible. However, the bond quality is reduced with decreasing the bonding temperature. The bond strength is low, and the bubbles or cavities are difficult to minimize or eliminate at the interface. Hardly any literature reports successful low temperature anodic bonding with high bond strength and a bubble-free interface.
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21.3.2 Low temperature wafer anodic bonding Recently, Si-glass wafer anodic bonding processes were successfully developed by Wei et al. [52, 53]. In their techniques, the bonding temperature ranges from 200 °C to 300 °C, and the voltage used is from 200 to 400 V. Siliconto-glass bonding takes place at once when the external voltage is applied. The current at the very beginning of the bonding process is high but starts to decay rapidly within a few seconds especially for high bonding temperature, as shown in Fig. 21.5. The initial current peak corresponds to the initial transport of sodium and potassium ions from the glass wafer to the cathode, which are consequently neutralized at the cathode. Once the equilibrium state is established, the current drops rapidly. The maximum current is limited to 30 mA in all experiments to protect the electronic device of the bonding system. A high temperature results in high ion mobility in a glass wafer. As the temperature is raised, the resistivity of the glass decreases exponentially and a rapid build up of the space charge occurs at the interface, giving rise to electrostatic forces which pull the two wafers into intimate contact. Therefore, high temperature causes high current, and consequently high bond quality. At bonding temperatures lower than 250 °C, the transition period required to reach the equilibrium state increases with increasing temperature. At bonding temperature higher than 250 °C, the transition period required to reach the equilibrium state decreases with increasing the temperature. At high temperature, more ions decomposed from Na2O and K2O migrate to the cathode at the very beginning. Therefore, the equilibrium state is easily obtained and the transition period is shorter. The bond interface integrity is evaluated by comparing the bonded area 35 Si-G-200 Si-G-225 Si-G-250 Si-G-275 Si-G-300
Current (mA)
30 25 20 15 10 5 0 0
200
400
600 Time (s)
800
1000
1200
21.5 The current-time relationship under different temperatures at voltage of 600 V, force of 200 N, and vacuum of 1 Pa; adapted from [52].
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with the whole wafer size. Figure 21.6 shows the SAM micrographs for the bonded wafers under different temperatures at 600 V, bonding force of 200 N, bonding time of 10 minutes and vacuum of 1 Pa. At a bonding temperature of 200 °C, one large bubble forms in the position 10 mm away from the center of the wafer. For the bonding tooling, the top round electrode consists of two pieces of graphite, an inner round electrode of about 20 mm diameter and an outer round electrode of about 100 mm diameter. These two pieces are connected with a metal strip. When the two wafers are brought into contact at the central point of the wafer, the small inner electrode is first pressed down, followed by the large outer electrode. Therefore, a large bubble is caused by such electrode configuration. The size of all other bubbles is less than 2 mm. At bonding temperature higher than 225 °C, the bubble size is reduced obviously, which is less than 1 mm, and the unbonded area also decreases. A higher bonding temperature can compensate for the drawback of the electrode configuration. Almost bubble-free interfaces have been observed with SAM whose resolution is 2.5 µm. Most of the wafer area has been bonded, which proves the feasibility of anodic bonding at low temperature from the bond efficiency point of view, particularly in the range of 200 °C to 300 °C. The unbonded area or the
(a)
(b)
(d)
(c)
(e)
21.6 Scanning acoustic micrographs of the silicon-to-glass bonded pairs at different temperatures. (a) 200 °C; (b) 225 °C; (c) 250 °C; (d) 275 °C; and (e) 300 °C, at voltage of 600 V, bonding force of 200 N, bonding time of 10 minutes and vacuum of 1 Pa; adapted from [52].
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bubbles are mainly due to the gas entrapped between the mating surfaces of two wafers, and no bubble due to particles can be found, which indicates that the cleaning procedure has effectively removed all the particles on the wafer surfaces. The SAM images are further analyzed using image analysis. The unbonded area is measured and plotted in Fig. 21.7. The unbonded area decreases from 1.22% to 0.6% when the bonding temperature increases from 200 °C to 225 °C. When the bonding temperature is increased to more than 250 °C, the unbonded area is less than 0.3%. Interface integrity is also observed by SEM. Figure 21.8 shows typical cross-sections of bonded silicon and glass pairs at bonding temperatures of 200 °C (a) and 300 °C (b), voltage of 600 V, bonding force of 200 N, bonding time of 10 minutes and vacuum of 1 Pa. No gap can be observed in the bonded interfaces. Bond strength is a very important factor related to bond quality and reliability. High bond strength indicates that a good bond has been formed. It is interesting that the bonded pairs at a higher bonding temperature cannot be separated, and they fracture in the glass wafer but not at the interface. Figure 21.9 shows the bond strength versus bonding temperature. The bond strength increases with increasing bonding temperature. At a bonding temperature of 200 °C, the bond strength of the bonded pairs is 10 MPa. At a bonding temperature higher than 225 °C, fracture occurred in the glass. The strength here is comparable to the bond strength (from 5 to 25 MPa) of the bonded pairs using higher a bonding temperature reported by other researchers [54– 56]. Although the fracture mainly happens inside the glass wafer rather than in the interface, the fracture strength still increases with increasing bonding temperature. It is believed that the glass is annealed during the bonding process, which improves the fracture strength of the glass. The bond strength between silicon and glass is therefore higher than the values plotted in the curve. 1.4
Unbonded area (%)
1.2 1.0 0.8 0.6 0.4 0.2 0.0 150
200
250 300 Temperature (°C)
350
21.7 Unbonded area versus bonding temperature; adapted from [52].
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(a)
(b)
21.8 Typical cross-sections of bonded silicon and glass pairs. (a) 200 °C and (b) 300 °C, voltage of 600 V, bonding force of 200 N, bonding time of 10 minutes and vacuum of 1 Pa; adapted from [52].
At lower temperatures, the reaction in the interface is largely dependent on the temperature since temperature is the major driving force of ion mobility. Besides, the low temperature can significantly hinder the reaction from occurring since less energy can be provided to promote the bonding reaction from the kinetics point of view. As temperature is increased, the dependence of temperature is released, and the electric-field-assisted factor becomes more important, which causes the drift of O2– and OH– to the surface.
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Bond strength (MPa)
50 40 30 20 10 0 150
200 250 300 Bonding temperature (°C)
350
21.9 Bond strength versus bonding temperature; adapted from [52].
The thermal residual stress generated in the low temperature anodic bonding process has been measured via the change in the curvature of the wafers using a Stylus profilometer. As expected, low temperature bonding largely reduces the induced stress. While all the bonded wafers under different conditions are measured, no residual stress is detected.
21.3.3 Anodic bonding mechanisms To determine the element distribution in glass and to avoid any contamination of the interface, the glass wafer of the bonded pair is mechanically and chemically thinned to several microns. Then sputtering starts from the glass layer and penetrates into the silicon wafer gradually. Figure 21.10 shows the element distribution (Si, O, H, Na and K) across the interface between Si and glass at bonding temperatures of 200 °C (Fig. 21.10(a)) and 300 °C (Fig. 21.10(b)), voltage of 600 V, bonding force of 200 N, bonding time of 10 minutes and vacuum level of 1 Pa. The sputtering rate is about 10 nm per minute. The depletion region at the bonding temperature of 200 °C is about 0.5 µm, and it is about 1 µm at the bonding temperature of 300 °C. Furthermore, the depletion region also increases with increasing bonding temperature. High temperature results in high ion mobility. More sodium and potassium ions drift to the cathode from the glass side in the interface and left a larger depletion region. However, the pile up of the ions at the surface of the glass is obvious, which is caused by the attraction by the hydroxyl group on the surface of the silicon wafer. The descriptions of bonding mechanisms reported are controversial and there is still no common agreement on this topic. These descriptions can be mainly categorized into two groups. One is the direct reaction between oxygen and silicon that causes bonding [57–58], while the other attributes bonding to the reaction between hydroxide groups [59].
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Intensity (arb. units)
1.0E+04 H O Na Si K
1.0E+03
1.0E+02
1.0E+01
1.0E+00 0
20
40 60 80 100 120 Sputtering time (minutes) (a)
140 160
Intensity (arb. units)
1.0E+04 H O Na Si K
1.0E+03
1.0E+02
1.0E+01
1.0E+00 0
20
40 60 80 100 120 Sputtering time (minutes) (b)
140 160
21.10 Element distribution in the interfaces of the bonded pairs at bonding temperatures of 200 °C (a) and 300 °C (b), voltage of 600 V, bonding force of 200 N, bonding time of 10 minutes and vacuum of 1 Pa; adapted from [52].
The wafer surfaces are treated to be hydrophilic prior to anodic bonding. Therefore, the wafer surfaces are terminated by hydroxyl groups. When the two wafers are brought into contact, the contact forces are the electrostatic force, the attraction between hydroxyl groups, and the externally applied force. In the anodic bonding process, a space charge region is formed at the glass side of the silicon-glass interface, leaving behind relatively immobile oxygen anions. This in turn creates an equivalent positive charge (image charge) on the silicon side of the silicon-glass interface, resulting in a high electrostatic force between the two wafers. The oxygen anions drift away from the Na+ depletion region to the silicon surface. Silicon is oxidized by the oxygen anions. Higher bonding temperature increased the mobility and
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the drift velocity of both Na+, K+ and oxygen anion as well as the reaction between silicon and oxygen, thus promoting the oxidation process. On the other hand, the glass wafer composed of SiO2, Al2O3, B2O3 and Na2O, and K2O. SiO2 is the major component. For Si-to-Si, Si-to-SiO2 and SiO2-to-SiO2 direct bonding [60–61], the surface terminated hydroxyl groups react by hydrogen bonding even at low temperature. In anodic bonding, Pyrex 7740 glass is used. SiO2 is the major component of such glass. Prior to stacking and bonding, both glass and silicon wafers are treated to be hydrophilic, the surfaces are terminated with hydroxyl groups. The top side is glass wafer and the bottom side is silicon wafer. After stacking, the –OH starts to dehydrate at a bonding temperature around 200 °C, and the hydrogen bonds are replaced by Si-O-Si bonds. Bonding between the hydroxyl groups on the glass and silicon can occur through hydrogen bonding between SiO2 and amorphous Si. In general, anodic bonding mechanisms consist of the oxidation of silicon and hydrogen bonding between hydroxyl groups. Higher temperatures can promote such reactions and result in better bonding quality.
21.4
Wafer bonding via intermediate layers
The applicability of the silicon direct bonding process is restricted since the bonding process involves excessively high process temperatures, typically in the range from 700 to 1100 °C, in order to achieve acceptable bond strengths. The use of thin intermediate layers to decrease the required process temperatures has been investigated.
21.4.1 Bonding via silicides Silicides have found wide application in the development of very large scale integration technology (VLSI), for low resistivity gates, interconnections, and ohmic contacts [62–66]. Recently, they have also been identified as the intermediate materials for wafer bonding because of their low resistivities and high temperature stability. The annealing temperatures for nickel, cobalt and titanium silicides are usually higher than 600 °C to obtain a stable silicide phase [63–66]. It is, however, reported that nickel will react with silicon and form one kind of nickel silicide when the temperature is higher than 280 °C (Ni2Si for 280–350 °C, NiSi for 350–750 °C and NiSi2 for temperatures higher than 750 °C, if the thickness of silicon is much larger than that of nickel), whereas the temperature for cobalt is 350 °C [62]. Ljungberg et al. bonded silicon wafer pairs with cobalt silicide [67], but the annealing temperatures of 700 to 900 °C are too high for wafers with the active readout electronics integrated circuits or with the integrated silicon sensors. Besides, the thermal stress of NiSi fromed from 600 to 750 °C and
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NiSi2 formed from 750 to 900 °C on polysilicon is much smaller than that of cobalt silicides formed from 350 to 1000 °C [62]. Ismail and Bower found that successful bonding occurred when the PtSi surface was rendered hydrophilic by a hot aqua region selective etching and cleaning process, however the annealing temperature is 700 °C [68]. In summary, Ni may be the suitable material to bond silicon to other materials at low temperatures through the formation of nickel silicide.
21.4.2 Bonding via silicates A diluted solution of sodium silicate, aluminum phosphate or ammoniasilica solutions in water was also reported as the intermediate layer [69, 70]. It was spun onto one of the two surfaces and the two wafers were brought into contact. The attraction force of the hydrophilic surfaces results in very close contact. The annealing temperature in the range of 200 and 350 °C for the silicates and aluminum phosphate, respectively, leads to maximum bond strengths. The surface energy of approximately 3 J/m2 can be obtained, which requires annealing temperature of 1000 °C by conventional silicon direct bonding [71]. On the contrary, Puers et al. reported that bonding performed at 250 °C using sodium silicate solution did not withstand the dicing of the wafers, whereas the samples bonded at 150 °C could not be separated by inserting a blade between the two wafers. Using ammoniasilica solution achieves less bond strength, however, it offers a full ICcompatible bonding process.
21.4.3 Bonding via solder Intermediate bonding by using Au-Si eutectic bonding has been proposed as another potential way to reduce bonding temperature. It involved a necessary process temperature of 500 °C that is higher than the Au-Si eutectic temperature of 363 °C [72]. On the other hand, an Au-Au4Al-Si intermediate bonding at 350 °C was reported for packaging infra-red microsensors successfully [73]. The same group also made the sealed cavity of capacitive pressure sensors using Au/Sn solder to bond Ni/Au metalized silicon and Sn/Pb solder to bond Ni/Au metalized silicon, while the respective treatment temperature were 300 °C and 250 °C for fluxless soldering in a vacuum oven with infrared light [74]. To further reduce the temperature, a novel low-temperature wafer bonding with process temperature lower than 160 °C was reported using In-Sn alloy to form the interface [75]. The helium leak test of 6 × 10–9 Torr l/s, and a tensile strength of 200 kg/cm2 could be passed using this technique. However, the flux used in conventional solder to improve wetting capability [76] may cause serious contamination problems for MEMS package. To
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advance the current wafer bonding processes, new bonding technology should be developed for low thermal budget, wafer-level processing, and insensitivity to surface roughness of the device substrate. Chiao et al. first introduced the rapid thermal processing as a bonding method for aluminum bonding [77], and it has been demonstrated that 100% survival rate for the encapsulated cavities to pass both the IPA and autoclave tests can be achieved when the aluminum solder is 4.5 µm thick and 150 µm wide.
21.4.4 Bonding via glass layer Processes for anodic bonding of silicon to silicon using thin glass layers have been reported [78–80]. Typically, Pyrex 7740 or Schott 8330 (Tempax) is used. The thermal expansion of these glasses is closely matched to the thermal expansion of silicon, resulting in low stress in the bonded devices. The thin glass films are usually deposited by sputtering with maximum deposition rates of 0.01 to 0.08 nm/s [79, 81]. For successful anodic bonding the thickness of the glass film should be more than 0.5 µm [79] and a thickness of a few micrometers is desirable to increase the yield and quality of the bond. This requires sputter deposition times of many hours which may be a drawback in volume production. The high stress generated in the deposited layer also limits the thickness of films. Alkali-containing borosilicate glasses, deposited at high rates by electron-beam evaporation, are commonly used in industrial applications. The use of evaporated glass films for anodic bonding purposes at 450 °C has been investigated [82, 83], and the deposition rates of evaporated glass are at least three orders of magnitude higher than those of sputter-deposited glass and therefore very attractive for volume production. However, it requires a relative large film thickness due to a large shift in the glass compositions (loss of sodium). Both methods have the disadvantage that the produced glass layers have a significant roughness. Since the quality of the glass layer is a very critical point for the anodic bonding process, very often the deposited glass layer has to be treated in additional processes. Either the glass film has to be annealed at temperatures over 540 °C to improve the electrical breakdown behavior [84] or even the surface of the glass layer has to be polished to guarantee high strength for the bonds. Wafer bonding by sol-gel intermediate layer has attracted much attention because sol-gel processing can provide active, uniform and homogeneous coatings on wafer surfaces in a simple way [85]. Spin-on glasses reported by Deng et al. offer a very economic and simple way for the fabrication of thin glass layers [85, 86]. In this process, a liquid sol solution is used within a spin coating process. The solution is a mixture based on silica sol, organic silicon containing compounds, and a sodium salt all dissolved in ethanol. The spin-on glass films show very smooth surfaces which improves the bonding quality. The bond efficiency, which refers to the percent of bonded
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area to the entire wafer area, is in the range of 40–90%, and the bond strength varies in the range of 8–35 MPa. At a bonding temperature of 100 °C, the bond efficiency is higher than 80% and bond strength is up to 35 MPa. However, when bonding temperature increases to 300 or 500 °C, the bond efficiency and bond strength decrease sharply to 40–60% and 8–21 MPa, respectively.
21.4.5 Glass-to-glass anodic bonding via silicon layer Glass-to-glass bonding has been found attractive for applications in liquid manipulation, on-chip chemical analysis systems [87], and field emission display [87]. Particularly in micromachined analysis systems for online monitoring of bioprocesses such as cell cultures, due to the high voltage applied on the liquid during the capillary electrophoresis separation [88], a completely electrically insulated channel has to be fabricated. Glass-to-glass bonding is a very suitable solution to such a problem. However, the glass-toglass processes published to date require either high temperature [89], which would damage aluminum patterns, or spin-on glass [90] as an intermediate layer, which would fill recesses, disturbing the channel definition. Therefore, the glass-to-glass anodic bonding process becomes the most suitable candidate. As glass-to-glass bonding is only possible with an intermediate layer, thin films of materials for conventional IC processes can be applied for the bonding, such as amorphous silicon, polysilicon, silicon nitride, silicon carbide, etc. Wei et al. have demonstrated the glass-to-glass bonding using amorphous silicon thin films, and successful bonds were achieved [91–92]. Similar to Si-glass anodic bonding [52, 53], glass-glass anodic bonding takes place immediately on the application of the external voltage. At the start of the bonding process, the value of the current is high but falls rapidly, especially for high bonding temperatures. The initial peak current corresponds to the initial transport of sodium and potassium ions from the glass wafer to the cathode. These ions are consequently neutralized at the cathode. Once an equilibrium state is established, the current remains at a low value. High temperatures generate high ion mobility in the glass substrate. As the temperature is raised, the conductivity of glass increases exponentially. A rapid build up of the space charge occurs at the interface, giving rise to electrostatic forces, which brings the two wafers into intimate contact. Therefore, high temperatures cause high current (Fig. 21.11(a)) and result in good bond quality. At high temperatures, more ions decompose from Na2O and K2O to migrate to the cathode. The equilibrium state is easily obtained and the transition period is shorter. A higher applied voltage is expected to increase the mobility of the Na+ ions. A higher voltage produces a higher electric field. The higher electric field will increase the drift velocity of the sodium ions. Higher voltages will
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Current (mA)
20 15
10 5
0 0
200
400
600 Time (s) (a)
800
1000
1200
12 G-G-200V G-G-400V G-G-600V G-G-800V G-G-1000V
Current (mA)
10 8 6 4 2 0 0
200
400
600 Time (s) (b)
800
1000
1200
21.11 The current-time relationship (a) under different temperatures at voltage of 600 V and (b) under different voltages at temperature of 250 °C. Force is 200 N and vacuum is 1 Pa; adapted from [92].
also accelerate the detachment of the sodium ions from the lattice matrix, and contribute to the concentration of free sodium ions. This is the likely reason for the higher current and the longer time required to establish the equilibrium state (Fig. 21.11(b)). Therefore, a higher applied voltage can generate more free sodium ions, and subsequently, contributes to a larger electrostatic force. The bonding interface integrity was evaluated by comparing the bonded area with the entire wafer surface. It was characterized using SAM. Figure 21.12 shows the SAM micrographs for the bonded wafers under different temperatures and voltages at bonding force of 200 N, bonding time of 10 minutes and vacuum of 1 Pa. At low bonding temperatures (say, 200 °C),
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(b)
(c)
(d)
(e)
(f)
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21.12 Scanning acoustic micrographs of the glass-to-glass bonded pairs under different temperatures: (a) 200 °C, (b) 250 °C and (c) 300 °C at voltage of 600 V and under different voltages: (d) 200 V, (e) 600 V and (f) 1000 V at a temperature of 250 °C. Bonding force of 200 N, bonding time of 10 minutes and vacuum of 1 Pa; adapted from [92].
more and larger voids are found at the interface. With increased bonding temperature, the number of voids and the void size decrease noticeably. When the bonding temperature is higher than 275 °C and the voltage is higher than 800 volts, no voids are found. The bonding temperature and the voltage have a significant influence on the voids. The unbonded area or voids are mainly attributed to gas entrapment between the mating surfaces of the two wafers. No voids arise from particle contamination, thus indicating that the cleaning procedure is effective. The SAM images were further analyzed with the image analysis method. The unbonded area was measured and plotted as in Fig. 21.13. It can be seen that the unbonded area decreases from 2.2% to 0.25% when the bonding temperature increases from 200 °C to 250 °C. With a further increase of the bonding temperature to more than 275 °C, the whole wafer area is bonded together. The unbonded area decreases from 1.65% to 0.02% when the voltage increases from 200 V to 1000 V. The interface integrity was also observed by SEM. No observable gap can be seen from the SEM micrographs, as shown in Fig. 21.14. Bond strength is an important factor for bond quality and reliability. High bond strength indicates that a good bond has been formed. Figure 21.15 shows the bond strength versus bonding temperature and voltage. The tensile
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Unbonded area (%)
2.5 2.0 1.5 1.0 0.5 0.0 150
200
250 300 Temperature (°C) (a)
350
Unbonded area (%)
2.5 2.0 1.5 1.0 0.5 0.0 0
200
400
600 800 Bias (–V) (b)
1000
1200
21.13 Unbonded area (a) under different temperatures at voltage of 600 V and (b) under different voltages at temperature of 250 °C (b). Bonding force of 200 N, bonding time of 10 minutes and vacuum of 1 Pa; adapted from [92].
strength of the bonded pairs is higher than 10 MPa for all the bonding conditions. The bond strength increases with an increase in the bonding temperature. Figure 21.16 shows the typical fracture surfaces after a pull-test. It can be seen that the fracture occurs inside the glass. The fracture initiates in the glass, then the crack propagates along the glass, and sometimes it propagates into the other wafer through the interface without damaging the bond. It further demonstrates that a good bond is formed. The thermal residual stress induced by the low temperature anodic bonding process was measured via the change in the curvature of the wafers. As expected, low temperature bonding largely reduced the induced stress. In this study, all the bonded
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Interface
(a)
Interface
(b)
21.14 Typical cross-sections of bonded glass and glass wafers under different bonding temperatures of (a) 200 °C and (b) 300 °C, at voltage of 600 V, bonding force of 200 N, bonding time of 10 minutes and vacuum of 1 Pa; adapted from [92].
wafers under different conditions were measured, and no residual stress was detected.
21.5
Bonding of dissimilar materials
Most materials, if properly polished and their surface is properly conditioned, do adhere to each other at room temperature and can thus be used for wafer
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Bond strength (MPa)
25 20 15 10 5 0 150
200 250 300 Bonding temperature (°C) (a)
350
400 600 800 Bonding voltage (V) (b)
1200
30
Bond strength (MPa)
25 20 15 10 5 0 0
200
1000
21.15 Bond strength (a) under different temperatures at voltage of 600 V and (b) under different voltages at temperature of 250 °C. Bonding force of 200 N, bonding time of 10 minutes and vacuum of 1 Pa; adapted from [92].
bonding approaches as has been shown by the group of Haisma in Philips [93, 94]. The materials are certainly not limited to semiconductors and can be present in single crystalline, polycrystalline or amorphous form. Examples of different material combinations fabricated by wafer bonding include Si on sapphire for improved SOS material [95], crystalline quartz on silicon for high frequency applications [96], Si on fused quartz or glass for HDTV projection masks [97], and GaAs or InP on Si for combining opto- and microelectronics [98–100]. Interesting applications are combinations of different III-V compounds for the fabrication of vertical cavity surface emitting lasers (VCSEL) [101] and light emitting diodes (LED) [102] pioneered at Hewlett-Packard. For the bonding of III-V compounds the wafer bonding
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(a)
(b)
21.16 Typical fractures at bonding temperature of 300 °C, voltage of 600 V, bonding force of 200 N, bonding time of 10 minutes and vacuum of 1 Pa. (a) and (b) show mating surfaces; adapted from [92].
process frequently is performed not at room temperature but at elevated temperatures (around 300–500 °C) in flowing hydrogen or at even higher temperatures under the influence of a weight [103, 104]. Wafer bonding may be performed via all kinds of intermediate layers such as oxides, nitrides, metals and silicides. Especially important is the bonding via polished polysilicon layers not only for advanced dielectrically isolated (DI) wafers [105] but also for the fabrication of complex three-dimensionally integrated devices or for micromechanics. In this context, the importance of
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chemical mechanical polishing (CMP) should be stressed which nowadays is widely available in microelectronics facilities. If two different materials are combined, a problem due to a mismatch of the thermal expansion coefficients may arise during the heating step following room temperature bonding. This may lead to debonding, interface sliding or breakage of the bonded pair at elevated temperature [99]. Only a few cases of interesting materials combinations are known for which the two materials have sufficiently close thermal expansion coefficients. Examples are silicon and silicon carbide, silicon and glasses specifically developed for this purpose and last but not least GaAs and sapphire which might become important for combining GaAs based electronics with superconducting microwave devices based on thin films of high temperature superconductors epitaxially grown on sapphire. In the case of GaAs/sapphire bonding the bonding has to be performed at elevated temperature around 500 °C in order to avoid undesirable surface bubbles [100]. Wafer bonding may also be used for many different areas of materials integration such as three-dimensional photonic crystals [106], bonding of diamond covered silicon structures for implantable biomedical devices, ferroelectric films directly on silicon without the presence of intermediate phases (by a low temperature wafer bonding approach), or magnetoelectronics (by allowing combination of ferromagnetic metals with silicon without the formation of silicides by appropriate bonding procedures), to name just a few. Such possible applications of materials integration are astonishingly diverse and apparently limited more by our lack of imagination than by technological restraints.
21.6
Concluding remarks
Wafer bonding has increasingly become a key technology for materials integration, three dimensional (3D) systems integration and packaging, as well as vacuum packaging and hermetic sealing for various micro/nanosystems, electronics and optoelectronics, which use a wide range of materials including Si, III–V semiconductor compounds, glass, and ceramics. The broad applications can be categorized in the following areas: • • • • • • • • •
nanoelectromechanical systems (NEMS) microelectronics optoelectronics nanosystems micro/nano-optoelectromechanical systems (M/NOEMS) micro-fluidics and bio-MEMS substrate fabrication semiconductor-on-insulator hermetic sealing
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• vacuum packaging • encapsulation. To bond wafers or substrates together, numerous techniques have been developed. These techniques can be categorized into direct bonding, intermediate layer bonding, and anodic bonding. Each of the processes has its merits and drawbacks. Although several technologies have been established for certain applications, there are still many challenges where more development is required. In order fully to utilize the potential of wafers, development of wafer bonding techniques becomes critical. One of the major wafer bonding research areas worldwide focuses on achieving strong wafer bonding at the lowest possible temperature. It is believed that with breakthroughs in bonding technology, more applications can be realized to meet both high quality and low cost requirements.
21.7
References
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22 Plastics microwelding I J O N E S , TWI Ltd, Cambridge, UK
22.1
Introduction
Microwelding of plastics is of practical importance as medical and electronic devices in particular are manufactured to smaller scales and with increased complexity. Following advancements in microinjection moulding, lithography and micromachining it is possible to manufacture polymer components with features a few microns in size, and joints of the order of 100 µm wide are often required. The devices are often intended to be single-use and disposable. Hermetic seals are often required, which can be achieved using either adhesive bonding or welding methods. Adhesives have the advantage that they can be specified to give little or no heating of the parts, but application and curing times can be significant. To be applied effectively, a welding method needs to be fast and have a small heat-affected zone to reduce distortion and avoid affecting thermally sensitive coatings or components in medical or electronic devices. In this chapter the principles of welding plastics are introduced and the methods available to the manufacturer of plastic microcomponents are described in detail. Of the 18 or so methods available for welding plastics only two have been used to any significant extent in microjoining – ultrasonic and laser welding. These two are discussed in detail. Some of the other methods have the potential to be scaled down to enable joints to be made at microscales. For these methods an outline description is provided.
22.2
Theory of welding plastics
22.2.1 Plastic materials and thermal effects Thermoplastics may be divided into two groups, amorphous and semicrystalline, each with its own set of properties. Amorphous thermoplastics have no order beyond that on the level of the repeat units, which make up the long chain macromolecule. The linear chains exist often in random positions 665 WPNL2204
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and orientations in the bulk polymer and are separated by a distance proportional to their dynamic state, which is governed by temperature. If a linear thermoplastic has very regular chemistry within repeat units (stereoregular) and little or no chain branching then it may exist in crystalline form [1]. Crystalline synthetic thermoplastics are invariably partly crystalline and partly amorphous and so the term crystalline polymer always implies partial crystallinity or semicrystalline [2]. For every amorphous thermoplastic homopolymer (one exhibiting no crystallinity) there exists a narrow temperature region in which it changes from a viscous or rubbery condition at temperatures above this region, to a hard and relatively brittle one (sometimes called glassy) below it [3]. This temperature region is called the glass transition temperature Tg (Fig. 22.1) and is usually obtained from a volume versus temperature plot of observations taken on cooling [2]. Since movement of whole polymer chain segments is a necessary prerequisite of welding, amorphous thermoplastics need to be above Tg before welding can take place. Semicrystalline thermoplastics are composed of crystalline regions and amorphous regions. For flow to occur in these polymers their temperature must be above the crystalline melting point Tm, as shown in Fig. 22.2, which is the temperature at which all of the crystalline regions have disappeared [1]. Semicrystalline thermoplastics will generally have a Tg associated with the amorphous regions and Tm associated with the crystalline regions, with Tm > Tg. Viscous flow will generally occur only above Tm.
Liquid (or rubber)
Specific volume
Glass
Tg Temperature (°C)
22.1 A typical graph of specific volume versus temperature for an amorphous thermoplastic. From [4].
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Tm
Temperature (°C)
22.2 A typical graph of specific volume versus temperature for a semicrystalline thermoplastic. From [4].
Thermal welding of thermoplastics relies on the temperature at the interface being above Tg for amorphous thermoplastics or above Tm for semicrystalline thermoplastics. Above Tm or Tg there is increased chain mobility of the thermoplastic in the region of the weld. It is this increased mobility of the chains that allows them to diffuse across the joint interface and entangle with chains on the other side of the interface. This is the mechanism for generation of strength at the joint and the formation of a weld between surfaces of similar thermoplastics.
22.2.2 Strength development at a joint (diffusion by reptation) A weld in a thermoplastic is generally formed when two pieces above their glass transition temperature or melting point are brought together for a period of time at temperature under sufficient pressure to ensure intimate contact at the joint interface. Based on this simple argument, early theories on welding were based around diffusion [5]. Molecular chains are said to diffuse across the interface and strength develops at the joint due to bridging by these chains (Fig. 22.3). Mathematics have been developed to account for the strength development as a function of time based on diffusion theory [6].
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22.3 Scheme for polymer chains crossing an interface by diffusion leading to coalescence of two surfaces. From [5].
The term ‘tack’ has been used to describe the build up of mechanical strength at joints between similar and dissimilar materials. Tack at joints in similar materials is also sometimes called autohesion and was studied extensively by Voyutskii and Vasenin. The main results of Vasenin’s derivation are that autohesion force for separation (F) has the following dependencies: F ∝ r (the rate of separation) F ∝ M–2/3 (molecular weight) F ∝ t1/4 (time of contact before breaking) Or in combination the force to separate two polymer surfaces is F ∝ t1/4M–2/3e–1/T
(1)
where: t = time of contact, M = molecular weight, T = temp. Since the work by Voyutskii and Vasenin, the next most significant advance chronologically in the theory of the welding of thermoplastics came from the development of reptation theory, which is important to an explanation of the development of weld strength. The word reptation was originally coined by de Gennes [7], from the Latin reptare (to creep), to describe the motion of polymer chains under certain circumstances. He first considered the snake-like movements of a linear polymer chain inside a strongly cross-linked polymeric gel. The gel provided a regular array of fixed obstacles through which the chain could not pass. Instead the linear chain had to wriggle between the obstacles. The reason for this approach was that theories for the motion of polymer molecules in their molten state could not be made to agree with the experimental observations of viscosity and self-diffusion. De Gennes therefore simplified the problem to one linear chain moving in a cross-linked gel. Amongst others, Wool developed the reptation theory into a model for welding. Wool’s theory for the welding of polymers has evolved since the late 1970s [8] and one incarnation involves dividing the welding operation into five stages [9]: 1. surface rearrangement 2. surface approach
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3. wetting 4. diffusion 5. randomisation. Using the five identified stages of welding, Wool next defined a welding function W, such that the mechanical energy G required to separate two welded polymeric parts, was given by: G = W (t, T, P, M)
(2)
where W is a function of t = contact time, T = temperature, P = contact pressure and M = molecular weight. The general conclusions of the theory include: • • •
long molecules take longer to diffuse – increased weld time at least half molecules to cross weld for optimum properties high temperatures lead to rapid welds.
22.2.3 Effect of main welding parameters From reptation theory and experimental observations, the following inferences can be drawn in relation to the effects of the main welding parameters.
Temperature Pressure Weld time Molecular weight
High
Low
Fast weld Thermal degradation Chain diffusion inhibited Oxidative degradation, flash Very long weld times and high pressures required Slow diffusion
Slow or no weld Insufficient surface wetting Insufficient diffusion, weak weld Rapid weld times, low pressures required, weak parent materials and welds
22.2.4 Modelling of polymer welding Molecular simulation of polymer welding has been studied using a Monte Carlo model at the mesoscale, i.e., with detail of polymer chain motion, but not with atomic scale detail [10]. The models described the polymer chain interdiffusion and could predict completion times for welds when surfaces were brought into contact at a temperature greater than Tg. The models were successful for linear amorphous polymers and there is a suggestion that the theory can be extended to semicrystalline and/or branched polymers. More complex is the extension to all the commercial welding processes which can involve flow of the material and non-isothermal heating conditions in combination with chain interdiffusion. Figure 22.4 shows the output from a mesocale model of polymer welding after different periods of heating following contact of the two pieces.
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t = 1000 ~ 10–8 s
t = 1500
22.4 Simulation of polymer welding showing chain interdiffusion at a joint. The two parts are shown in different colours for clarity [10].
Thermal modelling of polymer welding (i.e. using heat flow considerations rather than chain interdiffusion) is applied most easily when the process is predominantly one of diffusion at the surfaces and has minimal melt flow, which complicates the heat flow calculations. Transmission laser welding is one such process. A finite element thermal model for transmission laser welding has been developed [11] and used to simulate the welding of thermoplastic polymers. Cooling rate curves generated experimentally and using finite element analysis have shown good agreement, as have the weld dimensions measured and calculated.
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The model takes details of the materials’ radiation absorption coefficient, thickness, reflectivity, melting point, density, heat capacity and thermal conductivity. It also takes details of the laser source including the laser beam profile, laser power and welding speed. Through analysis of the heat flow it is possible to generate the thermal history for all points in the joint during the welding cycle. Thus it can provide the weld dimensions, the peak temperature and, with details of the materials’ mechanical properties, the residual stress remaining after welding can be estimated. Figure 22.5 shows a typical output from a finite element thermal model.
°C 208 164 143 123 102 82 61 41 20 0
22.5 Typical temperature distribution for a laser welded specimen resulting from the finite element analysis.
22.3
Introduction to welding processes
Before considering the plastics welding methods which may be applied for microjoining, it is useful to review all the methods available. Only a few of these can be applied at small scale as is summarised in Table 22.1. Techniques for thermally welding polymers can be broadly divided into three groups. Those where heat is generated by mechanical movement of components to be joined, those where heat is generated by an external source and this source heats the joint by thermal conduction, and those using electromagnetism directly.
22.3.1 Techniques where heat is generated by mechanical movement • spin welding • vibration welding • ultrasonic welding
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• orbital welding • friction stir welding. When two surfaces are brought into contact and rubbed together under pressure, there is generally a temperature increase at the interface. The thermal welding of thermoplastics may most efficiently be accomplished by the generation of heat at the joint interface and for this reason welding techniques employing mechanical movement have found many commercial applications. The rate of temperature increase at the interface is generally related to the speed of the mechanical movement such that the greater the speed the more rapid the temperature increase. High heating rates generally imply high welding rates and for this reason techniques where heat is generated by mechanical movement are often exploited in mass production. If the amplitude of vibration is small and correspondingly the frequency is high then the methods become applicable to parts with small dimensions. Ultrasonic welding provides such conditions and is the only method in this group to be considered for microjoining applications.
22.3.2 Techniques employing an external heat source • • • • • • •
hot plate welding hot bar welding impulse welding hot gas welding extrusion welding forced mixed extrusion welding flash free welding.
Thermoplastics generally have low coefficients of thermal conductivity particularly when compared to metals. This means that if heat can be applied to the joint area prior to joining, the temperature gradient normal to the surface may be high. However, in joining small parts the loss of heat from the surface may be too rapid for the joint to be completed before cooling. Therefore methods, such as hot plate welding, where heating is carried out with the parts separated are normally only applied to large components. The only methods applicable to microjoining are hot bar or impulse welding where the heat is applied externally to the parts as they are pressed together, and hot gas welding where a narrow gas stream may be used to provide a precise heating position to heat two small parts simultaneously.
22.3.3 Techniques which directly employ electromagnetism • resistive implant welding • induction welding
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dielectric welding microwave welding infra-red welding laser welding.
The welding techniques that directly employ electric or magnetic fields are listed above. Although most thermoplastics exhibit dielectric (non-electricallyconducting) properties, they generally produce some response when excited by an electric or magnetic field at a certain frequency. Heat is generated at the joint line by an electric or magnetic field or both. In some cases the field heats the thermoplastic directly, e.g., dielectric (high frequency) welding and in other cases the field heats a different material at the joint which in turn heats the thermoplastic, e.g., resistive implant welding. Of these methods only laser welding has found application for joints at small scale. This is as a result of the precise control available on the position and amount of the energy applied by the process at the joint.
22.3.4 Comparison table Table 22.1 summarises the plastics welding processes and their applicability at microscale.
22.4
Processes for microwelding of plastics
Many of the methods used for macrowelding of plastics could be miniaturised, as shown in Table 22.1. However, the two main methods that have been developed and made commercially available for microwelding are ultrasonic and laser welding. The process, equipment, welding parameter effects and examples for these two methods are described in Sections 22.5 and 22.6.
22.5
Ultrasonic welding
22.5.1 Process description Ultrasonic welding is a process which uses mechanical vibrations to soften or melt the thermoplastic material at the joint line. The parts to be joined are held together under pressure and subjected to ultrasonic vibrations, usually at a frequency of 30 kHz or more for small parts. Higher frequencies than the more usual 20 kHz are used for small components because they are less likely to damage any small internal parts. The mechanisms responsible for generating heat at the joint line are not well understood and the heating effect of the ultrasound varies with the degree of crystallinity of the material being welded. Generally, semicrystalline plastics require more energy
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Table 22.1 Summary of plastics welding processes Method
Description
Microscale joining equipment available
Spin welding Vibration welding Ultrasonic welding Orbital welding Friction stir welding Hot plate welding
Rotational friction welding Reciprocal motion friction welding
Potentially No
Heating through the ultrasonic vibration of one part against another Frictional heating using an orbital motion Mixing material at a joint using a rotating or reciprocating tool Heated plate between the two parts, that is removed to allow the parts to be brought together Heated tool applied to the outside of the parts. One part must be thin Resistively heated wire applied to the outside of the parts with a short heating pulse applied Stream of hot gas used to heat parts directly with or without addition of a filler rod at the joint Melt extruded into the joint between the two parts For thick section materials, a combination of extruded material, heated gas stream and rotation of the material at the joint Heated tool surrounds the joint and supports it during the weld cycle Conductive mesh implant at the joint is resistively heated to weld the parts Conductive mesh implant at the joint is inductively heated to weld the parts Radio frequency field is applied across the parts that heats some plastic types directly Conductive implant at the joint has a heating current induced in it by a microwave field Similar to hot plate but with a noncontact heating element or lamp replacing the contact plate Laser heating either directly or using an absorber at the joint surface
Yes
Hot bar welding Impulse welding Hot gas welding Extrusion welding Forced mixed extrusion welding Flash free welding Resistive implant welding Induction welding Dielectric welding Microwave welding Infra-red welding Laser welding
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Potentially Potentially
Potentially
No No
Potentially
Potentially
Potentially Potentially Potentially
No
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to weld ultrasonically than amorphous plastics. The ability to weld a component successfully is governed by the design of the equipment, the mechanical properties of the material to be welded and the design of the components. Components can be welded in either a projection or a shear joint configuration. In the projection joint, the ultrasonic energy is focused at a triangular point on the moulding known as the projection. The projection melts and forms the molten bead of material which makes the weld. In the shear joint configuration, the two components are, effectively, rubbed together at the joint line, albeit with a small amount of movement. The friction at the interface causes the component to locally heat and melt to form the weld. A further practical consideration is the thickness of material through which the ultrasonic energy needs to travel between the welding horn and the joint. In plastics that transmit ultrasound with low losses (amorphous plastics in general), the welding horn can be remote from the joint (more than 6 mm). This is termed ‘far-field’ welding. When ultrasonically welding other thermoplastics with higher losses, the plastic between the horn and the joint will heat. The horn should be placed as close as possible to the joint (less than 6 mm). This is termed ‘near-field’ welding. Ultrasonic welding times are short (typically less than one second) which makes the process ideal for mass production.
22.5.2 Equipment A typical ultrasonic welding machine (Fig. 22.6) consists of four main components: 1. microprocessor control system, ultrasonic generator and user interface 2. machine stand, consisting of the base-plate, pneumatic system and welding press 3. welding stack incorporating the ultrasonic transducer, booster and welding horn 4. component holding fixture.
22.5.3 Welding parameters Once the welding parameters have been optimised, ultrasonic welding is a consistent and reliable process. Amplitude, power, force, hold time and welding mode must all be correctly set to achieve the required joint quality. Amplitude selection depends on the type of material being welded, amorphous or semicrystalline.
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Transducer/ converter
Booster Welding horn Moulded parts Pneumatic system
Holding fixture Emergency stop button
Welding press Base-plate
Two-hand safety operation
22.6 Schematic of a typical ultrasonic welding machine.
Amplitude In general, semicrystalline materials require more energy and, therefore, more horn tip amplitude than amorphous materials. The gain in the horn and in the booster are altered to set the amplitude. Depending on the material being welded the amplitude set may be 20–60 µm for small parts. Power Typically, a small component produced in an easy-to-weld material, such as polystyrene, will require a machine with a maximum power output of 800 W. During the welding cycle, the machine may never draw the maximum power available. The majority of equipment has the ability to limit the amount of power consumed during the welding cycle. Welding mode The ultrasonic welding cycle consists of two phases, the ‘ultrasonics on’ period and the hold time. Throughout the welding cycle, the welding horn remains in contact with the components being welded and an axial force is applied. The period when the ultrasonic energy is applied can be controlled
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by a time setting (typically 0.1–0.3 s for small parts), energy applied or the movement (displacement) measured at the joint. Time, energy or distance are the three welding modes. Force The force applied to the joint must be enough to hold the parts together but not over-dampen the vibrations. For parts with less than 20 mm weld length typically forces less than 200 N are required. Hold time The parts are supported under pressure for typically 0.3–0.5 s to allow the joint to cool.
22.5.4 Advantages and disadvantages The main advantages of using ultrasonic welding over other welding processes are: • • • • • •
Short weld times – suits high volume production. Simple technique – straightforward to put into production. Consistent process (once weld parameters are optimised) – reliable operation. Clean process – can be used in a clean-room environment. No thermal damage to component – heating is confined to the immediate area around the joint. Relatively low cost.
The main disadvantages of ultrasonic welding are: • • • •
All aspects of the process must be considered together, and carefully, to ensure successful welding. Particles of flash are sometimes generated. Vibrations may affect sensitive components. The shape at the joint must be carefully designed and controlled to ensure the energy is directed correctly.
22.5.5 Examples Medical devices and electronic packages are the most common items welded using ultrasonic processing. The ultrasonic welding process is used where short weld times and a hermetic seal is required. The process can be easily automated. Medical sample flasks requiring hermetic seals and anaesthesia filters are ultrasonically welded.
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Ultrasonic welding is also used to join hydrophilic contact lenses [12]. The welding method has been designed such that directed and modified energy is applied to the (convex) output horn, which extends into the outer concave portion of one of the mould pieces. A shoulder has also been designed to rest upon the flange of the piece, providing centring. For applications in biotechnology in particular and potentially for the wider area of microelectromechanical systems (MEMS), ultrasonic welding has been shown to be effective for sealing covers on microfluidic devices [13]. In one case a micropump with a 14mm diameter and 500µm channel widths was welded from PMMA or PEEK. Ultrasonic sewing can be used to join textiles together, replacing the use of thread and staples. Typical applications include hook and loop straps [14]. A variation on the process uses a two-part horn with a gap between the tips of 0.15–0.55 mm, allowing independent vibration of the two [15]. The result is a weld of a similar size to the gap between the horn parts, which is effective for microwelding of thin materials over a limited area of a plastic part.
22.6
Laser welding
22.6.1 Process description Laser welding was first demonstrated for thermoplastics in the 1970s, and since the late 1990s has been used in mass production. The technique, suitable for joining sheet, film and moulded thermoplastics and textiles, uses a laser beam to melt the plastic in the joint region. Lasers are well suited to microwelding due to the small beam size available (down to less than 10 µm) and the potential for precise positioning of the heat source. Two general forms of laser welding exist; direct laser welding and transmission laser welding. In direct laser welding, the materials are heated from the outer surface possibly to a depth of a few millimetres. Normally, no specific radiation absorber is added to the plastics. Laser sources from 2.0–10.6 µm wavelength are used. At 10.6 µm (CO2 laser) radiation is strongly absorbed by plastic surfaces, allowing high speed joints to be made in thin films. At 2.0 µm, where the absorption is less strong, a fibre laser might be used to make welds in sheet a few millimetres thick. Direct laser welding is not widely applied for joining plastics. Transmission laser welding is widely used. Laser sources from 0.8–1.1 µm wavelength such as diode, Nd:YAG and fibre lasers are applied. The radiation at this wavelength is less readily absorbed by natural plastics without absorbing additives. Laser absorbing additives are put into the lower part or applied as a thin surface coating at the joint. The parts can be positioned
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together before welding and the laser beam passes through the upper part to heat the joint at the absorbing surface of the lower part (Fig. 22.7). The absorber in or on the lower plastic is typically carbon or an infra-red absorber with minimal visible colour (Clearweld®) [16]. These and other variations allow a wide range of part colours and appearances to be welded. Transmission laser welding is capable of welding thicker parts than direct welding, and since the heat affected zone is confined to the joint region no marking of the outer surfaces occurs.
22.6.2 Equipment and variations The main elements of a laser welding system are: • • • •
power supply including a chiller for higher power laser sources laser source beam delivery optics (lens, mirror or fibre based) beam focusing or shaping optics including masks if required
Laser beam Transparent to infra-red laser
Clamping pressure
Infra-red absorber Weld zone
Clamping pressure
Transparent or opaque to infra-red laser
22.7 Diagram of transmission laser welding showing the movement of a beam over a workpiece. The lower part can be arranged to be an infra-red absorber or the absorber may be placed at the joint surface.
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• • •
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beam manipulation optics such as galvanometer controlled mirrors (if used) workpiece clamping and support workpiece manipulation.
Laser types for transmission laser welding The main types of NIR laser used for through-transmission laser welding, and selected properties, are listed in Table 22.2. Additional details of each type of laser are given below. Nd:YAG lasers are widely use in industry for materials processing. Highpower systems are bulky, but lower-power systems are relatively compact. Water-cooling is usually required. The beam is transferred from the laser to the workpiece via an optical fibre. It is feasible to combine the beam from more than one laser to produce higher powers if required. The high beam quality allows a relatively small spot size to be produced. Focused spot sizes of less than 50 µm are available for sources with power greater than 10 W if required. Diode lasers produce radiation at a wavelength of 808 nm (InGaAlAs) or 940 nm (InGaAs). Water-cooling is usually required. Their relatively low beam quality means that they cannot be used to produce spot sizes as small as Nd:YAG or fibre lasers. However, this is rarely a problem for plastics laser welding, where the relatively low purchase and running costs have attracted a great deal of interest. The beam may be delivered by an optical fibre, but the diode is sufficiently small and light that it is often feasible to use a direct system, in which the diode is included with a lens system in a single unit, typically ~150 × 150 × 300 mm. This unit can readily be mounted on a gantry system or robot arm to manipulate the beam. Rare-earth doped fibre lasers typically supply a single wavelength in the range 1000 nm to 2100 nm. In the field of materials processing, much interest has focused on wavelengths around 1100 nm to provide a direct replacement for Nd:YAG lasers, with equivalent beam quality, but greater efficiency. Systems are relatively compact and can be air-cooled. In the field of plastics welding the use of fibre lasers has been demonstrated for a range of applications, including precision welding of films, textiles and larger moulded parts. Table 22.2 NIR laser types Laser
Nd:YAG
Diode
Fibre
Wavelength (nm) Efficiency (%) Approximate cost for 100 W system (£k) Beam quality
1064 3 40
780–980 30 10
1000–2100 20 30
High
Low
High
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Beam or workpiece manipulation The beam or workpiece manipulation equipment for microwelding will take one of the following forms as illustrated in Fig. 22.8 and described in more detail below. With the laser fixed (a), the part can be manipulated to form a continuous weld. This can be achieved, for example, with rollers, or a single or two-axis moving table. This type of system is relatively simple to set up and program. It would not normally be used if three-dimensional welds are required. The optical system for a fibre delivered laser or the laser head for a direct diode laser can be mounted on a variety of robotic systems (b). These range from simple two-axis gantry systems to multiple-axis robotic arms. The laser is then manipulated around the part to be welded, potentially allowing complex, three-dimensional welds to be produced. To facilitate automatic production, it is feasible to combine a moving laser with a moving part, for example by using a rotating table to present different faces of a component to a laser mounted on a robot arm. The resolution is limited by the size of the laser beam which could be as small as 10 µm. Alternatively, the laser energy can be spread into a line (c) and then passed over the component, either by moving the laser or by moving the
(a)
(b)
(d)
(c)
(e)
22.8 Welding methods (from [17]): (a) moving workpiece, (b) moving laser, (c) curtain laser, (d) simultaneous welding, (e) scanning laser.
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part. A mask is typically used to ensure that only the relevant areas of the component are exposed to the radiation. This is particularly suited to small components with a complex weld geometry. The process would usually be used only to produce two-dimensional welds. The welds are completed very quickly with a single sweep of a line source but the resolution is limited to around 100 µm. If a large number of identical welds is required then an array of diode lasers can be assembled in the shape of the weld (d). This is then used to irradiate the whole joint simultaneously, with a typical cycle time of 1–3 s. This approach is well suited to automated assembly. The equipment used is frequently based on ultrasonic welding equipment, and this process is typically used in place of ultrasonic welding where a good cosmetic appearance is required or for components that are sensitive to vibration. Two- and threedimensional welds can be produced. The entire joint is welded at the same time, allowing more collapse of the polymer at the joint, and therefore allowing wider part tolerances. The laser radiation can also be very rapidly manipulated or scanned by a pair of orthogonal rotating mirrors (e), over an area that may range from 50 mm × 50 mm up to approximately 1000 × 1000 mm. In general, a larger working area implies a longer working distance and a larger spot size. It is possible to co-ordinate an assembly of a number of scanning systems to give a larger working area. Two- and now three-dimensional welds can be produced using scanning optics. Repeatedly scanning the laser at high speed over the same path can be used to give quasi-simultaneous welding. As for simultaneous welding, this welds the entire joint area at the same time, allowing more collapse of the material in the joint and potentially allowing wider tolerances. Clamping systems A wide variety of clamping systems have been used for through-transmission laser welding. They are mostly variants of the two systems illustrated in Fig. 22.9. Variants of the fixed clamp include systems using mechanical fastenings, rather than an actuator, to apply a load. In the simplest variant, if the part design allows it, a bolt can be passed through the workpiece to apply the load. The transparent cover must be rigid enough to provide the clamping pressure. Thick acrylic or plain plate glass can be used. Borosilicate glass is less vulnerable to thermal shocks during welding, but more expensive. For welding of high-temperature polymers, quartz glass may be used. In all cases it is important to ensure that suitable safety precautions are taken to avoid the risk of injury if the transparent cover breaks while it is under load. The moving clamp can use bearings, rollers, or a simple sliding shoe to apply a clamping load. Because the load is applied only at the point where
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Laser radiation
Transparent cover
Pneumatic actuator
(a)
Movement of laser and clamp
Laser radiation Roller or bearing clamp
Support table (b)
22.9 Clamping systems for through-transmission laser welding (from [17]): (a) fixed clamp, (b) moving clamp.
the joint is irradiated, clamping loads may be much lower when a moving clamp is used. There is therefore less risk of distorting the workpiece, and equipment can be less bulky. This is particularly advantageous for large components, where application of a suitable clamping pressure to a large area can require large loads.
22.6.3 Welding parameters As with all plastics welding processes there are three critical process parameters, the temperature, time and pressure. The energy density used during welding combines the process parameters
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of temperature and time. It is determined by the laser power, the spot size at the joint, and the irradiation time (for fixed processes) or welding speed (for processes in which the part moves with respect to the laser). Energy density = Power × Time Spot size or
Energy density =
Power Spot width × Speed
If the energy density is too low then insufficient heating takes place and the material at the joint is not held at a high enough temperature for a sufficiently long time to form a strong weld. If the energy density is too high, then excess heating can degrade the polymer at the joint, resulting in porosity, or, in extreme cases, burning or charring of the polymer. Either case results in a weld of lower strength than the optimum. In practice, a relatively wide processing window can usually be found within which satisfactory welds can be produced. Typically laser welding applications use an energy density within the range 0.1–2 J/mm2, although this will vary depending on the depth of melt required to ensure a satisfactory joint. Although the energy density can be used to characterise the welding process, it should be treated with caution. The conduction of heat away from the joint during welding means that using the same energy density with different laser powers will not necessarily result in the same quality of weld. For example, with a constant spot width, doubling the power will usually allow the speed to be more than doubled, whilst retaining the same performance from the weld. The pressure applied is controlled using the clamping system. If the workpieces are not clamped together during welding, or if the pressure at the joint is insufficient, then the joint faces will not be in intimate contact. This results in poor conduction of heat to the upper workpiece and limited interdiffusion of polymer chains on either side of the joint. Both effects result in a weld of lower strength than the optimum. Care is needed to ensure that a clamping load actually provides pressure at the joint. Typically clamping pressure in the range 0.1–1 N/mm2 is used. If the workpieces bend under the clamping load in such a way that the joint is distorted then a poor weld can result. For this reason, it is often useful to have some compliance, for example an elastomeric element, in the clamping system.
22.6.4 Laser microwelding When laser welding small or complex parts, precise control of the process is required. Precision welding is dependent on the following factors: 1. The amount of energy applied, controlled by laser power, welding speed or time and the amount and absorption properties of the absorber.
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2. The location of the energy application, controlled by the dimensions of the heat source (either laser beam or absorber size), and the accuracy of the manipulation equipment. Manipulation equipment is available to position components (and the laser source) to an accuracy better than 1µm [18]. Accurate control of the laser power is also well understood and relatively easy to achieve. The limiting factor in providing precise location of the weld is therefore the localisation of the heat source either by the laser beam dimension or by the placement of laser absorber material at the joint. Additionally, the control of the energy applied depends on the amount and position of the absorber material and how and how it behaves in the laser beam as it delivers heat to the joint. The precision of weld location achievable in laser welding plastics may be estimated by measuring the minimum weld width possible with different methods and equipment systems. To provide an understanding of the limits of precision that are related to laser beam size, Nd:YAG and fibre laser sources with small focused spot sizes (less than 100µm wide) have been used to make narrow welds at high speed as an alternative to high power diode lasers, which typically have beam dimensions greater than 500µm [19]. It should be noted that lower power diode sources (around 1W power output), also providing spot sizes less than 100µm wide, can be used for transmission laser welding of plastics. In addition to weld localisation using the laser source, attempts have been made to precisely locate the absorber material either using a liquid coating process or a solid form as fine filaments to define the weld position. In practice it is found that it is much more effective to use the laser source and have a wide area of absorber preferably mixed into the plastic of the lower substrate. In tests of the minimum weld widths achievable from diode, Nd:YAG and fibre laser sources are dependent on the minimum laser beam dimension (or mask dimensions if they are used). A focused beam width of approximately 60 µm gave weld widths as small as 60 µm in trials. An example is shown in Fig. 22.10. Smaller beam widths are available that may well allow welds smaller than this. This is true when the absorber is incorporated into one of the parts being welded. When, alternatively, a film absorber is placed between the two parts then the minimum weld width is increased by at least as much as the thickness of the film. So welds using a 70 µm thick film and a 60 µm beam width were as small as 165 µm in width. Limiting the weld width using the absorber material either by applying liquid as a narrow coating track or by using narrow filaments of absorber material was less successful in tests. The spread of heat from the local region of absorber tends to lead to wider welds.
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Weld
500 µm
250 µm
22.10 Section and top view of weld in PC with Clearweld absorber using an Nd:YAG laser with a beam width of 54 µm (11W, 9.0 m/min) providing a weld width of approximately 70 µm. The high consistency of the weld size is shown by the top view on the right.
22.6.5 Advantages and disadvantages The main advantages of transmission laser welding over other processes are: • • • • • • • • • • • • • •
joint designs are simple flat to flat surfaces in general hermetic seal possible fast, <1 sec weld possible non-contact no vibration no particulate generation precise placement of welds no surface damage low residual stresses complex shapes possible localised heating – no thermal damage to sensitive features close to weld multiple layers can be welded simultaneously thin or flexible substrates can be welded little or no flash.
The main disadvantages of transmission laser welding are: • • • • •
it is expensive compared to other plastics welding equipment joint surfaces must be of good quality part clamping must be designed carefully to ensure contact at the whole of the joint area during welding laser absorbing material must be added to one of the plastics or at the joint surface the top part must transmit the laser radiation. This can limit the thickness of the top substrate when welding plastics with low transmission. For example, PEEK or some filled plastics. WPNL2204
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22.6.6 Examples Direct laser welding can be used to make small welds in thin materials at high speeds [20]. Figure 22.11 shows an example using a CO2 laser for welding polyethylene film. Transmission laser welding has been considered for a wide range of application areas. Some of these are as follows: • • • • • • •
electronic packages textiles biomedical devices windows and signs food and medical packaging visual displays automotive components.
A variety of methods using transmission laser welding have been used to create small and complex welded structures in plastics. A spin coating of infra-red absorber dye has been used to weld 250 µm thick polycarbonate (PC) foils using a Nd:YAG laser [21]. Minimum weld widths of around 150 µm were achieved and it was concluded that weld width consistency could be improved by improving the spin coating consistency. It has also been shown that a film could be welded over square channels (50 × 50 µm) in a polymethylmethacrylate (PMMA) base plate using a 15 mm wide diode laser line source covering many channels at once [21]. The process was applied to sealing of capillary gel electrophoresis chips made of PMMA containing 300 channels with 50 µm width and depth to be used for low-cost and high-throughput routine analysis of protein mixtures. [22]. Low power (<1 W) fibre-coupled laser diodes are available which provide a focused spot width of 25 µm and can be used to produce weld widths of 50 µm in clear to carbon black filled PC [23]. Weld strengths from 6–16 MPa were measured
22.11 A CO2 laser weld in 100 µm polyethylene film at 100 m/min with 100 W laser power. The weld is approximately 0.5 mm wide.
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and thermal degradation of the PC was identified as a problem as well as limitations in process speed. Trials have been carried out in the welding of a polycarbonate dental prosthesis part 2.75 × 1.5 mm in dimension. Both mask welding and quasisimultaneous welding provided good results and weld completion in less than 1 s. [24]. A line source diode laser and mask based system has also been applied to assemble microfluidic devices. Accuracy in positioning of the welding seam in the order of 5–10 µm were claimed [25]. Microtitre components have been sealed with a 80 µm width diode laser source [26]. The seals provided higher quality than similar samples which had been adhesively bonded. Low power diode laser sources have also been reported giving welding seams as narrow as 10 µm in welds in polyethylene terephthalate glycol (PETG). The PETG was coated with narrow band IR absorber [27]. A process using a thin carbon coating (5–20 nm) as an absorber has been developed for welding channels in the 100 µm size range. It was also suggested that smaller features (10–30 µm) could be welded in the same way and potentially multi-layer structures [28]. The feasibility of welding parts with a complex joint pattern, such as round a microfluidic channel system has been demonstrated [19] using a scanning laser source (Fig. 22.12). The parts were in polystyrene and the welds were made by placing a thin film of Clearweld absorber between the base and the lid. The weld width is less than 1 mm.
Weld
♦
10 mm
♦
Unwelded region – polystyrene parts separated by interleaving polystyrene film
22.12 Polystyrene sample (SMB Excelerator® platform) showing complex laser welded joint profile made using the Clearweld process.
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Future trends
In microwelding of plastics there is potential for alternative methods to arise and for developments on ultrasonic and laser welding to lead to yet smaller scale joints. In ultrasonic welding it is expected that smaller scales will arise from the use of higher frequency machines operating at lower amplitude of vibration. The limitation on smaller scales may well be the control of the plastics surfaces where the shape of the energy director is important to the success of welding. The shape will need to be controlled to greater precision generally using micromoulding. Laser welding will benefit from the development of moderately powered lasers with high beam quality and therefore small focused spot sizes, which will lead to a potential for smaller welds. Nd:YAG and fibre lasers are available with spot sizes less than 10 µm. Similar weld precision would be expected. If the other plastics welding methods are considered, then there are the opportunities for conversion from macrowelding to smaller scales. Spin welding, normally applied to tubes or hollow spheres with rotational symmetry could be applied to solid rods of small diameter, though probably not less than a millimetre. Friction stir welding has been applied for welding thin metal films. A process that could be transferred to thin plastic films. Hot bar and impulse welding is already used for films less than 50 µm thick and could be applied to complex forms with small complex shaped heating elements. There is also potential for using small heated tools on the outer surface of thin films to attach them to thicker substrates. This could potentially weld features less than a millimetre in size. The methods that use an electrically conductive mesh implant at the joint will rely on development of small implants to provide opportunity for smaller scale joints. Resistance, induction and possibly microwave welding would provide the means of making the joints where the parts may not be suitable for laser or ultrasonic welding.
22.8
References
1. Stevens M P, Polymer Chemistry, Cornell University Press 1953. 2. McCrum N G et al., Principles of Polymer Engineering, Oxford University Press 1988. 3. Wool R P, Polymer interfaces: Structure and strength, Carl Hanser Verlag, 1995. 4. Wise R J, Thermal Welding of Polymers, Abington Publishing, Cambridge, 1999. 5. Voyutskii S S, Autohesion and adhesion of high polymers, John Wiley and Sons, 1963. 6. Vasenin R M, ‘The sticking pressure in the diffusion theory of adhesion to polymers’, Vys. Soed. 3, 5, 1961 pp 679–685. Rapra Translation 1010 by R Moseley 1962. 7. De Gennes R P, ‘Reptation of a polymer chain in the presence of fixed obstacles’, Journal of Chemical Physics, 55, 2, pp 572. 1971.
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8. Wool R P and O’Connor K M, Craze healing in polymer glasses, ACS Polymer Preprints 21, p 40, 1980. 9. Wool R P and O’Connor K M, ‘A theory of crack healing in polymers’, Journal of Applied Physics 52, 10, pp 5953–5963, 1981. 10. Haire K R and Windle A H, ‘Monte Carlo simulation of polymer welding’, Computational and Theoretical Polymer Science, Volume 11, Number 3, June 2001, pp. 227–240(14), Elsevier Press. 11. Jones I A, Olden E, ‘A thermal model for transmission laser welding of thermoplastic polymers’, TWI Members report 708/2000, July 2000. 12. US Patent 5 759 318. Filed 6 Dec 1996. ‘Apparatus and method for releasably fusing moulded lens pieces’. Johnson and Johnson Vision Products Inc. 13. Truckenmuller et al., ‘Micro ultrasonic welding: Joining of chemically inert microparts for single microfluidic components and systems’, Microsyst Technol (2006) 12: pp 1027–1029. 14. Prescriptions for plastics joining Technical sales brochure from Branson Ultrasonics. 15. Tsujino et al., ‘Ultrasonic plastic welding with a welding tip pair’, Japanese journal of applied physics, Vol 25 (1986) Supplement 25–1, pp 168–170. 16. Jones I A, Wise R J, ‘Welding Method’. Patent WO 00/20157, 1 Oct 1998. 17. Warwick C M and Gordon M, ‘Application Studies using Through-Transmission Laser Welding of Polymers’, Proc. Joining Plastics 2006, 25–26 April 2006, NPL, London. 18. Klein H and Haberstroh E, ‘Laser beam welding of plastic micro parts’. Proc. ANTEC conference of the Society of Plastics Engineers, May 1999, pp 1406–1409. 19. Jones I A, ‘Precision laser welding of plastic components’, TWI members report 876/2007, July 2007. 20. Jones I A and Taylor N S, ‘High speed welding of plastics using lasers’, ANTEC ’94 conference proceedings, 1–5 May 1994, San Francisco, USA. 21. Klotzbuecher T et al, ‘Microclear – A novel method for diode laser welding of transparent microstructured polymer chips’. Proceedings of the 23rd International Congress on Applications of Lasers and Electro-Optics 2004. 22. Griebel A et al., ‘Integrated polymer chip for 2D capillary gel electrophoresis’, Lab Chip 4: 18–23, 2004. 23. Grewell D and Benatar A, ‘Experiments in Micro-Welding of Polycarbonate with Laser Diodes’, Proc. ANTEC conference of the Society of Plastics Engineers, May 2003, pp 1039–1044. 24. Hustedt M et al., ‘ Laser micro welding of polymeric components for dental prostheses’, Proc. of SPIE Vol. 6107, 2006. 25. Chen J-W and Zykbo J M, ‘Diode laser bonding of planar microfluidic devices, MOEMS, bioMEMS, diagnostic chips and microarrays’, Proc SPIE 5718, pp 92–98, 2005. 26. Curtis et al., ‘The uses and advantages of laser sealing in high density assay plates’, Society of Biomolecular Sciences 10th Conf. 11–15 Sep 2004. 27. Ussing T et al., ‘Micro laser welding of polymer microstructures using low power laser diodes’, Proc int conf on multi-material micro manufacture, Karlsruhe 2005, pp 291–293. 28. Pfleging W et al., ‘Laser assisted welding of transparent polymers for micro-chemical engineering and life science’, Proc SPIE 5713 pp 479–488, 2005.
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23 Microjoining in medical components and devices K J E L Y, Boston Scientific, USA and Y Z H O U , University of Waterloo, Canada
23.1
Introduction
The ageing of America and many of the industrialized countries has placed a burden on the healthcare industry. A key area within healthcare today is the burgeoning field of implantable medical devices and components. These devices include the now well-known pacemakers, but also include implantable defibrillators, heart failure cardiac care, neurostimulators, pumps, drug delivery, arterial stents, and other smart electronics or components. All of these devices contain a wide variety of materials and joining processes. Due to the implant size and their intended purpose, the raw materials and assembly processes must be of the highest quality and reliability. Medical devices can be some of the most demanding applications due to the requirements for biocompatibility. The degree of compatibility required is driven by the type of application and the time requirement for interaction with the human body. These applications are typically divided into three classes per FDA specification, with Class I being minimal risk, such as a monitoring device or external pump, to a Class III device such as a permanently implanted arterial stent or pacemaker [1]. An additional requirement for medical devices is the need for extremely high reliability, to avoid any adverse risk to a patient. The combination of these two requirements, biocompatibility and reliability make medical devices a challenge to design and manufacture. Some medical devices seem simple, yet have high degrees of complexity to predict final form and function. An example of this type of device is a bare metal stent. They are simple only in terms of their manufacture (laser cut) from a single tube of specialty metal, typically a cobalt bearing alloy, or other special alloys of nickel-titanium. These alloys can be fabricated with an optimal surface finish for the adhesion of supplemental polymer coatings that are doped with steroid drugs to control the normal inflammation response of the body and to control chronic restinosis. Larger more complex medical devices would include prostheses, including knee and hip replacements, polymer finger joints, and new vertebral disc 691 WPNL2204
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implants. Many of these types of implants use multiple materials and forms of materials to enhance the human body’s response to the material. The hip implant is an excellent example. The ball of the acetabular insert is covered in small cobalt alloy spheres, which yield high surface area for bonding into the pelvic bone. An insert of high molecular weight polyethylene is used in the acetabular cup to provide a lubricious smooth surface for the mating ball socket of the hip implant. In this example the cobalt spheres are joined to the main body of the implant by sintering, while the high molecular weight polyethylene cup is ultrasonically staked through an opening in the metal housing [2]. Implantable cardiac rhythm devices, cochlear implants, and neurostimulator devices are some of the most complex medical devices. Not only do they comprise all of the materials of electronic assemblies, they require the biocompatibility and extreme reliability required for human use. So the selection of the outer case materials is typically driven by the ability to provide a hermetically sealed enclosure, and long-term corrosion resistance. Typical materials now in use include 316L stainless steels, MP35N, titanium, niobium, and alloys of platinum-iridium. Microjoining in this application includes the welding of pacing lead coils to terminal pins via resistance and laser spot welding, laser seam welding thin titanium housings, and spot welding of thin electrode rings to interconnect wires. This chapter will review key aspects of material choices and the typical joining processes for medical devices.
23.2
Materials challenges
Implantable devices are approved for use by the FDA according to the devices’ ability to provide therapeutic value to the patient. Since the full device is approved in the process, individual materials alone are not approved for use. Therefore, most designs leverage off past proven designs where a material has been shown to perform well within a system. Materials that have passed device level approvals for a wide variety of implants are shown in Table 23.1. These materials all have the following key characteristics: • • • •
biocompatibility corrosion resistance consistent surface quality suitability for welding and joining.
These materials are generally chosen at the beginning of the design process due to their corrosion, electrical, and biocompatibility performance. However, that does not mean they can be joined in any combination using any methods. Further, while macroscopic size assemblies can often be joined, small-scale microscopic assemblies require more effort to maintain a tight joint during
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Table 23.1 Implantable materials Material
Application
Joining methods
Platinum Platinum/ iridium Tantalum Titanium
Wires, coils, electrodes, pins Wires, coils, electrodes
Resistance, laser, ultrasonic Resistance, laser, ultrasonic
Wire, coils Wires, coils, terminal pins, PG cases Wires, coils, terminal pins, PG cases, battery cases Standoffs, interconnect pads Wires, terminal, battery cases, capacitor cases Surface mount bumps, general soldering Tubings, seals Tubings, seals, adhesives
Resistance, laser, brazing Resistance, laser, brazing
MP35N Kovar 316 SS Lead free solders Polyurethane Silicone
Resistance, laser Resistance, laser, brazing Resistance, laser, brazing Brazing, soldering, ultrasonics Laser, adhesive Adhesives
joining processes. Tables 23.2(a) and (b) highlight many common medical grade materials and their ability to be joined by various welding technologies. The tables assume a full fusion joint is developed between the parent materials. Many solid-state welds can be formed from combinations that under full fusion would develop unusable intermetallics. Similarly, the speed of cooling in some small-scale laser welds allows the weld pool to freeze before bad microstructure can develop, allowing combinations of materials not allowed in the table. Typical material properties for many common metals and alloys are shown in Table 23.3 [3, 4]. Materials to be joined will be affected by all of the significant variables (process parameters, welding consumables, etc.). For example, the ductility of titanium welds can be greatly reduced through picking up interstitial elements, such as oxygen, nitrogen and hydrogen, if shielding is inadequate during welding. A study of laser welding of titanium indicated that addition of oxygen in pure argon shielding gas (as in the case of shielding deficiency) caused further changes in weld microstructure (Fig. 23.1) and fracture morphology (Fig. 23.2) [5]. The addition of oxygen in the shielding gas increased the amount of acicular and platelet alpha compared to equiaxed alpha in the base metal. As the oxygen content continued to increase (above 1.5–2.0%), weld microstructure was dominated by the acicular and platelet alpha. At very high oxygen content, the alpha platelets formed in colonies, giving a basketweave appearance (Fig. 23.1(f)). Accordingly, fracture surfaces during tensile testing changed from ductile, with classical dimple morphology for the base metal, to brittle for welds with oxygen addition. When the oxygen content was high (Fig. 23.2(f)), the weld fractures as intergranular along colony boundaries or cleavage across colonies, which indicates a complete brittle fracture.
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Denote compatibility in some cases (usually blends)
Amorphorous Polymers ABS ABS/polycarbonate alloy Acrylic Acrylic multipolymer Butadiene styrene Phenylene oxide based resins Polyamide-imide Polyanytate Polycarbonate Polyetherimide Polyethersulfone Polystyrene (general purpose) Polystyrene (rubber modified) Polysulfone PVC (rigid) SAN-NAS-ASA Xenoy (PBT/polycarbonate alloy) Semi-crystalline polymers Acetal Cellulosics Fluoropolymers Icomer Liquid crystal polymers Nylon Polyethylene terephthalate – PBT Polybutylene terephthalate – PBT Polyetherethertetone – PEEK Polyethyene Polymethylbenetene Polyphenylene sulfide Polypropylene
Denotes compatibility
Amorphous polymers ABS ABS polycarbonate alloy Acrylic Acrylic multipolymer Butadiene styrene Phenylene-oxide based resins Polyamide-resins Polyorytate Polycarbonate Polyethermide Polyethersulfene Polystyrene [general purpose] Polystyrene [rubber modified] Polysulfene PVC (rigid) SAN-NAS-ABA Xenoy (PBT/polycarbonate alloy) Semi-crystalline polymers Acetal Cellulosics Fluoropolymers Kremer Liquid crystal polymers Nylon Polyethylene terephthalate – PBT Polybutylene terephthalate – PBT Polyetheretheketone PEEK Polyethylene Polymethyloenelene Polyphenylene sulfide Polypropylene
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Note: This chart depicts compatibility as a complete mixing of materials resulting in a homogeneous bond. Other combinations of materials may be used to create mechanical or partial bonds. Consult your Branson representative if you have a combination not listed in this literature.
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Table 23.2 (a) Weldability matrix for plastics (Courtesy EWI)
Table 23.2 (b) Weldability matrix for metals (Courtesy Unitek Miyachi Corporation) → Weldability of materials
Help
→ Power supply capabilities and optimal solutions
m
iu
in
um
l
Titanium (1870°) Stainless steel (1450°) Platinum (1770°) Phosphor bronze (900°) Niobium (2470°) Nickel (1450°) MP35N (1400°) Molybdenum (2000°) and tungsten (3400°) Iconel, Kanthal, (1400°) Kovar, Nichrome-1500°) Gold (1060°) Galvanised steel (1450°) Copper (1080°) Cold rolled steel (1450°) Brass (900°)
A A,E 4 c A,B A,E 4 c A,B A,E 4 A,B A,E 3 c A A,E 4 c A,B A,E 3 c A A,E 4 c A A,E 4 e A,B
A,E c A,E c A,E c A,E c,d A,E c,d A,E c A,E c,d A d
4 A,B 2 A,E 4 A 2 E 4 B 2 A 2 A
d l, an ze ha e l nt om um on ee a r br d n d st K h e r e e , s s o ll N el ic bd n ni es um l um um er ro ph s on r, N ly ste nl 35 in ni bi d va ke ld pp os as ai P at nc va Mo ng ta al io ic ol r o o t l h I i B C C G P N N G P S T M tu Ko A 2 A,B 4 E 3 A 4 A,E 1 A,B 2 A,B 2 A 2 A 2 A,B 5 A 2 A 2 A,B 1 A,B c A,B a A,B c A,B d A,B d A,B a A e A a A,B A,B a A,B A,B A,B c A,B A 1 A,B 4 E 2 A 4 A,E 1 A,B 3 A,B 2 A,B 1 A 2 A,B 5 A 2 A,B 1 A,B a c A,B a A,B c A,B d A,B d A,B a A,B e A,B a A,B a A,B c A,B A,B A,B A 2 B 4 E 3 A 3 A,E 2 A,B 2 A 2 A,B 1 A 2 A,B 4 A 1 A A,B a A,B c A,B d A,B A,B c A A,B a A b,e A,B a A,B a A A 4 B 3 E 4 A 3 A,E 4 A,B 5 A,B 5 A,B 4 A 5 A,B 2 A c A c A c A d A c A c A c,e A c,d A c A c A c A 3 B 4 E 3 A 4 A,E 2 A,B 2 A,B 2 A,B 3 A 2 A,B c A,B A,B a A,B c A,B d A,B c A,B a A,B e A,B a A,B Weldability A 2 B 3 E 2 A 3 A,E 1 A,B 2 A,B 2 A,B 1 A codes A a A c A d A A a A e A a A a A 3 B 4 E 3 A 4 A,E 2 A,B 2 A,B 1 A,B 1 – Excellent c A c A a A e A a A c A d A 2 – Good A 4 B 4 E 4 A 4 A,E 3 A,B 2 A,B 3 – Fair e A,B e A,B c,e A,B d,e A,B e A,B e A,B b,e 4 – Difficult A 2 B 4 E 3 A 4 A,B 1 A,B 5 – Very difficult c A,B a A,B c A,B d A,B c A,B a A 4 B 2 E 4 A 2 A,E Comments Key A,E c A,E c A,E d A,E A 2 B 4 E 2 A c A d A c A d a – High joint strength is Electrode A 3 B 2 E possible E d E Weldability material b – Use power supply with choice A 1 B closed loop feedback B a c – Low joint strength is A possible Electrode Designing d – Electrode sticking may occur material Comments Electrode Links e – Short weld times may be parts for choice materials necessary. weldability
er pp o c
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ee
st
ee
st
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Beryllium copper (980°) Aluminium (660°)
5 A,B 4 A,B 4 A,B 4 A 4 A,B 4 A 4 A 4 A,B 4 A,B 4 A,E 4 A 4 E 4 B 4 A 4 A 3 A
um lli ry e B A 4 A,B A 4 A,B A 4 A,B A 3 c A A 4 c A,B A 3 A A 4 c A A 4 e A,B A 4 c A,B A 2 a A,E A 4 c A A 2 E A 4 B A 2 A A
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Microjoining in medical components and devices
Materials to be welded (melting point °C)
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Table 23.3 Typical metallic material properties
Aluminum Beryllium Cadmium Chromium Cobalt Copper Gold Hafnium Indium Iridium Iron Magnesium Molybdenum Nickel Niobium Osmium Palladium Platinum Rhenium Rhodium Silicon Silver Tantalum Tin Titanium
Melting point °C
Density
660 1278 321 1857 1495 1083 1064 2227 156 2447 1538 648 2617 1453 2468 3045 1552 1772 3180 1966 1410 962 2996 232 1660
g/cm3
Thermal conductivity W/mK
Specific heat cal/gramC
Coefficient expansion cm/cmC x 10E6
2.702 1.848 8.642 7.2 8.9 8.92 19.3 13.31 7.3 22.42 7.86 1.74 10.2 8.9 8.57 22.48 11.97 21.45 20.53 12.4 2.34 10.5 16.6 7.28 4.5
247 210 97.5 67 69 398 318 23 83 147 80 155 142 83 52 – 70 71 71 150 156 428 54 63 11
0.215 0.436 0.0555 0.107 0.107 0.0922 0.0308 0.035 0.056 0.0312 0.107 0.244 0.0598 0.106 0.064 0.0311 0.0584 0.0317 0.0329 0.0582 0.168 0.564 0.0334 0.0543 0.125
23.5 12 31 6.5 12.5 17 14.1 6 24.8 6.8 12.1 26 5.1 13.3 7.2 4.57 11 9 6.6 8.5 7.6 19.1 6.5 23.5 8.9
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Resistivity @20°C ohmscm 10E-6 2.69 4 7.4 12.9 6.24 1.67 2.3 30.6 9 5.3 9.71 3.9 5.7 6.84 14.5 9.5 10.8 10.6 19.1 4.7 2.30E+11 1.6 13.5 12.8 55
Elastic modulus Gpa
Poisson ratio
69 288 62 253 – 125 77 – 13.5 517 213 43 330 202 106 558 115 171 460 319 163 74 186 50 105
0.33 – 0.3 – – 0.33 0.42 – 0.45 0.26 0.29 0.26 0.37 0.28 0.38 – 0.39 0.39 – 0.26 0.22 0.37 0.35 0.35 0.33
Microjoining and nanojoining
Metals (pure)
3410 1890 419 1852
19.35 5.96 7.14 6.49
1425 1450 1400 1315 1660 1800 1300
8.1 8.0 8.0 8.43 4.43 21.53 6.45
160 31 113 21
11.0 16.2 16.3 11.2 6.5 3.1 10
0.032 0.116 0.0925 0.0666
4.5 8.3 31 5.9
2.15
3.39 17.2 16.2
2.09 0.0317 2.20
8.6 8.8
5.5 26 5.92 44.6
720 740 10.33 178 70-100
398 128 100 99
0.28 0.35 0.25 –
150 193 193 233 114
0.33 0.29 0.32 0.33
21-110*
*21 = martensite *110 = austenite
0.30
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Tungsten Vanadium Zinc Zirconium Metals (alloys) Invar 304 SS 316L SS MP35N Ti 6V4Al PtIr NiTi
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30 um
30 um (a) Base metal
(b) 0%
30 um
30 um (c) 0.5%
(d) 1.5%
30 um
30 um (e) 3.0%
(f) 10.0%
23.1 Micrographs of base metal and laser welds made with various oxygen contents in argon shielding gas.
The suitability (or weldability) of materials to be joined is greatly reduced for dissimilar material combinations, which are commonly required in medical device manufacturing. For example, laser welds between dissimilar Pt and Ti alloys are susceptible to cracking [2]. In the welds, regions with 66–75% Ti, i.e., consisting of primary Ti3Pt and/or Ti3Pt + TiPt eutectic, were found to have hardness above 700 VHN, while regions with 42–66% Ti, i.e., consisting of primary TiPt, had hardness between 400 and 700 VHN. The regions
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(a) Base metal
(b) 0%
(c) 0.5%
(d) 1.5%
(e) 3.0%
(f) 10.0%
699
23.2 Fracture surfaces of base metal and laser welds made with various oxygen contents in argon shielding gas.
consisting predominately of Ti3Pt were very brittle, with cracks formed either during welding or hardness indentation (Fig. 23.3). It is believed that Ti3Pt formed in the laser welds is the phase most susceptible to cracking [6]. While the obvious choice from a welding/joining standpoint is to make a design using one consistent material and practical design, cost and machining ability often dictate the use of many materials within a design. This use of a combination of materials is also evident even within wires and coils now
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(a)
(b)
23.3 SEM micrographs of laser weld between Ti and Pt alloys, with Ti percentages indicated, showing (a) cracks formed during welding in the regions predominantly consisting of Ti3Pt but arrested in the regions predominantly consisting of TiPt and (b) cracks formed during hardness indentation in the regions consisting predominately of Ti3Pt [2].
used in many catheters, pacing leads, and sensors. In many of these components the material is a cored wire, sometimes referred to as a drawn-filled-tube (DFT). Common wire and coil materials in this category include tantalum cored MP35N, silver cored MP35N, platinum clad tantalum, and platinum clad titanium. Materials introduced during welding and joining, such as filler metals during brazing, need to be biocompatible as well. For example, in brazing of zirconia tube to Ti-6Al-4V ferrule in developing an implantable neurostimulator, existing brazing filler metals, such as Ag-Cu-Ti series, cannot be used due to the toxicity of the fillers [7]. Instead, a TiNi-clad braze, consisting of Ni/Ti/ Ni laminated layers, is used since the braze becomes nitinol, which is biocompatible, when the overall Ti/Ni ratio is 50/50. A more thorough review of some medical devices and components, and the associated joining/processing challenges is given in the next section.
23.3
Medical components and devices
Medical components and devices come in many shapes, sizes, and functions. The following section will highlight many current medical devices and the unique joining challenges they present.
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23.3.1 Vascular devices Vascular devices cover a broad range of the medical device industry from catheters and guidewires to balloon angioplasty and stents. While guidewires and catheters are generally single-use devices, stents are obviously permanent implants. For catheters the challenge is joining small diameter wires in a butt weld configuration, often with subsequent sections having a smaller diameter. An additional challenge is that some sections may be made from unique materials. One such example is the joining of a 316L stainless steel wire to a nitinol (nickel-titanium) end affecter. The nitinol can be formed or shaped to provide a certain function, while the 316L provides good torque transmission and lower cost for the long section of the wire. Alignment of the diameters to a center axis and control of the diameter in the weld zone can be difficult. An example of guidewire joining is shown in Fig. 23.4. Balloon catheters require the joining of the polymer balloon to a polymer tube on a guidewire. This can be achieved via direct laser welding or adhesive bonding. Challenges are the consolidation of the thin weld zone to develop high strength, and control of the initial gap between the components during fixturing. An example of a balloon welded to a lumen is shown in Fig. 23.5. Stents are complex devices, typically laser cut from tube stock or welded assemblies of small wires. The stent is collapsed and delivered to the precise arterial location via a catheter and fluoroscopy, and then deployed and expanded in place. These devices pose the challenges of excellent surface and edge finish, control of heat-affected-zones (HAZ), and extreme control of orientation and laser cutting path. A laser cut stent is shown in Fig. 23.6.
23.4 Guidewire for catheter applications. Note the change in diameter and smooth diameter across the joint zone (raw weld unpolished). The materials are NiTi and stainless steel.
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23.5 Balloon catheter joined to lumen expanding a laser cut stent. Photo courtesy of AngioTech Corporation.
23.6 Laser cut stent.
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23.3.2 Pumps and sensors A wide variety of external and internal pumps and sensors are now in use and in development. Insulin pumps and left ventricular assist devices have been in place for many years, with recent years seeing the development smaller scale silicon-based drug reservoirs, MEMs based pressure sensors, and implantable RFID tags. Challenges include the welding and joining of metal and plastic foils, bonding fragile semiconductor layers to each other, assembly of microscopic circuit assemblies, and wire or tab attachments for interconnects.
23.3.3 Cochlear implants Hearing implants are growing in use and are complex multicomponent devices. These units require sub-miniature voice coil and electronic assemblies, in addition to implantable circuit cans, lead wires, and even RF transmit units. Small-scale joining of wires, hermetic seals of enclosures, and assembly of small-scale wire coil assemblies are unique challenges.
23.3.4 Neurostimulators, pacemakers, and defribrillators These devices, while providing unique therapy to the patient, are similar in that they generally comprise a hermetic ‘can’ which encloses the critical electronics, and metal wire leads that provide sensing or therapy delivery. The cans are implanted in a pocket in the upper pectoral region, or low in the abdomen. Wire leads are then threaded through the vasculature to the point where therapy is to be delivered. In the case of cardiac care devices the heart lead is implanted directly inside the right atrium or right ventricle to pace or shock the heart into normal rhythm [8]. In neurostimulators, the lead tip is implanted near the spine, or deep in the brain to control electrical signals that signal pain or to control tremors. Table 23.4 lists typical microjoining processes used to manufacture an implantable pacing system, which includes mainly a pulse generator and leads. The pulse generator includes a battery to generate electricity and circuitry to generate, control, and deliver the pulses, which uses many of the microelectronics packaging processes. The battery, however, is hermetically laser sealed to prevent leakage of chemicals (with a leak rate less than 1 × 10–9 atm-cc/sec). Both the battery and circuitry are inside a titanium case that, again, is hermetically laser sealed to prevent human body fluids entering and causing corrosion. The internal circuits are connected through brazed joints such as Al2O3/Pt-Ir, to an outside connector block encapsulated in a biocompatible polymer such as polyurethane, and interconnected via Pt or Ti wires/ribbons using about 10–20 laser and resistance microwelds. These hermetic cans are Class III devices and designed for long-term
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Table 23.4 Typical microjoining processes in pacemaker manufacturing Interconnecting and attaching
Battery
Circuitry Header Lead Hermetical sealing
Pacemaker, battery cases Lead feedthrough
Wire bonding, soldering, adhesive bonding, resistance and laser microwelding Wire bonding, soldering, adhesive bonding, BGA Brazing, laser and resistance microwelding, plastic joining Laser microwelding and adhesive bonding Laser welding Ceramic/metal brazing
implant in the body. The cans are typically thin walled titanium or MP35N, while the leads are straight wires or coils of MP35N, titanium, or platinum clad cored wires of tantalum, MP35N, or titanium. Ensuring a hermetic seal of the thin walled case is not trivial, as the designs require welding near feed throughs, pins, or ceramic inserts. The case halves must fit together with minimum gap, and the welding process must track a non-uniform shape or seam. The weld penetration must be carefully controlled so as to not overheat nearby structures or penetrate too deeply and damage internal components. The wire leads also provide a multitude of challenges. The lead wires are often dual materials (clad cores), such as Pt-Ir coils coated with silicone rubber, and joined to pins, tubes and other structures that are dissimilar material. An example is the termination of a coil wire for a heart pacing lead. The wire coil is MP35N clad silver and the termination pin is 316SS. The weld is formed using a single laser spot weld. Controlling the melt flow of the silver core can be a significant challenge if improper laser energy or poor targeting of the weld zone is used. An example of a neurostimulator is the BION® microstimulator: a RF powered single-channel ceramic-cased implantable stimulator, which weighs just 0.75 g with an overall volume of only 0.19 cm3 (3 × 28 mm) [9]. It includes a ceramic to metal case, where the ceramic provides a window for RF transmission. The case, 2.44 mm in diameter, is composed of an 11.28 mm long zirconia tube brazed onto a 0.76 mm long Ti cap (cathode) and a 1.98 mm long Ti ferrule (anode). The wall of the ceramic tube and the ferrule is 0.28 mm thick. After the electronic module is inserted into this case, a Ti lid is laser welded to the ferrule for the final hermetic seal.
23.3.5 Radioactive seed implants Implant or internal radiation therapy, or brachytherapy, has been developed as an alternative to the traditional treatments such as radical surgery and
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external beam irradiation for cancer treatments. This technique uses a radioactive substance (e.g., in the form of pellets) sealed in needles, seeds or wires that are implanted directly inside or near the cancer. The cancer cells can be killed by the energy given off as the radioactive material decays. Figure 23.7(a) shows an example of such a seed implant developed for prostate cancer treatment. The radioactive substance is sealed inside Ti tubes by laser microwelding, as shown in Fig. 23.7(b), in this case as laser is well developed and suitable for such precision manufacturing.
23.4
Joint design and process selection
As described in Sections 23.2 and 23.3, medical devices have unique requirements and use a specialized material set for biocompatibility and implant conditions. Many of the material and joint design configurations, however, are common across device types. The following sections will highlight specific assembly challenges and the selection of the microjoining processes best suited to achieve high quality reliable assemblies.
23.4.1 Wire to wire assembly Microjoining challenges manifest themselves at the assembly and joining stage. While small diameter wire and cable has been available for years, joining these small gauge wires to each other in either butt, overlap, or Pt-Ir rod for visualization
Ti spacer 4.4 mm 0.8 mm
Radioactive ceramic bead (I-125 or Pd-103) Ti annulus (a) Laser microweld
(b)
23.7 Radioactive seed implant: (a) schematic design and (b) laser welded assemblies [5].
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parallel (tipping as in a thermocouple) joint configurations and forming operational assemblies is a true challenge. Typical wire diameters may range from large at 0.100″ to small at 0.0005″ for the finest interconnect wires. Typical range for semi-automated assembly is 0.003–0.007″. Wire to wire assembly can be as simple as an overlap weld where both wire ends are flattened in the weld zone, to a uniform iso-diameter butt joint. One of the best processes for joining butt wire assemblies is percussive arc welding. To form this type of weld, the following process steps are completed: 1. Cut the wire ends to form a tapered bullet or square face on the wire. 2. Straighten a short length of wire on each side of the joint. Careful axial alignment is required. 3. Accelerate the cut wire faces toward each other, while simultaneously energizing the wires with opposing voltages and a current limit is also imposed. An alternative is to place a separate electrode between the wire faces with both wires opposite in charge to the center electrode. As the wire faces approach each other or the center electrode, an electric arc will form in the gap, and melt the opposing wire faces. 4. The wire faces are then allowed to forge together. Careful control of the force, the arc time, the closing velocity, and axial alignment are required. Tipping of catheter wire, formation of thermocouples and sensor wires can be completed in a similar manner. The heat source can be an electrical arc, or even the heat of the plasma formed at the tip of a micro-plasma torch. Laser energy can also be directed at the tip where a small portion of the wire end is melted. Natural surface tension causes the molten pool to pull back and form a uniform sphere. This approach works best when the two wires are straight, parallel with each other, and form an isomorphous alloy in the melt zone. An example of catheter tipping is shown in Fig. 23.8. For these types of small-scale wires, arc processes work very well. The small size of the wire helps direct or draw the arc only to the wire tip. Laser weld spots are generally much larger in diameter than the wire. Because the laser can have a non-uniform energy distribution across the weld spot (gaussian), the location of the wire to the beam can be non-optimum. Further, the laser welds and cools over a short timeframe, and minimizes the weld zone. In these types of assemblies this can lead to a marked gradient in grain size from the fusion ball to the parent material in the wire. Arc and plasma processes tend to add more heat, softening the grain transition from the welded ball to the parent wire. This improved gradient in microstructure can result in improved toughness and fatigue performance of the final assembly.
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23.8 Tip forming of a catheter wire. The formed end can be made using laser or arc processes followed by precision grinding to final shape.
23.4.2 Wire to pin assemblies In this application the wire can be straight or in the form of a coil. The wire is then joined or terminated to a pin or other solid substrate, such as a jumper pad on a hybrid circuit or a header block where pacer leads are plugged in. The wire is often in the shape of a coil spring to improve flex fatigue and other mechanical performance attributes. One of the best processes for this type of assembly is resistance welding. In this process oppositely charged electrodes are applied under force to the wire, and current is discharged through the wire to the interface. Resistance to the flow of current causes localized heating, resulting in weld formation where the resistance is highest (usually at the interface of the wire to the substrate). The clamping force used in resistance welding is a key advantage in that it directs the weld location, and consolidates the molten or softened metal to form a concise nugget. Gap is not such a problem as the forces used to clamp the electrode against the parts close any excessive gap. Further, as the materials are heated and soften during the first few milliseconds of the current application, additional joint consolidation occurs. Key challenges are the control of the current path in the often small diameters, especially when direct access for the lower electrode is not possible. Electrode size, position, and force are also issues in that the electrode face is often larger than the wire to be joined. Careful construction of the electrode face and radiusing of the electrode corners can help define the precise weld zone to just a few wire diameters. Disadvantages are the development of small electrodes, electrode wear, and positioning the
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wire-substrate assembly prior to joining. An example of a titanium wire resistance welded to a Ti6/4 terminal block for a pacemaker application is shown in Fig. 23.9. Laser welds are also used for wire to substrate bonds, as they are readily automated and simple to align and target. A disadvantage of the laser weld process is the lack of clamping force holding the wire in place against the surface. If the wire is not in good contact with the surface, there is no driving force to ensure wetting and flow of the wire to the surface. Conversely, the driving force is surface tension in the molten wire, which pulls back from the surface, forming a balled end to the wire. An example of this common defect is shown in Fig. 23.10.
23.4.3 Hermetic sealing of electronic enclosures A key similarity in many medical devices is the need for a small thin walled can or container to hold the necessary electronics, power supply, batteries, capacitors, and other circuit functions. This package must be hermetic to at least moisture ingress, but is typically gas tight to even low volumes of helium. Typical housings in medical applications are MP35N, 316L or, the most common, commercially pure titanium. Hermeticity is required to ensure moisture attack cannot cause corrosion and eventual shorts or opens of the circuits. Note that many hermetic enclosures in use today also incorporate getters for oxygen and moisture, to improve long-term reliability and quality.
23.9 Resistance spot weld of 0.005″ titanium wire to Ti6Al4V terminal block in the header of an implantable pacemaker.
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23.10 Example of surface tension ball defect in a small-scale laser weld. When the wire is not in contact with the surface of the groove, laser energy melts and forms a ball on the wire end.
Regardless of the sealing technique, a minimum of 50% spot to spot overlap is recommended, with 75–80% overlap common practice. Also note that the important aspect is to achieve overlap of the weld zones at the interface of the joint. Overpenetration of individual detrimental spots and penetration does not correlate to hermeticity, it is the continuity of the seal at the interface. This is shown schematically in Fig. 23.11. Three processes have been widely used to hermetically seal thin walled medical cans: • • •
resistance seam welding plasma seam welding laser seam welding.
Resistance welding is an excellent choice when the can materials are resistive as in the case of MP35N or 316L stainless steel. Opposed electrodes are clamped against a flange of the housing and slowly rotated to traverse the length of the package. Current and voltage is applied to form overlapping resistance weld spots. After the first pass along two parallel sides, the electrodes retract, the part rotates 90°, and the process is repeated along the two remaining sides. Circular packages can also be sealed with this technique. Disadvantages of the process are the requirements for a flanged cover, which is often nickel and gold plated. In this manner the resistance weld is actually a resistance braze, forming a eutectic alloy of gold-nickel at the interface. This allows the
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OL
dw
23.11 Overlap condition requirement to achieve full hermetic weld at the interface of the weld joint. In general 50% overlap on the surface is the minimum to achieve uniformity at the interface.
process to run somewhat faster and use less heat input than full fusion welding of the higher melting point MP35N or 316l stainless steel. Plasma welding has hermetically sealed more medical pressure bellows than any other process. Pressure bellows are formed from a stack of preformed thin metal discs, sealed along their periphery. In a manner similar to arc processes for sealing wire tips, the arc is drawn to the sharp edge of the disc, melting and sealing the disc edge. The process produces a weld seam very similar to pulsed laser welding, but with slightly increased heat input. It is an excellent process for sealing flanged enclosures as well. Its disadvantage is the heat input for deeper penetration welds, and the close proximity of the torch face to the weld zone itself. Laser welding is today the most popular and easy to implement for hermetic sealing. Advantages are the ease of ‘point-and-shoot’ line of sight welding capability, and the ease with which automation can track odd package shapes and configurations. Deeper penetration welds can be developed without overheating critical glass or ceramic feed-throughs, and weld cycle time is the shortest for laser welding. Disadvantages are often the initial start up capital cost, training and laser safety requirements, and the need for excellent part fit-up. A general rule of good laser welds is the joint fit-up must have a maximum joint gap of 10% of the thinner material. Excessive gaps cause the laser to punch through one layer causing a hole, or the edges of the seam pull back due to the molten material surface tension. An example of a seam weld with too much gap is shown in Fig. 23.12. Note the smooth balled edges on each side of the seam. An example of a well-formed hermetic seam weld is shown for a pacemaker pulse generator in Fig. 23.13.
23.4.4 Cutting and drilling shapes and holes The laser can also be used to cut and shape tubes to form unique scaffoldlike structures such as stents or abdominal aortic aneurysm (AAA) grafts.
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23.12 Laser welding application showing surface tension effects causing pull back of the weld zone opening a larger gap.
23.13 Laser seam weld in a thin walled titanium enclosure. The surface has been treated post welding to yield a matte finish.
These lasers are unique configurations that develop very small spot sizes with extremely high peak powers. The small spot size minimizes the cut kerf and improves the cut edge quality, while the high peak pulse power fully penetrates and ejects a plasma plume of material. The automation and computer
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control of the laser allows for the development of full 3-D structures. The cut edge quality is critical in these applications, and post-cutting polishing and honing processes are often used to ensure smooth burr free edges. However, the proper selection of pulse rate, table feed rate, cover gas type, and cover gas pressure all contribute to cut edge quality.
23.4.5 Ring to ring or tube to tube assembly There are many instances in medical devices where the mating components are thin walled tubes that slide into one another. Arthroscopic tools, catheters, and electrodes on pacing leads or terminals are examples of this application. Two joining processes are generally used to join these thin wall structures. First and foremost is laser welding. The laser can be used to form single or overlapping spot-welds that join the outer tube to the inner tube. The wall thickness of the electrode can range from 40–75 microns. Here again the advantages of the laser are the ease of point-and-shoot process development, with the disadvantage of joint fit-up. If the inner ring is not in close contact to the outer tube, the laser will simply cut a hole in the outer layer, with little or no mating connection developed. A newer process originally developed for larger-scale automotive applications is magnetic pulse welding. In this process magnetic fields are set up in the inner tube such that the repulsive force to an inner mandrel explosively forces the inner layer outward forming a solid-state bond to the outer layer. The larger scale version of magnetic pulse welding has been used to form flares and expanded regions in cooling tubes and muffler components. The small-scale version is an excellent choice for bonding inner tube structures to outer electrode rings. An advantage of this process is the low heat input, rapid cycle time, and continuous weld over the entire mating surface. Disadvantages are that the application works primarily for round conductive components and requires unique fixturing.
23.4.6 Polymer tube to tube assembly Polymers are used throughout medical devices, especially for single use disposable devices such as catheter sheaths, packaging, and insulators over wire components. One of the simplest and most widely used joining processes is adhesive bonding. Joint design challenges are especially evident where polymeric materials must be joined to metal components. Common practice is to clean and prepare the surface of the metal, and then join a tube over the metal and bond it in place with liquid adhesive. Here again the gap at the interface can define success or failure. In most cases the design extends the bond area to ensure adequate surface area for the bond to occur. Studies have shown that for optimal bonds, a uniform gap should exist to provide a uniform
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adhesive layer thickness. Too thin an adhesive layer yields poor surface coverage and low strength. Similarly, too thick an adhesive layer yields low strength due to failure of the adhesive itself. Careful analysis of the bond failure is critical to development of the proper process and achievement of a robust design. Balloon catheters are another example of polymer-to-polymer joints; this time joined using full fusion processes. A simple heated mandrel/clamp can be used to fuse the joint region, but newer methods using focused infra-red light in the form of a laser or IR lamp are commonly used. The non-contact feature of the infra-red welding process is an advantage for medical grade devices [10]. Cross-contamination of the heated tool approach is a serious defect, and careful tool and heat controls are required to keep the tool face clean and to avoid transfer to the next welded device. CO2 laser systems are common in this application as most polymers have strong absorption bands near the 10.6 micron output of the CO2 laser [11]. In CO2 and some diode welding, the polymer itself absorbs the laser energy, causing the development of a fusion zone. Other wavelength diode lasers and standard ND:YAG systems can also be used, especially when welding onto metallic components.
23.4.7 Polymer tube to solid pin assembly When Nd:YAG is used, its 1064 nm wavelength is usually transmitted through the polymer with very little absorption and heating. In this case the Nd:YAG is focused on an internal metallic mandrel and heated with the laser. The thermal energy then conducts to the polymer layer forming a bond [12]. In this manner the process is referred to as through transmission infra-red welding (TTIR). Additional research has been completed on specific laser absorbing dyes that can be molded into the plastic parts or provided as a thin interlayer. In this application the laser is absorbed at the dye layer, and thermal conduction then heats the surrounding materials to form a weld zone. Note that in all polymer welds, the fusion zone must be held in compression while the nugget cools in order to achieve the best weld strength. The high expansivity and poor thermal conduction of polymers make this compression requirement vital to the development of good welds.
23.4.8 Polymer bags and enclosures Numerous medical devices are fabricated using all polymer systems. These include intravenous fluid bags, drug delivery, microfluidic canisters, pumps, filters, etc. Sealing and welding of thin films to each other and to structural plastics is a common requirement. Early processes used simple heated bar or heated formed tools to compress the weld zone. However, cross-contamination
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from the tool face is a serious issue so non-contact methods were developed. Non-contact includes the aforementioned laser welding, but two of the most popular methods are radio frequency dielectric welding and ultrasonic plastic welding. In RF dielectric welding the process relies on the chemistry of the polymer itself. In this process an oscillating RF electric field is imposed on the weld joint. The polymer chains react to the imposed field, and the dipoles of the polymer attempt to align with the field. The polymer dipoles cannot oscillate in time with the applied field and this hysteresis loss causes rapid heating [13]. Benefits of the technology are fast cycle time, good weld zone control with the applied clamping force, and the tool is cold so transfer and crosscontamination is eliminated. This process is commonly used for final packaging of devices prior to shipping. Ultrasonic plastic welding uses a medium frequency transducer horn and clamping force to develop a weld zone. The ultrasonic energy is applied typically at 20 kHz or 40 kHz depending on the component size, shape and wall thickness. Similar to RF dielectric heating, the ultrasonic energy vibrates the polymer chains, causing rapid heating and development of a weld zone [14]. Here again the applied clamping force forges the molten zone together to form an excellent seal and weld. This process can be completed on virtually all thermoplastics.
23.5
Testing and verification
Microjoining applications are often tested in manufacturing using smallscale tensile testing equipment. During pull testing, the actual joint is often exposed to a complex stress state, resulting most often in tensile shear of the weld zone. It is important in small-scale welds to consider the effect of the weld process on the raw parent material strengths. As a general rule, full fusion welds yield approximately 60% of the raw parent material strength, in tension. In practice, the ‘pull test’ can apply off axis loads and stresses, leading to failure at lower than expected loads [15]. Peel testing is especially stressful, as the load is focused on a narrow zone of material, which tends to tear and unzip along the weld interface. Even though the weld zone formation is the same, peel test samples may fail at 25% of the load applied in direct tension. A schematic of this effect is shown in Fig. 23.14. As the stresses are effectively applied to smaller and smaller areas of the joint or weld zone, the test strength decreases. The point here is to carefully document the testing procedure and provide adequate controls to apply the test force in a consistent and uniform manner. In practice, small-scale joints typically have two failure modes. In the first, insufficient mixing across the interface leads to a small interface area to carry the applied load, so the joint fails. In this case the mating faces of the
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Tensile = 100 lb
Butt weld
Lap weld
Shear = 75 lb
Peel = 50 lb
Lap weld
Lap weld
Torque = 25 lb
23.14 Schematic of test conditions and the effect of loading condition on measured weld strength for a uniform area weld zone.
joint show very little transfer of material and it is obvious to the process and quality engineer that welding was insufficient. Overwelding is also an issue. Weldments generally exhibit a horseshoe form when plotted as strength versus power. At low power the weld has no strength due to lack of weld zone development at the interface. At some point, optimum welds will yield the highest strength. This occurs where the weld zone at the interface is optimized relative to the loss of parent material strength. With the application of too much power, the weld zone penetrates deeper and overheats the parent materials, causing parent material failure adjacent to the weld zone. This is often thought of as a ‘good’ weld, even though the parent material has been seriously degraded from its original condition. The reader is encouraged to optimize the weld zone and parent material performance, and neither under or over weld these critical assemblies.
23.6
Summary and future trends
Microjoining is a critical and enabling technology for manufacturing of medical implantable components and devices. The small scale of many emerging designs requires a deeper understanding of joining technologies and thin materials so that the proper joining method can be applied. Lack of standardization in polymeric materials will slow the development of many new devices. Polymeric materials generally must meet a melt flow index and density, but joining properties can be greatly affected by flow enhancers, regrind content, lubricants, etc. Material purchased to similar
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specifications from multiple vendors often does not perform with equal results. Polymeric materials must approach the compositional rigor imposed by metals standards to enhance their use in medical devices. Joint designs for metals and plastics should also be developed and shared across the industry. Guidelines for joint-fitup, gap tolerance, weld penetration, and heat input would help improve overall quality and reliability of medical devices. The weldability data given in Tables 23.2(a) and (b) should be expanded upon, especially for dissimilar material combinations. Further, that database should be driven by welding technology, as some small-scale laser welds cool so rapidly, the process can violate what the handbooks would indicate as weldable. Biocompatibility and corrosion performance of joints should be studied and a database of performance presented to the medical industry. Many materials could be used for certain applications, but since compatibility is governed from a device level, some material choices are neglected because they must replicate what has been used in the past. This regulation limits the use of new materials until an overall design is approved. Lastly, a key future trend in medical devices is a push toward more automated assembly processes. The current practice of manual assembly places high requirements on training and inspection processes. The use of automation has the potential to improve overall quality and reliability by removing the human element and inherent variability of manual processes.
23.7
References
1. FDA Classifications for Devices www.fda.gov 2. B.D. Ratner, A.S. Hoffman, F.J. Shoen, J.E. Lemons, Biomaterials Science, Academic Press 1996. 3. G.F. Carter, Principles of Physical and Chemical Metallurgy, American Society of Materials 1979. 4. ASM Handbook, Properties and Selection of Non-ferrous Alloys and Special Purpose Materials, Vol. 2, Copyright 1990, ASM International. 5. X. Li, J. Xie, Y. Zhou, ‘Effects of Oxygen Contamination in the Argon Shielding Gas in Laser Welding of Commercially Pure Titanium Thin Sheet’, J. Mater. Sci., 40, 2005, 3437–43. 6. N.J. Noolu, H.W. Kerr, Y. Zhou, J. Xie, ‘Laser Weldability of Pt and Ti Alloys’, Mater. Sci. Eng. A, 397, 2005, 8–15. 7. www.guidant.com 8. Peytour, Berther, Barbier, Revcolevschi, J. Mater. Sci. Lett., 9 (1990), 1299–1131. 9. www.draximage.com 10. Robert A. Grimm, IR Welding of Polymers, MDDI May 2001. 11. P.C. Heimenz, Polymer Chemistry, Marcel Dekker Inc. 1984. 12. V.A. Kagan, R.G. Bray, W.P Kuhn, ‘Laser Transmission Welding of Semicrystalline Thermoplastics’, Journal of Reinforced Plastics and Composites 2002:21:1101.
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13. K. Mistry, ‘Plastic Welding Technology for Industry’, Journal Assembly Automation September 1997, Vol. 17, Issue 3, pp. 196–200. 14. Ultrasonic Welding of Plastics used in Medical Devices, Welding Technology Institute of Australia, TGN-MS-04, Rev 0, April 2006. 15. G.E. Dieter, Mechanical Metallurgy, McGraw Hill Company 1986.
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24 Hermetic sealing of solid oxide fuel cells M B R O C H U McGill University, Canada and R E L O E H M A N Sandia National Laboratories, USA
24.1
Introduction
Globally, the demand for electricity is expected to double in the next 20 years, with the bulk of that demand coming from developing nations whose electric power infrastructure is modest or nonexistent. Solid oxide fuel cells (SOFCs) could play a major role in meeting this worldwide expanding demand for energy by providing efficient, environmentally friendly electrical energy. For instance, fuel cell use in transportation will reduce our dependence on fossil fuels, which will limit economic dislocations from the inevitable rise in the cost of transportation fuels. Because of the importance given to this technology, there has been a major worldwide investment in fuel cell R&D to bring them to more immediate commercialization. The targeted areas for the use of SOFCs are numerous, including stationary units, transport vehicles, and military applications. Among the many technical challenges in SOFC development, the fabrication of hermetic seals that function for long periods at service condition is the main limitation to manufacturing reliable units. A SOFC is a low pollution electrochemical device that converts the chemical energy in a fuel (natural gas, coal-derived synthetic gas, reformed gasoline or diesel) into electrical energy without combustion [1]. A typical planar SOFC possesses a sandwich design composed of a porous perovskite cathode, a dense yttria-stabilized zirconia (YSZ) electrolyte and a porous nickel (Ni)YSZ anode. Figure 24.1 presents a schematic representation of the SOFC stack configuration. Figure 24.2(a) presents a schematic, with typical dimensions, of one of the earliest cell designs, namely electrolyte supported cells [2]. To reduce the total electrical resistance of the cell, developers have tended to reduce the thickness of the electrolyte while increasing the anode thickness to provide mechanical support, hence the name, anode supported cell. The thickness of the electrolyte in current cells is typically around 20 µm while the thickness of the anode is 1000-2000 µm. The thickness reduction of the electrolyte reduces significantly the resistance of the cell and improved performances are observed. 718 WPNL2204
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Current flow
End plate Anode Electrolyte matrix Repeating unit
Cathode Bipolar separator plate Anode Electrolyte matrix Cathode Fuel flow Oxidant flow
24.1 Schematic representation of the SOFC stack configuration.
50 µm
Cathode
200 µm
Electrolyte
50 µm
Anode
Anode
(a)
(b)
Cathode Electrolyte
50 µm 20 µm
1000–2000 µm
24.2 Schematic representation of (a) electrolyte supported cell and (b) anode supported cell.
To ensure sufficient oxygen atom mobility across the electrolyte, SOFCs are operated in a temperature range of 700 to 1000 °C. An electric current is generated when molecular oxygen dissociates at the cathode and diffuses as O2– to the anode where it reacts with hydrogen to produce water, carbon dioxide, heat, and electrons. The electrons given up by the reduction of O2– travel through the external circuit to do useful work. This chemical reaction releases about 1 volt during operation at 1000 °C [3]. Similar to battery technology, a stack configuration is required since it generates a significant increase of the output power when compared to a single cell [1]. The absence
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of the combustion process in fuel cells eliminates the formation of pollutants including NOx, SOx, hydrocarbons and particulates, and significantly improves electrical power generation efficiency (60–70% for SOFCs vs. 30–40% for combustion engines) [1]. Thus, to effectively market and to obtain acceptable performance from SOFCs, there is an urgent need to develop reliable, robust joining techniques. However, prior to presenting the current research directions in sealing SOFCs, a brief description of the different cell materials will be presented to illustrate the challenges faced in joining these dissimilar materials under the severe service conditions imposed on the seals.
24.2
Materials involved in the fabrication of a SOFC
24.2.1 Cathode material The cathode is the layer in contact with the electrolyte that is exposed to air or oxygen. The cathode performs three functions [4]. First, it provides a mechanism for breaking the covalent bond of the oxygen molecule. Second, it accepts electrons from the external circuit and carries them to the cathode reaction site where oxygen ions are formed. Third, it transports the oxygen ions to the electrolyte interface. The commonly used materials have the perovskite crystal structure, and are typically a lanthanum manganite composition doped with rare earth elements such as Sr or Ce. These perovskites have matched expansion coefficients (CTEs) to the electrolyte, and are mixed ionic and electrical conductors.
24.2.2 Electrolyte material The electrolyte is a fully dense ceramic layer that must possess high ionic conductivity but no electronic conductivity to avoid current leakage, and it must be chemically and mechanically compatible with the other components of the cell. The thickness of the electrolyte is minimized to reduce internal cell resistance. Yttria stabilized zirconia (YSZ) and cerium oxide (CeO) are the two main materials considered for electrolyte applications. The addition of yttria to zirconia increases the ionic conductivity by increasing the concentration of oxygen vacancies. The contribution of grain boundary conduction is important in YSZ and becomes prominent at lower temperature. This combination explains the efforts in developing nanostructured YSZ electrolyte for lower operating temperature SOFCs [5]. Cerium oxide has been considered since is has a higher ionic conductivity than YSZ, which allows operation at lower temperatures. However, CeO is a mixed electronic conductor at low oxygen partial pressures and is more susceptible than YSZ to being reduced at the anode interface [6]. Development of other electrolyte materials in the perovskite family is presented by Fergus [7].
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24.2.3 Anode material Ni-YSZ cermet is the most commonly used anode material. The three-phase composite of Ni, YSZ, and porosity provides electrical conductivity through the metallic phase, ionic conductivity through the oxide phase, matching thermal expansion with the electrolyte and sufficient porosity providing increased paths for transport of the fuel. The material is engineered so the relative volume fractions of YSZ and pores adjust the CTE, while the YSZ particles provide a support to separate the Ni agglomerates. One of the drawbacks of the Ni-YSZ cermet is the catalytic formation of graphite from hydrocarbon gas, which deposits in the pores and may block the anode reaction [8-10]. The cermet is also susceptible to sulfur poisoning, where sulfide deposits reduce the anodic reaction [9]. Therefore, to prevent the loss of efficiency, novel single-phase anode materials, possessing both electrical and ionic conductivity are in development. These anodes are based on Ti-, Cr-, Mn- or Fe-chemistries. Fergus presents a thorough review of the development in this field in [11].
24.2.4 Interconnect material The primary function of the interconnect is to make electrical contact between the cathode and the anode and to provide a physical barrier between the oxidant and the reducing fuel atmosphere. Thus, the interconnect experiences the most severe operating conditions. The interconnect material must be electrically conductive, fully dense, it should have matching thermal expansion to both electrodes, and inertness in oxidizing and reducing environments to prevent the formation of phases that increase the electrical resistance or cause mechanical degradation. Candidate materials for the interconnect are constrained by the targeted operating temperature of the cell. For SOFC operating at 1000 °C, rare earth doped LaCrO3 materials are used [1, 12]. Cells operating between 900 and 1000 °C use interconnects made from Ni-based superalloys, while cells operating at 800 °C or lower use ferritic stainless steel materials, which are less expensive, easier to manufacture, and have CTEs closer to those of the cell materials. There is a trend toward reduced SOFC operating temperatures with the development of improved cell materials, which has led to significant attention to metallic interconnects. Oxide chemistry, reaction rates in both oxidizing and carboncontaining atmospheres, surface oxide growth, stability and vaporization at high temperatures, electrical resistance, and thermal expansion are crucial factors to consider in the selection of interconnect materials. A thorough review of potential metallic interconnects and a comparison of their properties are presented by Fergus [13].
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24.3
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Different approaches for sealing SOFCs
Similar to battery technology, a stack configuration is favored in planar SOFCs since an increase of the output power can be obtained [1]. New and advanced sealing materials and processes are required for manufacturing SOFC stacks with long term reliability. The seals must be inert to the materials that make up the SOFC stack, they must possess sufficient fluidity and adhesion to form the seal but with appropriate creep resistance to prevent failure or leakage during service, and they should possess a CTE matched to those of the cell materials to prevent formation of residual stresses from thermal cycling during manufacture and operation [14]. Figure 24.3 presents a schematic representation of different sealing approaches based on whether or not the seal CTE matches those of the other components of the cell. The following sections describe the status and advances of some of the most promising techniques: glass and glass ceramic seals, glass composites, metallic brazes, air brazes and compressive seals.
24.3.1 Rigid seals The development of rigid seals has received significant attention since they can provide the mechanical support required to obtain self-standing stack. Rigid seals do not require the presence of a load frame surrounding the cell, which ultimately lightens the stack and make them more attractive for the transportation industry. The main limitation of rigid seals is that they cannot relieve residual stresses by flow or creep. The CTEs of cell materials differ at least a little bit. Further, thermal gradients are inevitable during heating and cooling and even during operation. Thus thermal stresses will always be present and the seal has to adapt to those stresses to prevent mechanical failure.
Joining process for sealing SOFCs
Unmatched CTE
Compression seals
Matched CTE
Compliant seals
Glasses Glass ceramics Glass composites
Reactive brazes Air brazes
24.3 Schematic of some joining techniques for sealing SOFCs based on degree of CTE match.
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Glass and glass ceramic seals Seals using conventional glasses Glass sealing is a well known joining technique that has its roots in techniques of early glass blowers who used it in the decoration of glassware and in the fabrication of silver mirror in the 15th and 16th centuries. Since then, continual improvements in the technology have led to the development of examples such as electrical feed thrus in the 19th century, to current applications such as electrochemical devices and hermetic packaging. The earliest approach to hermetic SOFC seals used glass sealing. Glass seals were operated for more than 100 hours without significant degradation in stack performance [15]. Multiple glass systems have been developed and are available for use in SOFC seals but to date none of them has shown the range of properties necessary to meet cost and reliability requirements. According to Stevenson [16], specific families of glasses investigated showed the following trends. Phosphate-based glasses possess intrinsic low CTE and low strength. Borate-based glasses have low softening temperatures but they have been thought by some to be volatile in the high temperature steam atmosphere characteristic of SOFC operation and their reliability for stack operating temperatures above 800 °C, has been questioned. However, the suitablility of borate glasses for SOFC sealing is still being studied and whether or not they have the required properties is still an open question. Wide varieties of silicabased glass chemistries exist and represent the potential candidates available. Other examples within these glass families have properties outside the presented ranges and also represent potential material system for sealing SOFCs. Minimizing the reaction between the glass seals and the SOFC materials they contact during manufacture and operation is a main concern, but other criteria used in the selection of a glass for a seal system are the glass transition temperature, Tg, and the CTE. The glass transition temperature and softening temperatures should be sufficiently low to allow flow, wetting and adhesion to the parent materials. However, while spreading, the glass seal must retain sufficient mechanical strength and creep resistance to maintain mechanical integrity of the stack. As much as possible, the CTE of the glass seal should match the CTEs of the materials to be joined, such as the electrolyte and the interconnect, to minimize the interfacial stresses generated during start-up and cool down of the stack. Geasee et al. [17] proposed a selection window for glass seal systems. According to their analysis, the glass transition temperature should be between 600 and 740 °C and the glass should possess a CTE ranging between 10.4 and 12 × 10–6/°C. Figure 24.4 shows a small collection of silica-based glass systems, originally presented by Fergus [18], that agree with the criteria proposed by Geasee. In this list, the promising compositions are bariumcontaining glasses, some of which are known to have large CTEs.
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Coefficient of thermal expansion (× 10–6 /°C)
12 11.8 11.6 11.4 11.2 11 10.8 10.6 10.4 600
620
640 660 680 700 Glass transition temperature (°C)
720
740
Ba aluminosilicate [19]
BaCa borosilicate [17]
BaMg silicate [24]
Ba boroaluminosilicate [20–23]
BaMg borosilicate [17]
BaZn silicate [24]
24.4 Glass compositions satisfying the criteria of glass transition temperature and CTE suggested by Geasee et al. [17].
Glasses that have suitable thermal and mechanical properties for high temperature SOFC seals usually are not very good glass formers. Thus, they are susceptible to crystallization when exposed to high temperatures, which can lead to changes in properties such as strength, CTE, and glass transition temperature. The crystallization of a glass produces a glass-ceramic, usually composed of a residual glassy matrix and one or more crystalline phases. The development of glass compositions whose properties do not change after crystallization is very improbable. A better approach is to engineer the properties of the glass ceramic itself and not to be very concerned with the properties of the parent glass beyond initial flow, wetting, and lack of undesirable reactivity. Glass ceramics generally require a two-stage heat-treatment composed of a lower temperature nucleation stage and a higher temperature crystallization stage. The phases that crystallize may be metastable, or they may be the equilibrium phase(s) given by the relevant phase diagram. Metastable phases may be susceptible to further change during high temperature exposure as the glass ceramic tends toward the equilibrium composition. If the crystalline phases are thermodynamically stable in the temperature range of interest, the seal composition should not change over time at temperature. In the latter case, a well-defined heat treatment to nucleate and grow crystalline phases in the glass will result in the formation of a glass-ceramic seal possessing stabile high temperature properties. Such an approach has been studied for several glass systems [17–18, 21, 24]. One of the limitations of this approach is that the control of the seal properties is generally restricted by the formation
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of thermodynamically favorable compounds, which limits the window for engineering particular seal properties. In principle, the glass ceramic can be formed by crystallizing a preplaced glass preform during initial heatup of the SOFC stack. Exposing a SOFC stack to the heat treatment schedule used to develop the in situ glass-ceramic composite seal may favor the formation of a thick reaction layer, which can result in the formation of large residual stresses on cooling. In addition, the phases formed during the heat treatment might induce mismatch in CTE with the rest of the seal components. As mentioned, several glass ceramic systems for sealing SOFCs have been reported. Only some selected systems are presented here to illustrate the main features of this joining approach. Meinhardt et al. [25] developed a glass composition range composed of BaO, SrO, CaO, MgO with Al2O3, B2O3 and SiO2 as glass forming agents possessing matching CTE with YSZ under its crystallized form since the crystals are Ba-silicates and Baaluminosilicates, which possess high CTE. Table 24.1 presents the first phase formed during the crystallization of various silicate glass systems [26] and the respective CTE of the crystal. As observed, the CTE of a glass ceramic is generally different from that of its base glass. The change in CTE depends upon the phase and its volume fraction. As observed, the precipitation of monoclinic celsian (named monocelsian) crystals for the BAS and the SAS glasses causes a significant reduction in CTE, which may be deleterious to cell performance. Because the initial crystallization may not give the equilibrium phases the microstructure of the glass ceramic may change with time. Figure 24.5 presents the evolution of the crystalline phases, measured by X-ray diffraction, for two BaO-Al2O3-B2O3-SiO2-based glass ceramics [23]. The experiments were carried out at 800 °C and the results show that the volume fraction of monocelsian increases with time, which lowers the CTE Table 24.1 Primary crystalline phase formed during crystallization of various glass systems and their respective CTEs Glass system
Primary crystal
CTE (×10–6/°C)
Activation energy (kJ/mol) [26]
References
BAS
BaAl2Si2O8
473–560
27
BCAS SAS
BaSiO3 SrAl2Si2O8
259 473–560
28 27
MABS
MgSiO3
7–8 (Hexacelsian) 2–3 (Monocelsian) 5–7 (Orthocelsian) 9–13 8–11 (Hexacelsian) 3 (Monocelsian) 5–8 (Orthocelsian) 7–9
420
29
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Intensity (A.U.)
726
500 450 400 350 300 250 200 150 100 50 0
Hexacelsian Monocelsian BaSiO3
0
200
400 600 Time (hours) (a)
800
1000
350 Intensity (A.U.)
300 250
Monocelsian
200 150 100
Hexacelsian
50 0 0
200
400 600 Time (hours) (b)
800
1000
24.5 Explanation of change in CTE of two glass-ceramic systems as a function of the evolution of the crystalline phases [23].
of the seal. The glass ceramic illustrated in Fig. 24.5(a) had a reduction of CTE of 1.6 × 10–6/°C, changing from 9.9 × 10–6/°C to 8.3 × 10–6/°C during the 1000 hours of heat treatment. For the system presented in Fig. 24.5(b), the CTE decreased from 11.5 × 10–6/°C to 7.3 × 10–6/°C for the same heat treatment. The volume fraction of residual glass was not reported. Other examples are given in Table 24.2, which presents four glass-ceramic compositions, the heat treatment imposed to study the stability at high temperature, the phases observed after heat treatment and the measured change in CTE. Unfortunately, the volume fractions of the crystal phases were not reported. The data illustrate the general concept that the microstructures of glass ceramics evolve during long term exposure to high temperature through crystallization and phase transformation, and the change in microstructure will alter the seal properties. Moreover, it illustrates the difficulty of engineering stable microstructures for prolonged service periods. It is also worth mentioning that glass compositions containing Mg are prone to form crystals possessing low CTE, like cordierite, rendering the glass seal inappropriate for SOFC applications [18].
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Table 24.2 Change in CTE after long-term heat treatment Glass chemistry (mol%) B 2O 3
SiO2
30
5
50
45
5
50
35
10
16.7
33.3
36.1
10.3
21.5
30.1
Notes:
La2O2
ZrO2
5 2
Heat treatment CaO Time (hrs)
Temp (°C)
15
24
900
24
900
1000
800
1000
800
(1) Monocelsian crystal: BaAl2Si2O8 (2) Hexacelsian is a polymorph of monocelsian
Crystals
BaAl2Si2O8 Ba5Si8O21 BaCaSi3O9 BaAl2Si2O8 BaSiO3 Ba2Si3O8 Monocelsian1 Hexacelsian2 Monocelsian1 Hexacelsian2 BaSiO3
CTE before heat treatment (×10–6/°C)
CTE after heat treatment (×10–6/°C)
Ref.
10.2 (50–700 °C)
9.7 (50–700 °C)
19
10.7 (50–700 °C)
13.2 (50–700 °C)
19
11.1 (RT–700 °C)
9.3 (RT–800 °C)
23
10.6 (RT–700 °C)
6.4 (RT–800 °C)
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BaO Al2O3
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Crystallization kinetics also are a consideration when using the glassceramic seal approach. In some glass ceramics, the residual glass has a lower viscosity than the parent glass, implying that if crystallization occurs prior to the formation of a proper interface with the cell components, poor adherence and a potential site for leakage will be present. Table 24.1 illustrates that that the activation energy is function of the glass composition, and that the activation energy of crystallization is a good indicator of the rate of crystallization of the glass. Barium calcium aluminosilicate glasses will quickly crystallize and thus should be more stable at high temperature. Since in most glass systems, the composition of the crystalline phase differs from the liquid, the rate of crystallization is controlled by the rate of diffusion through the liquid phase and the number of nucleation sites. By increasing the number of nucleation sites, e.g. by using slow heating rates, the crystallization peak temperature will be reduced. The addition of nucleating agents is one of the approaches to engineer the crystallization kinetics of the glass phase. If properly selected, the nucleation agent can be used to speed crystallization kinetics of other particular crystalline phases. Table 24.3 presents the effect of various nucleating agents on the activation energy of crystallization for a Mg aluminosilicate glass, which has a crystallization activation energy of 420 kJ/mol (see Table 24.1). In this particular glass composition, Zr also acts as a network modifier, explaining the lower activation energy [30]. Microstructure investigations have also shown the positive effect of Cr as a nucleating agent by accelerating the crystallization of other phases than cordierite (Mg2Al4Si5O18), a crystalline phase possessing a low CTE (1 × 10–6/°C). The previous discussion presented the potential of using Ba-containing silicate glasses for sealing SOFCs since the crystallization can be controlled to match the CTEs of the other components of the cell. The main drawback of this family of glasses is its high reactivity with the interconnect materials. Ba-containing silicate glasses form thick BaCrO4 reaction product at the interface [31–32]. Mg-containing silicates form MgCrO4 [33] and calciumcontaining silicate glasses form Ca3Cr2Si2O8 crystals [34], while in contact with ferritic stainless steel interconnect material. The reaction products have Table 24.3 Activation energy for crystallization of Mg aluminosilicate glasses containing various nucleating agents [18] Nucleating agent
Mol%
Activation energy for crystallization (kJ/mol)
ZrO2 TiO2 Cr2O3 Ni
2 1.5–7 0.6 2
310–410 310–450 500 575–620
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a negative effect on the integrity of the joints, which makes the long-term reliability of a SOFC stack questionable. Invert glasses Invert glasses are a family of glasses whose properties differ from those of conventional glasses due to the nature of the glass network. The glasses, developed by University of Missouri-Rolla, have shown great potential for SOFC sealing applications. An invert glass is a glass with <45mol% of glass-forming oxides (SiO2 + B2O3). Figure 24.6 presents a schematic representation of the glass network in an invert glass [35]. Their structures are less highly crosslinked than conventional glasses. When invert glasses crystallize they tend to form pyro- and orthosilicates, while conventional glasses are prone to crystallize silicate and metasilicate phases. Reis and Brow [36] reported compositions of invert glasses with Tg-CTE properties that match the criteria proposed by Geasee (see Fig. 24.4). The CTE for crystallized invert glasses remained stable or slightly increased after
(Na, K, Ca, Sr, Ba) SiO4 tetrahedron Bridging oxygen ion Conventional modifying ion
Si-ion Non-bridging oxygen ion
24.6 Schematic representation of an invert glass network.
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crystallization of pyro- and/or orthosilicates. No significant changes in CTE and crystal volume fraction were observed for these glasses after exposure of up to 670 hours at 750 °C. In addition, these glasses form strong bonds with Cr-steel substrates and, compared to BaO-containing glasses, no significant interfacial reaction products formed at the glass-metal interface after an exposure of 1440 hours at 750 °C [36]. Glass composites Glass composites are another useful method for sealing SOFCs. The microstructures of glass composites are similar to those of glass ceramics, i.e., that the material is composed of a glassy matrix with dispersed crystalline particles. The main advantage of glass composites resides in the fact that the crystalline and glassy phases can be chosen independently, rather than being restricted to the metastable or equilibrium phases available from a given glass composition. Thus, the composition, particle size distribution and volume fraction can be varied over wide ranges to engineer the properties of the composite [38]. The chemical and mechanical properties of composite seals (viscosity, adhesion, reactivity with parent materials and CTE) can be engineered independently through the careful selection of the additive phase, which is not limited to the equilibrium phase relationship derived from the glass chemistry, as in the glass-ceramic approach. In addition, the flow of glass composites can be predicted by equations such as eqn 24.1, which allows values to be tailored precisely [38]: 2 κφ η = 1 + 1 – ( φ / φmax )
(1)
Where η, κ and φ are the viscosity, 1/particle size and particle packing density (volume fraction), respectively. From eqn 24.1, the viscosity of a glass composite is expected to increase with the reduction of the particle size and/or the additive volume fraction. Figure 24.7 shows the effect of glass/ ceramic powder ratio on composite flow and viscosity. As depicted, increasing the volume fraction of additive progressively reduces spreading up to the point where it is totally inhibited. The CTE also can be engineered using composite theories. In particular, eqn 24.2 can be used to calculate the composition that adjusts the CTE of a composite material to the desired value: αc = αg + (αa – αg) Va
(2)
where αc, αg and αa are respectively the coefficients of thermal expansion of the composite, the glass matrix and the additive and Va is the volume fraction
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731
(b)
(c)
24.7 Effect of glass/ceramic powder ratio on composite flow and viscosity, (a) pure glass, (b) 85/15 glass/powder ratio and (c) 65/35 glass/powder ratio.
of additive. This equation predicts that in an inert system, the CTE of a composite varies linearly with the volume fraction of additive. Figure 24.8 illustrates the agreement between the predicted CTE of a glass composite (dashed line) and the measured CTE (solid points) for two inert glass composite systems. Brochu et al. [39] have shown that with the addition of ZrO2 particles, desired control of the fluidity of a borate-based glass can be achieved and that the reactivity between the glass and the cell materials is inhibited, leaving a clean and reaction-free interface. They also demonstrated that a finer particle size of additive is advantageous since a lower volume fraction is required to control the flow and the glass properties and that Tg and CTE do not vary during the sealing cycle [40]. Two vol% of nano-YSZ additive would have a similar effect on viscosity as 10 vol% of micron scale YSZ. Nielson et al. have demonstrated similar results using MgO additives in a sodium aluminosilicate glass system [41]. Their MgO-glass composite showed constant low levels of gas leakage even after multiple thermal cycles and after more than 5000 h run time in stack testing. Metallic particles are also attractive additives for glass composites. Loehman et al. have shown that Ag particles are inert in borate-based glasses and thus are a useful alternative to the ceramic additives [42]. Beatty also reported successful uses of Ag additives in an alkali borosilicate glasses to control CTE mismatch [43]. Loehman et al. [42], showed that Ni-additive (APS 7 microns) in an invert silicate glass was unreactive after long term exposure at SOFC service temperature and thus its possible utility for adjusting the CTE of composite glass seals to match the cell components. Figure 24.9 shows the long-term stability of two invert glass-Ni composites for exposures of up to 500 hours at 800 °C. As depicted, no significant variation of CTE
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12 10 CTE (× 10–6 / °C)
Invert glass + Ni additive 8 Borate glass + YSZ additive
6 4 2 0 0
5
10
15 20 25 Additive volume fraction (%)
30
35
24.8 Predicted and measured CTEs of two inert glass composite systems.
11.5 11
CTE (× 10–6 /°C)
20 vol% Ni 10.5 10 9.5
5 vol% Ni
9 8.5 8 0
50
100
150
200 250 300 Time (hours)
350
400
450
500
24.9 Effect of time at 800 °C on the CTEs of two invert glass/Ni composites.
was measured. The seals were also subjected to thermal cycling between 25 °C and 800 °C and no deterioration of the interface was observed. Metallic brazes Brazing is one of the simplest joining processes, based on melting and solidification of an inert filler metal between the parent materials. Brazing of ceramics or ceramics to metals, for example in SOFC stacks, requires the
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addition of an active element into the filler metal to destabilize the covalent or ionic bond of the ceramic and form an intermediate reaction layer, which is wetted by the filler metal. This modification of the process is called reactive brazing. The most common active brazing elements are Ti and Zr. Compared to the glass seal approach, the lower stiffness of the metallic brazing alloy allows the possibility of absorbing residual stresses created during thermal cycling of the component by plastic deformation. Because the SOFC service environment is both oxidizing and reducing, excellent oxidation resistance of the filler metal will be critical. Thus, most alloys studied for this application are mainly composed of silver and gold. The presence of an active element that promotes wetting of the ceramic surface is problematic since it will significantly react in an oxidizing environment. Weil et al. [44] studied the microstructures and oxidation resistance of joints between the YSZ electrode and a SS430 interconnect brazed with two active brazing alloys, namely gold ABA (96.4%Au; 3%Ni; 0.6%Ti) and Nioro ABA (82%Au; 15.5%Ni; 1.75%V; 0.75%Mo), at 800 °C for 200 and 50 hours, respectively. Significant changes in the microstructures were observed. Figure 24.10 presents the microstructures of both alloys after the oxidation tests. Figure 24.10(a) shows significant diffusion of Fe, Ni and Cr from the interconnect toward the gold ABA filler metal. A chromium oxide layer of ~4 microns in thickness at the ceramic interface and a number of FeNi rich precipitates were observed. Pressure tests showed weak interfacial strength since leakage was observed during the test. Figure 24.10(b) depicts the interface brazed with Nioro ABA. A 50-micron thick Fe-Ni oxide layer was detached from the ceramic interface and significant dissolution of the interconnect was observed. Pressure tests showed instantaneous leakage of
Fe-Ni oxide Cr2O3 SS430
YSZ
SS430
YSZ Filler metal
Filler metal 50 µm
50 µm
(a)
(b)
24.10 Microstructures of YSZ/SS430 interconnect brazed interfaces after oxidation test at 800 °C in air (reproduced with permission of Elsevier) [44].
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the seal. The experiments on reactive brazing of SOFCs showed that this joining route is not suitable for manufacturing SOFC stacks as significant diffusion between the filler metal and the interconnect and internal oxidation of the filler metal readily occur. Air brazes The application of air brazing to seal SOFCs was first investigated by the Pacific Northwest National Laboratories (PNNL), and has been advanced as an alternative to the conventional reactive brazing used to join ceramics [45]. The bonding mechanism in air brazing is based on the formation of a thin oxide layer at the surface of the metallic substrate, which can then be wetted by a molten metallic filler metal [14]. Weil et al. used the Ag-CuO air brazing system to join an anode-electrolyte bi-layer to a FeCrAlY interlayer [14]. They reported a direct correlation between the wetting angle of the braze and the rupture strength of the seal for a narrow filler metal compositional band at 4 mol% CuO. The air brazed seals sustained 40 thermal cycles with a heating rate of 75 °C/min without any reduction in the strength. Failure of the assembly occurred through the bi-layer, indicating that the adhesion strength of the seal was not limiting. The strengths before and after thermal cycling were reported to be four times higher than for a reference seal composed of a barium calcium aluminosilicate.
24.3.2 Compressive seals The main developments of compressive seals were achieved by PNNL, and have shown potential applicability for sealing SOFCs [45–47]. As opposed to the rigid seals, the compressive seals have lower requirements regarding strength, adhesion and CTE since no bonding is present at the interface. Thus this joining process employs materials to serve as a gasket between each cell. This joining approach requires the application of a compressive load to maintain hermetic sealing. Since no interface is created, the problem of matching CTEs of components is avoided as the seal can slide on the material of each individual cell. Among the advantages of this technique is the possibility of using other materials than ferritic stainless steels for interconnects; compressive seals are easy to apply and offer the possibility of repair of the stack by releasing the pressure. However, in order to employ compressive seals in a SOFC stack, a load frame is required to maintain the desired level of compression over the entire period of operation, and the stack components must be capable of withstanding that load. The load frame introduces several complexities in stack design, including: oxidation of the frame material, load relaxation due to creep, and increased weight and thermal mass (and therefore reduced specific power and thermal response of the
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overall system). In addition to increased system cost, these factors seriously limit the use of compressive seals in mobile applications. Initially, oxidation resistant metallic gaskets, such as gold or silver were investigated [48–50]. One significant problem of metallic gaskets is the solubility of oxygen, diffusing from the air side, and hydrogen, diffusing from the fuel side. Reaction between oxygen and hydrogen forms water within the seal, and leads to internal deformation, leakage and failure of the seal. Thus, before using metallic gaskets, a full understanding of the methods to control gas adsorption in metals will be required. Mica-based compressive seal are significantly more oxidation resistant when compared to metallic gaskets. Figure 24.11(a) presents a typical compressive mica seal. The major leaks observed in these seals occur at the
Metal
Mica
Ceramic
Metal
Mica
Compliant layer (Ag or glass)
Ceramic
Metal Mica Infiltrate
Compliant layer (Ag or glass)
Ceramic
24.11 Various configurations of mica-based seals: (a) plain seal, (b) with compliant layer and (c) infiltrated mica and compliant layer (reproduced with permission of Elsevier) [18].
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interfaces with the metal and the ceramic. Thus, the application of a compliant layer, such as shown in Fig. 24.11(b), helps reducing leakage. Improvements in leakage resistance were obtained for an infiltrated mica-based seal combined with the application of a compliant layer, as presented in Fig. 24.11(c). This type of seal is somewhat similar to the glass-ceramic seals. Strong bonds with all components were obtained, but reaction between the mica seal and the interconnect in hydrogen reduces the cell efficiency over time. Such compressive seals have shown success in manufacturing stacks, but requires the application of an applied compressive load during operation to maintain a gas-tight seal.
24.3.3 Compliant seals A compliant seal is a flexible interlayer that is sufficiently ductile to absorb deformation during heating or cooling of the assembly. A number of such flexible interlayer designs are possible, such as corrugated washers, honeycombs and dimples. The key properties for the selection of a material for a compliant seal suitable for SOFCs are high oxidation resistance, low stiffness, high ductility and low cost [14]. Weil et al. [14] selected a ferritic stainless steel as a compliant layer and manufactured a stack assembly between a Haynes 214 interconnect and a YSZ electrolyte. A cross-section and a schematic of the assembly are presented in Fig. 24.12. The seal thickness is 1.1 mm. A BNi-2 brazing material was used to join the compliant layer to the interconnect, while a Ag-4mol%CuO brazing alloy was used to join the electrolyte to the compliant material. The seals were hermetic and showed constant rupture strength after 20 thermal cycles. In all cases, the fracture occurred through the YSZ electrolyte, indicating that the compliant design is not the weakest part of the cell assembly.
24.4
Summary
This chapter presented the main joining routes envisaged for sealing SOFCs. As discussed, sealing SOFCs is a complex problem caused by the competitive interactions between both fuel and air service environments, the high temperatures of operation, and the multitude of materials involved in a SOFC stack. The research involved in solving the sealing problem is increasing every year and if current trends continue, a viable solution should be available in the not too distant future. Developments are also targeting lower-temperature SOFC operation (600 °C) in order to decrease the materials cost and ease the fabricating of stacks, which will enable the use of metallic materials with better mechanical properties and thermal conductivity.
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BNi-2 braze
Durafoil
Haynes 214
* Mid-section of the specimen not shown.
737
Ag-4mol% Ceramic PEN, CuO braze YSZ side down
(a) Ceramic bilayer PEN, YSZ side down
Air
Metal interconnect
Compliant foil H2 flow through gas manifold (b)
24.12 (a) Cross-section of a compliant seal assembly and (b) schematic representation of proposed compliant seal design to be used in SOFC manifold (reproduced with permission of Elsevier) [14].
24.5
References
1. N.Q. Minh, Ceramic Fuel Cells, Journal American Ceramic Society, 76, 3, 1993, p. 563–588. 2. D. Simwonis, A. Naoumidis, F.J. Dias, J. Linke and A. Moropoulou, Material characterization in support of the development of an anode substrate for solid oxide fuel cells, Journal of Materials Research, 12, 6, 1997, p. 1508–1518. 3. O. Yamamoto, Solid oxide fuel cells: fundamental aspects and prospects, Electrochemica Acta, 45, 2000, p. 2423–2435. 4. K. Huang and J. B. Goodenough, A solid oxide fuel cell based on Sr- and Mg-doped LaGaO3 electrolyte: the role of a rare-earth oxide buffer, Journal of Alloys and Compounds, 303–304, 2000, p. 454–464. 5. X. Xin, Z. Lü, Z. Ding, X. Huang, Z. Liu, X. Sha, Y. Zhang, W. Su, Synthesis and characteristics of nanocrystalline YSZ by homogeneous precipitation and its electrical properties, Journal of Alloys and Compounds, 425, 1–2, 30 November 2006, p. 69– 75. 6. V.V. Kharton, F.M.B. Marques, A. Atkinson, Transport properties of solid oxide electrolyte ceramics: a brief review, Solid State Ionics 174 (2004) p. 135–149.
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7. J.W. Fergus, Electrolytes for solid oxide fuel cells, Journal of power sources, 162, 2006, p. 30–40. 8. A. McEvoy, in: S.C. Singhal, K. Kendall (eds), High Temperature Solid Oxide Fuel Cells, Elsevier Ltd, Oxford, UK, 2003, p. 149. 9. S. Tao, J.T.S. Irvine; Discovery and characterization of novel oxide anodes for solid oxide fuel cells; The Chemical Record; 4, p. 83–95. 10. J.T.S. Irvine, A. Sauvet, Improved Oxidation of Hydrocarbons with New Electrodes in High Temperature Fuel Cells, Fuel Cells, 1, 3–4, 2001, p. 205. 11. J.W. Fergus, Oxide anode materials for solid oxide fuel cells, Solid State Ionics, 177, 2006, p. 1529–1541. 12. J.W. Fergus, Lanthanum chromite-based materials for solid oxide fuel cell interconnects, Solid State Ionics, 171, 1–2, 2004, p. 1–15. 13. J.W. Fergus, Metallic interconnects for solid oxide fuel cells, Materials Science and Engineering A 397, 2005, p. 271–283. 14. K.S. Weil, C.A. Coyle, J.S. Hardy, J.Y. Kim, G.G. Xia, Alternative planar SOFC sealing concepts, Fuel Cells Bulletin, May 2004, p. 11–16. 15. R. Barfod, S. Koch, Y.-L. Liu, P.H. Larsen, P.V. Hendriksen, Long-term test of dksofc cells, in: S.C. Singhal, M. Dokiya (eds), Electrochemical Society, SOFC VIII PV 2003–07, 1158–1166. 16. Jeff Stevenson, SOFC Seals: Materials Status, Proceedings of the SECA Core Technology Program – SOFC Seal Meeting, July 8, 2003. 17. P. Geasee, T. Schwickert, U. Diekmann, R. Conradt, in: J.G. Heinrich, F. Aldinger (eds), Ceramic Materials and Components for Engines, Wiley-VCH Verlag GmbH, Weinheim, Germany, 2001, p. 57–62. 18. J.W. Fergus, Sealants for solid Oxide fuel cells, Journal of power sources, 147, 2005, p. 46–57. 19. C. Lara, M.J. Pascual, M.O. Prado, A. Durán, Sintering of glasses in the system RO– Al2O3–BaO–SiO2 (R=Ca, Mg, Zn) studied by hot-stage microscopy, Solid State Ionics, 170, 3–4, 2004, p. 201–208. 20. N. Lahl, L. Singheiser, K. Hilpert, K. Singh, D. Bahadur, Aluminosilicate glassceramics as sealant in sofc stacks, Proceeding Electrochemical Society 1999–19 (SOFC VI) p. 1057–1066. 21. N. Lahl, K. Singh, L. Singheiser, K. Hilpert, D. Bahadur, Crystallization kinetics in AO-Al2O3-SiO2-B2O3 glasses (A = Ba, Ca, Mg), Journal Material Science, 35, 12, 2000, p. 3089–3096. 22. S.-B. Sohn, S.-Y. Choi, G.-H. Kim, H.-S. Song, G.-D. Kim, Stable sealing glass for planar solid oxide fuel cell, Journal Non-Crystalline Solids, 297, 2–3, 2002, p. 103– 112. 23. S.-B. Sohn, S.-Y. Choi, G.-H. Kim, H.-S. Song, G.-D. Kim, Suitable Glass-Ceramic Sealant for Planar Solid-Oxide Fuel Cells, Journal American Ceramic Society, 87, 2, 2004, p. 254–260. 24. C. Lara, M.J. Pascual, A. Durán, Glass-forming ability, sinterability and thermal properties in the systems RO–BaO–SiO2 (R = Mg, Zn), Journal Non-Crystalline Solids, 348, 2004, p. 149–155. 25. K.D. Meinhardt, J.D. Vienna, T.R. Armstrong and L.R. Pederson, Glass-ceramic joint and method of joining, U.S. Patent # 6,532,769. 26. N.P. Bansal, E.A. Gamble, Crystallization kinetics of a solid oxide fuel cell seal glass by differential thermal analysis, Journal of power sources, 147, 2005, p. 107– 115.
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27. N.P. Bansal, M.J. Hyatt, C.H. Drummond III, Crystallization and properties of Sr-Ba aluminosilicate glass-ceramic matrices, Ceramic Engineering Science Proceedings, 12, 7–8, 1991, p. 1222–1234. 28. K.S. Weil, J.E. Deibler, J.S. Hardy, D.S. Kim, G.-G. Xia, L.A. Chick, C.A. Coyle, Fuel Cells: Materials, Processing and Manufacturing Technologies – Rupture Testing as a Tool for Developing Planar Solid Oxide Fuel Cell Seals, Journal of Materials Engineering and Performance, 13, 3, 2003, p. 316–326. 29. W. Höland, G. Beall, Glass–Ceramics Technology, The American Ceramic Society, Westerville, OH, 2002. 30. D. Bahadur, N. Lahl, K. Singh, L. Singheiser, K. Hilpert, Influence of Nucleating Agents on the Chemical Interaction of MgO-Al2O3-SiO2-B2O3 Glass Sealants with Components of SOFCs, Journal Electrochemical Society, 151, 4, 2004, p. 558–562. 31. Z. Yang, J.W. Stevenson, K.D. Meinhardt, Chemical interactions of barium–calcium– aluminosilicate-based sealing glasses with oxidation resistant alloys, Solid State Ionics, 160, 3–4, 2003, p. 213–225. 32. Z. Yang, K.D. Meinhardt, J.W. Stevenson, Chemical Compatibility of Barium-CalciumAluminosilicate-Based Sealing Glasses with the Ferritic Stainless Steel Interconnect in SOFCs, Journal Electrochemical Society, 150, 8, 2003, p. 1095–1101. 33. K. Eichler, G. Solow, P. Otschik, W. Schaffrath, Degradation Effects at Sealing Glasses for the SOFC, pp. 899-906 in Proceedings of the Fourth European Solid Oxide Fuel Cell Forum, Oberrohrdorf, Switzerland, edited by A. J. McEvoy. European Fuel Cell Forum, Oberrohrdorf, Switzerland, 2000. 34. T. Schwickert, U. Reisgen, P. Geasee, R. Conradt, Electrically Insulating HighTemperature Joints for Ferritic Chromium Steel, Journal of Advanced Materials, 35, 4, 2003, p. 44–47. 35. R.K. Brow, Thermochemically stable sealing materials for sealing solid oxide fuel cells, Proceedings of the 6th Annual SECA workshop. 36. S.T. Reis, R.K. Brow, Designing Sealing Glasses for Solid Oxide Fuel Cells, Journal of Materials Engineering and Performance, 15, 4, 2006, p. 410–413. 37. H.J.L. Trapp, J.M. Stevels, Conventional and Invert Glasses containing titania, Part 1, Phy Chem Glasses, 1, 1960, p. 107–118. 38. K.G. Ewsuk, L.W. Harrison, Densification of Glass-Filled, Alumina Composites, in Sintering of Advanced Ceramics in Ceramic Transactions, Vol. 7, eds., C.A. Handwerker, J.E. Blendell, and W. Kaysser, The American Ceramic Society, Westerville, OH, 1990, p. 436–451. 39. M. Brochu, B.D. Gauntt, R. Shah, G. Miyake, R.E. Loehman, Comparison between barium and strontium-glass composites for sealing SOFCs, Journal of the European Ceramic Society, 26, 15, 2006, p. 3307–3313. 40. M. Brochu, B.D. Gauntt, R. Shah, R.E. Loehman, Comparison between micrometerand nano-scale glass composites for sealing solid oxide fuel cells, Journal of the American Ceramic Society, 89, 3, 2006, p. 810–816. 41. K.A. Nielson, M. Solvang, S.B.L. Nielson, A.R. Dinesen, D. Beeaff, P.H. Larsen, Glass composite seals for SOFC application, Journal of the European Ceramic Society, 27, 2007, p. 1817–1822. 42. R.E. Loehman, M. Brochu, B.D. Gauntt, K. Malone, Composite Seals for Solid Oxide Fuel Cells, 30th International Conference on Advanced Ceramics and Composites, Cocoa Beach, USA, January 22nd–27th, 2006. 43. C.C. Beatty, Compliant glass–silver seals for SOFC application. in Proceedings of the IX ECS on Solid Oxide Fuel Cells, ed. S.C. Singhal, J. Mizusaki, 2005, p. 1949.
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44. K.S. Weil, J.S. Hardy, J.P. Rice, J.Y. Kim, Brazing as a means of sealing ceramic membranes for use in advanced coal gasification processes, Fuel, 85, 2006, p. 156– 162. 45. Y.-S. Chou, J. W. Stevenson, Phlogopite mica-based compressive seals for solid oxide fuel cells: effect of mica thickness, Journal of Power Sources, 124, 2, 2003, p. 473–478. 46. Y.-S. Chou, J. W. Stevenson, Novel infiltrated Phlogopite mica compressive seals for solid oxide fuel cells, Journal of Power Sources, 135, 1–2, 2004, p. 72–78. 47. Y.-S. Chou, J. W. Stevenson, P. Singh, Thermal cycle stability of a novel glass–mica composite seal for solid oxide fuel cells: Effect of glass volume fraction and stresses, Journal of Power Sources, 152, 2005, p. 168–174. 48. M. Bram, S. Reckers, P. Drinovac, J. Moench, R.W. Steinbrech, H.P. Buchkremer, D. Stoever, Characterization and evaluation of compression loaded sealing concepts for SOFC stacks, Proc. Electrochem. Soc. 2003–07 (SOFC VIII), 2003, p. 888–897. 49. M. Bram, S.E. Bruenings, F. Meschke, W.A. Meulenberg, H.P. Buchkremer, R.W. Steinbrech, D. Stoever, Application of Metallic Gaskets in SOFC-Stacks, Proc. Electrochem. Soc. 2001–16 (SOFC VII), 2001, p. 875–884. 50. J. Duquette, A. Petric, Silver wire seal design for planar solid oxide fuel cell stack, J. Power Sources, 137, 1, 2004, p. 71–75.
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25 Joining of bulk nanostructured materials M B R O C H U , McGill University, Canada
25.1
Introduction
The intensive research in the field of nanotechnology in recent years has led to the emergence of several novel materials that possess improved properties when compared to conventional materials. Of these, bulk nanomaterials are most interesting, from a practical point of view, as they can expand the use of current materials with their superior properties [1]. Bulk nanomaterials are single-phase or multi-phase polycrystal materials possessing a grain size distribution ranging between 1 and 100 nm in at least one dimension [2]. The equal or higher proportion of atoms at the grain boundaries, as compared to the number of atoms within the grains, is responsible for the novel microstructures, which lead to the extraordinary potential of nanocrystalline materials. Assuming the grains are spherical, the volume fraction of interfaces (grain boundaries) can be as much as 50% for 5 nm grains, 30% for 10 nm grains and 3% for 100 nm grains [3]. In particular, bulk nanomaterials possess enhanced diffusivity, superior strength and improved formability, which also include ductilization of ceramics and intermetallics, improved electrical, optical and magnetic properties, high thermal expansion and improved corrosion behaviour [1]. Despite the current low number of commercially viable applications of large-scale nanomaterials, the wide spectrum over which they possess improved properties allows for the potential of revolutionizing several industrial fields. Applications such as structural materials (yield in specific strength), electronics, energy (fuel cells), wear and corrosion applications (cermet for cutting tools, nanostructured coatings) are a few of the many areas targeted by the current research. Research performed over the last fifteen years resulted in the development of several processes to fabricate nano-powders from ceramics or metals. The next step in the commercialization of bulk nanomaterials is the creation of a cost-effective fabrication method followed by the development of joining processes that maintain the nanostructure of the part or assembly to be joined. Typically, joining processes are left to the end of the manufacturing 741 WPNL2204
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flow chart and therefore receive lower attention and funding; however, this latter step can jeopardize quality and reliability of the component. Despite the fact that the science of joining nanomaterials is in its infancy, the present chapter is aimed at highlighting the main direction of research taken by the community. Prior to presenting the discussion on current trends and the small number of advancements in the field, a brief description of the commonly used manufacturing processes and the relation between the processing conditions and the thermal stability of the materials will be discussed. The knowledge and understanding of this relation is a crucial factor in the selection of a joining process and the appropriate joining parameters.
25.2
Fabrication processes
As in any industrial field, joining processes are only necessary once material processing reaches a stage where availability and reliability of the material is well established. To attain this type of material, rigorous and reliable manufacturing techniques are required. The main challenge in fabricating large-scale nanomaterials, as it is for joining, is to ensure that the final part maintains a nanostructure. This implies the use of process parameters that will mitigate recrystallization or grain growth. The main routes for manufacturing large-scale nanomaterials, i.e., identified as possessing a high potential for commercialization, are the powder metallurgy (PM), the electrodeposition and the severe plastic deformation routes.
25.2.1 The powder metallurgy (PM) approach The PM route allows the fabrication of large-scale nanomaterials from powder precursors. In the present case, two types of nano-powders are available: nano-size crystals and nanostructured powders. The former consists of a single crystal possessing a size distribution ranging between 5 and 100 nm. This type of powder is generally fabricated by inert gas condensation [4, 5] and plasma processing techniques [6, 7]. The use of nano-crystals is not trivial since metallic powders are highly reactive with the environment and difficult to handle and these powders can be difficult to consolidate due to the oxides that form at the surfaces of the crystals. The nanostructured powder is micron-sized and possesses a nano-grained structure. They are generally obtained by milling micron grained powders, where a repetition of welding, fracturing and re-welding events take place [8]. The welding and fracturing events create subgrains which then lead to a reduction in overall grain size. Milling operations can be performed under dry (controlled atmosphere) or wet (liquid nitrogen, ethanol, etc.) conditions [9]. It has been shown that nanometre-sized grains can be obtained in almost any metal-base material after sufficient milling time [3, 8–9]. The grain size
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decreases with milling time down to a minimum value that appears to be inversely proportional to the melting temperature. Mechanical milling is a process allowing medium volume production and the resulting nanostructured powders are feedstock for bulk nanomaterials and thick nanostructured coatings. The adapted-conventional powder metallurgy processes used in some approaches are pressure-assisted processes such as hot pressing, hot isostatic pressing and hot extrusion. The pressure-assisted processes are generally preferred since pore closure is enhanced by the external forces resulting in densification at lower temperatures and lower microstructure gradients. Numerous studies in the literature reported the fabrication of bulk nanomaterials using conventional processes [8]. The shockwave consolidation technique is another viable process, which imposes a travellng planar shockwave front into a contained-porous body. Under the presence of pressures above 1 GPa applied for short durations in the micro-second range, bonding between particles occurs [10]. The extreme processing conditions prevent heat transfer and thus inhibit recrystallization and grain growth. Despite the fact that few systems have been studied, it has been proven that this process maintains nanostructure after consolidation [10–13].
25.2.2 Electrodeposition Electrodeposition is a simple, inexpensive, and versatile method to produce dense metallic nanocrystals [14]. Figure 25.1 presents a schematic of the pulsed process. During the electrodeposition process, the applied current is pulsed and as the current spikes, the metal cations present in the solution are deposited in crystalline or amorphous patches on the substrate. The recent developments in electrodeposition allow for the fabrication of a wide variety
V
V
V
Time
Voltage
Voltage
Voltage
nm
Time
Time
25.1 Schematic representation of pulsed electrodeposition process.
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of metals, binary and ternary alloys as well as nanostructured composites [15]. It has been shown that electrodeposition yields grain sizes in the nanometre range when the electrodeposition variables are chosen such that nucleation of new grains is favoured over the growth of existing grains [15]. This was achieved by using high deposition rates, formation of appropriate complexes in the bath, addition of suitable surface-active elements to reduce surface diffusion of ad-atoms, etc. This technique can yield porosity-free finished products that do not require subsequent consolidation processing. This process requires low capital investment, provides high production rates, has few shape and size limitations and possesses a high probability of technology transfer to existing electroplating and electroforming industries.
25.2.3 Severe plastic deformation The severe plastic deformation (SPD) technique is characterized by introducing large amounts of plastic strain, over one thousand percent for some materials, at low homologous temperatures (typically below 0.3 of the melting temperature) to a bulk micron-grained material. The severe plastic deformation leads to a subdivision of the initially coarse-grained microstructure into a hierarchical system of cell blocks and dislocation cells. With increasing straining of the material, the size of the microstructural elements decreases to the nanometric scale. At the same time the disorientation – difference in crystallographic orientation – is increased. In order to obtain the smallest microstructural sizes, plastic strains of more than 600 to 800% are necessary depending on the material. The special feature of all variants of SPD is that the cross-section of the material remains constant during or after SPD processing. Thus high degrees of plastic deformation are possible as one sample can be subjected to the SPD process several times in order to accumulate the total amount of plastic strain required for a nanometric grain size. SPD is a proven technique to fabricate nanostructured materials [16] and it encompasses a family of processing methods that can impose extremely large strains (> 4) on materials with the intent of grain refinement. Processes such as equal-channel angular pressing, cyclic extrusion-compression, high-pressure torsion, accumulative roll-bonding are some examples of processes imposing large deformation and a schematic representation of SPD processes is presented in Fig. 25.2 [17]. Equal-channel angular pressing has an important advantage over processing by standard powder-metallurgy procedures because it may be applied directly to bulk samples prepared by ingot metallurgy. Nanomaterials fabricated by SPD possessing increase in tensile strength from between 25 to 100% have been reported [18].
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CEC
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HPT
Plunger
Anvil
Die Φ
dm
Ψ Sample Sample
Extrusion compression
Support
25.2 Schematic representation of equal-channel angular pressing, cyclic extrusion-compression and high-pressure torsion processes.
25.3
Non-equilibrium state of microstructure
As mentioned earlier, the main challenge in processing or joining bulk nanostructured materials is to prevent any grain growth or recrystallization. Failing to do so transforms the part into an ultra-fine grained material (grain size ranging between 100 and 300 nm) or ultimately reverts to conventional materials, i.e., when the grain size will be above 500 nm. Thus, the key point in consolidating or joining bulk nanomaterials is to perform the operation within a time-temperature window that will mitigate changes in the microstructure. Several studies on thermal stability of various nanomaterials are available in the literature and only few were selected to show pertinent examples of the extensive information. Figures 25.3(a), (b) and (c) present the evolution of the grain size of iron, nickel and aluminum nanostructured systems, respectively, as a function of the homologous temperature. The homologous temperature is defined as the ratio of the tested temperature divided by the melting temperature of the metal/alloy. From these curves, key concepts regarding the interaction between the maximum joining temperature and the nanostructured system can be extracted. The thermal stability of the material changes dramatically depending on whether the consolidated material is in pure or alloy form. Both grain boundary solute and precipitate drag provide grain boundary pinning force which results in the enhancement of the thermal stability [8, 27]. The former effect is clearly observed in Figure 25.3(b), while electrodeposited Ni-1.2%P alloy has a better thermal stability than the corresponding pure Ni metal. The latter interaction is observed in Fig. 25.3(a), where during cryomilling, the presence of aluminium as solute (Fig. 25.3(a)) or solvent (Fig 25.3(c)) in the alloys was shown to enhance thermal stability. This stability is due to nanoscale
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350 Fe ball milled [19] Fe cryomilled [20] Fe milled [21] Fe10Al cryomilled [20]
Grain size (nm)
300 250 200 150 100 50 0 0.1
0.2
0.3 0.4 0.5 0.6 Homologous temperature (T / Tm)
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(a) 400 Ni cryomilled [22] Ni electrodeposited [23] Ni-P electrodeposited [24]
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(b) 250 5083 severe plastic deformation [25] 5083 cryomilling [26]
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(c)
25.3 Thermal stability of nanocrystalline (a) Fe, (b) Ni and (c) Al-5083 as a function of fabrication process.
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AlN precipitates formed during the milling process that pins the grain boundaries. For the case of pure metals having a lower tendency to form strengthening precipitates, their behaviour is a function of the fabrication process. Figure 25.3(b) and 25.3(c) demonstrate that the thermal stability of cryomilled and HIPed Ni is higher than for electrodeposited Ni1. Similarly, the thermal stability of cryomilled and HIPed 5083 alloy is significantly higher than samples processed by severe plastic deformation. The intent of this section was to demonstrate that joining bulk nanomaterials has several peculiarities. In addition to the joining parameters (time-temperature relation,…), the parent material chemistry and the fabrication process will have a significant influence on the joining procedure. In addition, the low homologous temperature at which grain growth becomes significant inhibits the use of almost all conventional welding processes.
25.4
Attempts in joining bulk nanomaterials
Through the natural sequence of material development, the need for reliable joining techniques is increased when the newly designed materials are readily available and start to be used commercially. In that respect, since the knowledge and understanding of the processing-property relationship of bulk nanomaterials increases, attempts to use them in structural applications are envisaged. At this point in time, research efforts focused on joining these nanomaterials is emerging in the scientific community. The following sections present some of the pioneering processes, as well as results obtained to date in this field. A person interested in the field should remember that the process used to manufacture bulk materials and its relation to the thermal stability of the material is a criterion for the selection of the joining process.
25.4.1 Spark plasma sintering The spark plasma sintering process (SPS) is based on the diffusion bonding processes using interlayers and would be applicable to nanomaterial systems possessing relatively good thermal stability. This process is a recent adaptation of the commonly used hot pressing technique. The two processes are similar; however, where they differ is in the method of incorporating heat during processing. In the SPS process, instead of placing the powder in a heated apparatus, it is placed within a graphite die and a pulsing electrical current is passed through the die and the powder bed all while being under the influence of an externally applied load. A schematic diagram of the apparatus
1
For the electrodeposited alloy, two curves are presented since the material has a bimodal grain size distribution at temperature below 0.3 homologous temperature.
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is presented in Fig. 25.4. The main advantages of this fabrication method are the lower processing temperature and shorter duration cycles, both of which are advantageous for the consolidation of metastable materials. Many studies reported the fabrication of bulk nano-metals or nano-intermetallics using this technique [28, 29] and it is then possible to envisage using SPS as a joining process. Thus, SPS has been used to fabricate and simultaneously join nanocrystalline materials. Liu and Naka [30] reported the simultaneous fabrication and joining of Ni3Al-TiC/Ni3Al nanocomposites at temperatures ranging between 900 and 1100 °C, which corresponds to a homologous temperature ranging between ≈0.68 and ≈0.78, for soaking time of 5 minutes. Figure 25.5 presents a SEM micrograph of the interface. Their results showed formation of fully densified samples with an average grain size after consolidation of 60.2±17.6 nm Upper punch/electrode
P
Controlled atmosphere chamber
Powder
Punch and die
Lower punch/electrode
DC pulse generator
P
25.4 Schematic of a Spark Plasma Sintering Apparatus.
TiC/Ni3Al
Ni3 Al
25.5 SEM micrograph of a SPS of the interface between TiC/Ni3AlNi3Al (reproduced with permission of Elsevier) [30].
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(measured by XRD broadening) for the sample bonded at 1100 °C. The samples had shear strengths of 765±37.1 MPa. Interestingly, fracture of the samples occurred mainly in the composite side and partly through the interface demonstrating the potential of this approach to fabricate reliable joints. The viability of SPS for fabricating and simultaneously consolidating and joining bulk nanomaterials can easily allow us to speculate that joining preexisting bulk nanomaterials will be possible using a nanostructured powder bed as an interlayer; however, this work has yet to be performed. Owing to the nature of this joining technique, it is expected that hot pressing should yield similar results but with a coarser microstructure since the heating schedule is more demanding. Similarly to advanced joining processes applied to conventional materials, no apparent limitation in combination of parent materials to be joined is expected if the powder interlayer is carefully selected. Unfortunately, the approach will have limitations, since the sample geometry must be fairly simple and the parent materials must be of similar dimension.
25.4.2 Diffusion bonding using a bed of nanoparticles as interlayer The notion that small particles (typically below 10 nm) possess very different physical properties, such as lower melting point than the micron-scale counterparts, has been studied for nearly 50 years [31, 32]. The significant surface free energy of small particles creates a relationship where the melting point decreases linearly with the inverse of the particle radius [32]. To reduce the total free energy of the system, very fine particles have a tendency to sinter or Ostwald ripen. This phenomenon is called the Gibbs-Thomson effect and occurs when fine nanoparticles undergo spontaneous coalescence under standard conditions (pressure and temperature) until the driving force becomes insufficient for further growth of the nanoparticles. Figure 25.6 shows this coalescence phenomenon applied to gold nano-powders.
50 nm
50 nm
Before reheating
50 nm
Tf = 200 °C
Tf = 400 °C
25.6 TEM images of gold nano-particles at different reheating temperatures (reproduced with permission of Elsevier) [40].
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This equilibrium between the surface and the volume energies determines the lower diameter limit to which a nanoparticle can survive. Interestingly, the exposure of an agglomerate of fine nanoparticles to a controlled heat source or imposing high external pressure level to the powder activates a self-sintering mechanism. Several studies have shown visual observation of the coalescence of metallic nanoparticles under various sources of heat [33–35]. Recently, numerous studies on fabricating bulk nanostructured materials through sintering of nanoparticles have demonstrated the possibility of engineering fabrication processes to take advantage of this phenomenon [36–39]. In light of the physical evidence of self-sintering at low to moderate temperature, the high surface energy can be utilized to join nanomaterials by sintering an interlayer composed of nano-crystals between the two parent materials and this, at a significantly lower bonding temperature than that in an ordinary process such as diffusion bonding. Two examples showing the possibility of using the surface energy contribution of nano-particles for bonding are currently available in the literature. Ide et al. [41] reported that careful decomposition of silver-based organometallics leads to a well-dispersed slurry containing metallic silver nanoparticles possessing an average diameter of 10 nm. This slurry can then be used as filler material and experiments using the slurry have demonstrated that bonding within the interlayer and between the parent materials can be obtained at temperatures as low as 300 °C (0.31 homologous temperature). To demonstrate the concept that only fine nanoparticles are effective interlayers, similar experiments were performed using coarser particles (100 nm). Interfaces, obtained for pressure levels of 1 and 5 MPa for both nanoparticles and coarse particles, respectively, are compared in Fig. 25.7. The filler metal slurry containing the largest particles has a significant volume fraction of pores, which is not the case for the slurry composed of fine particles. This lack of bonding is reflected in the shear strength results; the joint is 300 times stronger with the fine particle interlayer than for the coarser one. The cleanliness of the surface (free of contaminants) plays a crucial role in this approach, since the surface energy of all metallic nanoparticles is lowered by the presence of contaminants. In the former case, silver, being relatively noble compared to other metals, would be less affected by the presence of contaminants and would be more suitable. Nano-metallic particles from any noble metal would be expected to behave similarly. A good example of the lowered efficiency of surface contaminated nanoparticles is reported in the work of Onaka and Funkenbusch [42], where Cu nanoparticles were used as interlayer during diffusion bonding. Their results show bonding at temperatures as low as 400 oC (0.37 homologous temperature). Despite the fact that care was taken during sample preparation and bonding, the interlayer oxidized during the heat treatment cycle, leaving a porous copper/copper oxide composite interlayer. The presence of surface oxide
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10 nm Ag particles Sintered Ag layer
Sintered Ag layer Void Void
Cu
Cu 1 MPa
5 MPa
5 µm
100 nm Ag particles Sintered Ag
Void Sintered Ag layer Gap Void
Void Cu
Cu 1 MPa
5 MPa
5 µm
25.7 Interfacial comparison between 10 nm Ag powders and 100 nm Ag powders as interlayer during low pressure-low temperature hot pressing (reproduced with permission of Elsevier) [41].
significantly reduced the quality of the interlayer, which correlates to the low mechanical properties of only 40% strength of the parent material. Youngdahl et al. reported the same detrimental effect of surface oxides during trials for the fabrication of bulk nanostructured materials from nanopowders [43]. The two examples presented above demonstrated bonding at homologous temperatures of 0.31 and 0.37, respectively. Comparing these temperatures with the thermal stability results of nanomaterials presented in Fig. 25.3, this approach seems viable since the reported bonding temperature is lower than the critical grain growth temperature. Thus, engineering joining processes for bulk nanostructured metals using the high surface energy of a nanometallic particle is viable, but the known advantage of the high surface energy of the particle is at the same time its worst enemy since metallic nanoparticles react vigorously with contaminants. This will limit the utilization of most metallic nanoparticle systems since the removal of the surface oxide scale is difficult, even almost impossible at such small diameters.
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As was seen in the case of metallic nanoparticles, ceramic nanoparticles possess accelerated sintering behaviour resulting from enhanced diffusivity. Hahn et al. have demonstrated that TiO2 nanoparticles of 14 nm in size could be fully densified at temperatures as low as 900 °C [44]. Thus, diffusion bonding nano-ceramic materials using an interlayer of nano-ceramic particles is an attractive route for joining these metastable ceramics. Cina and Eldror [45] have demonstrated the feasibility of bonding stabilized zirconia (Y-TZP) using Y-TZP nanoparticles of 42 nm at a temperature as low as 1090 °C. This represents one of the numerous examples presented in the literature, which are all based on the high surface energy of fine particles. As opposed to the metallic nanopowders, the powder particles used for this approach are already oxides and therefore there is an absence of oxidation contamination. This means that nano-ceramic systems could be joined by this technique
25.4.3 Welding processes Despite the high heat input of conventional welding processes, few attempts in welding bulk nanostructured 5083 alloy have been reported. Three welding processes, namely tungsten arc welding, inertia welding and friction stir welding were used to joined plates of the bulk n-5083 alloy fabricated from HIPed and hot extruded cryomilled powder [46]. The thermal stability of this alloy is presented in Fig. 25.3(c). Unfortunately, only hardness results are reported and no significant detail on the welding procedure was described. The hardness results of the fusion zone for both the TIG and inertia welding remained similar to the parent materials after joining. On the other hand, the friction stir welding process had a reduction in hardness in the fusion zone of nearly 20%. Some care must be taken in interpreting these results, since the hardness in nanomaterials or UFG depends on any substructure formed during the process, i.e., precipitation of Mg from the supersaturated solid solution. In addition the materials processed from cryomilled powders possess a large fraction of nanometric precipitates of AlN which significantly aid in maintaining the fine grain structure. This conservation of the hardness in the fusion zone might not be observed in materials fabricated by severe plastic deformation or electrodeposition since their grain stability is lower.
25.5
Metallic glasses
Amorphous metal can be obtained by imposing a sufficiently rapid cooling rate on a pure metal or an alloy to prevent any ordering (crystallization) of the structure during solidification. They can also be produced by solid state amorphization or during electrodeposition, as previously described in Section
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25.2.2 [47]. Thus the crystallization of any amorphous metal under an optimized heat treatment schedule is another option to fabricate polycrystalline materials possessing ultra fine grain structure. The driving force for the crystallization is the difference in Gibbs free energy between the amorphous and crystalline state, and where the driving force for crystallization decreases when the size of the crystal increase. The amorphization of structural alloys, as discussed so far, is feasible through techniques such as melt spinning, vapour deposition or electrodeposition; but their composition will produce a strong driving force for crystallization and grain growth, which will limit their joining to processes having a combination of short time and low temperature to avoid any grain coarsening. On the other hand, a significant reduction of the driving force for crystallization was observed during the development of bulk metallic glasses, a family of materials having a wide supercooled liquid region before crystallization and thus a high glass-forming ability. These materials possess a complex chemistry and preserving their stoichiometry is crucial to maintaining their high glass-forming ability. It also limits the number of potential nanosystems and dictates the composition and volume fraction of the crystalline phases (thermodynamic phase equilibrium). By nature, the melting point of these alloys is significantly lower than their respective elements. Kawamura [48] presented a thorough review of the viability of welding bulk metallic glasses with conventional welding processes such as explosive welding, pulse-current welding, electron beam welding and friction welding. A complete description of these processes can be found in [49]. For these four welding processes, he reported joining parameters resulting in bond strength equivalent to that of the parent materials, without inducing crystallization of either the fusion zone or the heat affected zone. He also demonstrated the possibility of welding bulk metallic glasses to some conventional grain size. Figure 25.8 shows a schematic of the melt-cooling and amorphous solid-heating crystallization TTT diagrams for two bulk metallic glasses. The process temperature cycle for a few welding processes is also incorporated. As illustrated, any welding process having a process temperature cycle shorter than the time required for crystallization will be viable, and proves the applicability of conventional welding process for these materials. Bulk metallic glasses can be an attractive filler metal to join bulk nanostructured materials, but only if the process temperature cycle is faster than the grain growth or recrystallization kinetic of the bulk nanomaterials. This approach could be attractive for bulk nanomaterials having a high thermal stability with the selection of low melting temperature bulk metallic glass as filler metal. Such experiments are yet to be carried out but this approach should lead to interesting results.
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Microjoining and nanojoining Higher energy concentration Explosion
Pulse-current
Electron-beam, Laser-beam Higher glass forming ability ZrAINiCu
ZrBeTiNiCu
Melt Temperature, T
Tm SCL
Tg Solid
10–4
10–2
100
102
Time, t /s
25.8 Melt-cooling and amorphous solid-heating crystallisation TTT diagrams for two bulk metallic glasses (reproduced with permission of Elsevier) [48].
25.6
Summary
It can be seen that the development of joining processes aimed at joining bulk nanostructured materials is an emerging field and will gain attention in the near future. The complexity of joining these materials is due to the metastable nature of the nanostructure and, in all cases, the joining operating window must take into consideration the kinetic of grain growth (temperature and time), to prevent the loss of the desired nanostructure. Further major advances in this area will require a better knowledge of the metastable grain structure and ways to increase the thermal stability of the nanograins. Any improvements in either of these areas will allow for an increase in the joining parameter window and may also enable the use of other joining processes.
25.7
References
1. Suryanarayana, ‘Nanocrystalline Materials’, International Material Reviews, 40, 2, 1995, p. 41–64. 2. R.W. Sielgel, ‘Nanostructured Materials – Mind over Matter’, Nanostructured Materials, 4, 1, 1994, p. 121–138. 3. M.A. Meyers, A. Mishra and D.J. Benson, ‘Mechanical Properties of Nanocrystalline Materials’, Progress in Materials Science, 51, 2006, p. 427–556. 4. H. Gleiter, ‘Nanocrystalline Materials’, Progress in Materials Science, 33, 1989, p. 223–315.
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5. H.J. Höfler, R. Tao, L. Kim, R.S. Averback and C.J. Altstetter, ‘Mechanical Properties of Single-Phase and Nano-Composite Metals and Ceramics’, Nanostructured Materials, 6, 5–8, 1995, p. 901–904. 6. R. Kalyanaraman, Sang Yoo, M.S. Krupashankara, T.S. Sudarshan and R.J. Dowding, ‘Synthesis and Consolidation of Iron Nanopowders’, Nanostructured Materials, 10, 8, 1998, p. 1379–1392. 7. C. Qin and S. Coulombe, ‘Synthesis of Organic Layer-Coated Copper Nanoparticles in a Dual-Plasma Process’, Materials Letters, 60, 16, 2006, p. 1973–1976. 8. D.B. Witkin and E.J. Lavernia, ‘Synthesis and Mechanical Behavior of Nanostructured Materials via Cryomilling’, Progress in Materials Science, 51, 2006, p. 1–60. 9. Suryanarayana, ‘Mechanical Alloying and Milling’, Progress in Materials Science, 46, 1–2, 2001, p. 1–184. 10. T.G. Nieh, P. Luo, W. Nellis, D. Lesuer and D. Benson, ‘Dynamic Compaction of Al Nanocrystals’, Acta Materialia, 44, 8, 1996, p. 3781–3788. 11. G.E. Korth and R.L. Williamson, ‘Dynamic Consolidation of Metastable Nanocrystalline Powders’, Metallurgical and Materials Transactions A, 26, 1995, p. 2571–2578. 12. T. Chen, J.M. Hampikian and N.N. Thadhani, ‘Synthesis and Characterization of Mechanically Alloyed and Shock-consolidated Nanocrystalline NiAl Intermetallic’, Acta Materialia, 47, 8, 1999, p. 2567–2579. 13. C.P. Dogan, J.C. Rawers, R.D. Govier and G. Korth, ‘Mechanical Processing, Compaction and Thermal Processing of α-Fe Powder’, Nanostructured Materials, 4, 6, 1994, p. 631–644. 14. F. Ebrahimi, G.R. Bourne, M.S. Kelly and T.E. Matthews, ‘Mechanical Properties of Nanocrystalline Nickel Produced by Electrodeposition’, Nanostructured Materials, 11, 3, 1999, p. 343–350. 15. U. Erb, ‘Electrodeposited Nanocrystals: Synthesis, Properties and Industrial Applications’, Nanostructured Materials, 6, 1995, p. 533–538. 16. V.M. Segal, ‘Equal Channel Angular Extrusion: from Macromechanics to Structure Formation’, Materials Science and Engineering A, 271, 1–2 1999, p. 322–333. 17. R.Z. Valiev, Y. Estrin, Z. Horita, T.G. Langdon, M.J. Zehetbauer and Y.T. Zhu, ‘Producing Bulk Ultrafine-Grained Materials by Severe Plastic Deformation’, Journal of Metals, April 2006, p. 33–39. 18. T.C. Lowe, ‘Metals and Alloys, Nanostructured by Severe Plastic Deformation: Commercialization Pathways’, Journal of Metals, April 2006, p. 28–32. 19. F. Zhou, J. Lee, S. Dallek and E.J. Lavernia, ‘High Grain Size Stability of Nanocrystalline Al Prepared by Mechanical Attrition’, Journal of Materials Research, 16, 12, 2001, p. 3451–3458. 20. R.J. Perez, H.G. Jiang, C.P. Dogan and E.J. Lavernia, ‘Grain Growth of Nanocrystalline Cryomilled Fe-Al Powders’, Metallurgical and Materials Transactions A, 29, 1998 p. 2469–2475. 21. C.H. Moelle and H-J Fecht, ‘Thermal Stability of Nanocristalline Iron Prepared by Mechanical Attrition’, Nanostructured Materials, 6, 1–4, 1995, p. 421–424. 22. J. Lee, F. Zhou, K.H. Chung, N.J. Kim and E.J. Lavernia, ‘Grain Growth of Nanocrystalline Ni Powders Prepared by Cryomilling’, Metallurgical and Materials Transactions A, 32, 2001, p. 3109–3115. 23. U. Klement, U. Erb, M. El-Sherik and K.T. Aust, ‘Thermal Stability of Nanocrystalline Ni’, Materials Science and Engineering A, 203, 1–2, 1995, p. 177–186. 24. K. Boylan, D. Ostrander, U. Erb, G. Palumbo and K.T. Aust, ‘An In-Situ TEM Study
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27. 28.
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30. 31. 32. 33.
34. 35. 36. 37.
38. 39. 40.
41. 42.
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Microjoining and nanojoining of the Thermal Stability of Nanocrystalline Ni-P’, Scripta Metallurgica et Materialia, 25, 1991, p. 2711–2716. Z. Horita, T. Fujnami, M. Nemoto and T.G. Langdon, ‘Equal-Channel Angular Pressing of Commercial Aluminum Alloys: Grain Refinement, Thermal Stability and Tensile Properties’, Metallurgical and Materials Transactions A, 31, 2000, p. 691–701. F. Zhou, R. Rodriguez and E.J. Lavernia, ‘Thermally Stable Nanocrystalline Al-Mg Alloy Powders Produced by Cryomilling’, Materials Science Forum, 386–388, 2002, p. 409–414. D.L. Feldheim and C.A. Foss Jr, editors, Nanoparticles: building blocks for nanotechnology; New York: Kluwer Academic/Plenum Press, 2004. S.H. Lee, K.I. Moon, H. S. Hong and K. S. Lee, ‘Microstructural and Mechanical Properties of Nanocrystalline (Al+12.5 at.% Cu)3Zr Alloys Synthesized by Planetary Ball Milling and Plasma Spark Sintering’, Intermetallics, 11, 2003, p. 1039–1045. Y. Minamino, Y. Koizumi, N. Tsuji, N. Hirohata, K. Mizuuchi and Y. Ohkanda, ‘Microstructures and Mechanical Properties of Bulk Nanocrystalline Fe-Al-C Alloys made by Mechanically Alloying with subsequent Spark Plasma Sintering’, Science and Technology of Advanced Materials, 5, 1–2, 2004, p. 133–143. W. Liu and M. Naka, ‘In Situ Joining of Dissimilar Nanocrystalline Materials by Spark Plasma Sintering’, Scripta Materialia, 48, 2003, p. 1225–1230. S.C Tjong and H. Chen, ‘Nanocrystalline Materials and Coatings’, Materials Science and Engineering R, 45, 1–2, 2004, p. 1–88. G.L. Allen, R.A. Bayles, W.W. Gile and W.A. Jesser, ‘Small Particle Melting of Pure Metals’, Thin Solid Films, 144, 1986, p. 297–308. S. Arcidiacono, N. R. Bieri, D. Poulikakos and C. P. Grigoropoulos, ‘On the Coalescence of Gold Nanoparticles’, International Journal of Multiphase Flow, 30, 7–8, 2004, p. 979–994. B. Giesen, H. R. Orthner, A. Kowalik and P. Roth, Chemical Engineering Science, 59, 11, 2004, p. 2201–2211. Y. Tamou and S.I. Tanaka, ‘Formation and Coalescence of Tungsten Nanoparticles under Electron Beam Irradiation’, Nanostructured Materials, 12, 1999, p. 123–126. D.M. Owen and A.H. Chokshi, ‘An Evaluation of the Densification Characteristics of Nanocrystalline Materials’, Nanostructured, Materials, 2, 2, 1993, p. 181–187. M. Jain, and T. Christman, ‘Synthesis, Processing and Deformation of Bulk Nanophase Fe-28Al-2Cr Intermetallic’, Acta Metallurgica et Materialia, 42, 6, 1994 p. 1901– 1911. T.R. Smith and K.S. Vecchio, ‘Synthesis and Mechanical Properties of Nanoscale Mechanically-Milled NiAl’, Nanostructured Materials, 5, 1, 1995, p. 11–23. J.R. Groza and R.J. Dowding, ‘Nanoparticulate Materials Densification’, Nanostructured Materials, 7, 1996, p. 749–768. K. Nakaso, M. Shimada, K. Okuyama and K. Deppert, ‘Evaluation of the Change in the Morphology of Gold Nanoparticles during Sintering’, Journal of Aerosol Science, 33, 7, 2002, p. 1061–1074. E. Ide, S. Angata, A. Hirose and K.F. Kobayashi, ‘Metal-Metal Bonding Process using Ag Metallo-organic Nanoparticles’, Acta Materialia, 53, 2005, p. 2385–2393. S. Onaka and P.D. Funkenbusch, ‘Low Temperature Diffusion Bonding Using NanoScale Materials, Advances in powder metallurgy & particulate materials – 1992’, vol 8 Properties of emerging P/M materials, San Fransisco, Ca, USA 21–26 June, p. 33–44. J. Youngdahl, P.G. Sanders, J.A. Eastman and J.R. Weertman, ‘Compressive Yield
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44. 45. 46. 47. 48. 49.
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Strengths of Nanocrystalline Cu and Pd’, Scripta Materialia, 37, 6, 1997, p. 809– 813. H. Hahn, J. Logas, R.S. Averback, Journal of Materials Research, 5, 1990, p. 609– 614. B. Cina and I. Eldror, ‘Bonding of Stabilised Zirconia (Y-TZP) by Means of Nano Y-TZP Particles’, Materials Science and Engineering A, 301, 2001, p. 187–195. A. Piers Newbery, S.R. Nutt and E.J. Lavernia, ‘Multi-Scale Al 5083 for Military Vehicles with Improved Performance’, Journal of Metals, April 2006, p. 56–61. F.L. Luborsky, Amorphous metallic alloys, Butterworth and Co. Ltd, 1983, pp. 534, 831154–1074.. Y. Kawamura, ‘Liquid Phase and Supercooled Liquid Phase Welding of Bulk Metallic Glasses’, Materials Science and Engineering A, 375–377, 2004, p. 112–119. ASM Handbook Volume 6, Welding, Brazing and Soldering, ASM International, 1993.
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26 Ceramic/metal bonding A W U , Tsinghua University, P.R. China
26.1
Introduction
In crystalline form, ceramics are generally considered to be the inorganic compounds between metals (or semi-metals) and nonmetals. Ceramics are covalently and ionically bonded materials and have some unique and excellent properties, such as electrical insulation, chemical and size stability and resistance to the effects of heat. Due to ceramic materials’ wide range of properties, they are used for a multitude of applications. For example, alumina, aluminum nitride and beryllium oxide are the base materials most widely used in electronic packaging because of their high electrical resistance. Zirconia is used as the solid electrode of SOFCs due to its ionic conduction properties at high temperature and as the oxygen sensor material for its sensitive electrical properties. It is also an ideal biomaterial because of its chemical stability, biocompatibility, size stability and high strength and toughness. It is also the material of ferrules for optical fiber connectors because of its wear resistance and mechanical properties. Joining a metal to a ceramic is required in many advanced applications, such as microchip substrates, capacitors and heat sinks. For example, in power amplifier packaging1, ceramic aluminum nitride is required to be joined to a copper heat sink and lead frame. The application of ceramic zirconium oxide as biomaterial requires joining between the ceramic and metal, for example, to make implantable microstimulators2. In the miniature manufacturing field, the joining of ceramics and metals is necessary and unavoidable. This chapter gives an introduction to the characteristics of some typical ceramics used in miniature manufacturing, the major methods (mainly brazing and diffusion bonding) to join ceramics to metals and the main difficulties exhibited in the joining process.
26.2
Characteristics of typical ceramic materials
In miniature manufacturing, ceramics encountered are mainly alumina (Al2O3), aluminum nitride (AlN) and zirconium oxide (ZrO2). 758 WPNL2204
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26.2.1 Alumina Aluminum oxide, also called alumina, is a type of ceramic with α-Al2O3 as the main composition, the higher the content of crystalline aluminum oxide, the better the performance, offsetting the more complex processing and the higher cost. The most common features of aluminum oxide ceramics are high hardness (could reach HRA87 at 760 °C and could hold HRA82 at 1200 °C), good wear resistance and high heat resistance allowing them to be used at a high temperature 1600 °C for a long period. They also possess strong corrosion resistance and good electrical insulating properties, which are outstanding at high frequency. Alumina has good chemical stability, not reacting with most molten metals except Mg, Ca, Zr and Ti at high temperatures. Alumina can dissolve into hot concentrated sulfuric acid, hot hydrochloric acid and hydrofluoric acid just causing some corrosion. The evaporation and decomposition pressures of alumina are both low. The shortfalls in alumina’s properties are low plasticity and poor thermal shock resistance. Alumina ceramics are used mainly to manufacture vacuum devices and as the substrate of electronic circuits. Table 26.1 shows the typical characteristics of alumina ceramics3.
26.2.2 Zirconium oxide ceramics Zirconium oxide ceramics have a high melting and boiling point and high hardness. They are insulators at room temperature and conductors at high temperatures. They are usually used as functional and structural materials. The crystal structures of ZrO2 include three types, cubic(C-phase), tetragonal (t-phase) and monoclinic (m-phase). After adding some stabilizers, parts of t-phase could exist at room temperature in a metastable state, called partially stabilized zirconia (PSZ). The transition of t-phase to m-phase occurring under stress is called ‘stress-induced transformation’. This transformation will absorb energy, relaxing the stresses of the crack tips and increasing the retarding force of crack extension, thus improving fracture toughness. The fracture toughness of PSZ is much higher than other structural ceramics. Among the number of ZrO2 ceramics developed nowadays, stabilizers commonly used are MgO, Y2O3, CaO, CeO2. As a structural material PSZ ceramics are mainly used to produce grinding balls, cylinder sleeves, piston heads, valve seating, follow-up units of jaw, ball valve, bearings and debugging tools for electric appliances. ZrO2 ceramic is the main material for making optical fiber contact pins and sleeves because of its characteristic of easy machinability at micron dimension levels with reliable accuracy and similar physical properties to optical fibers. ZrO2 ceramics are also used as biomaterials, such as the caput of articular prosthesis, and as an oral implant material, substituting bone in oral repairs.
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Alumina content (%) Bulk density (kg/m3)
90 3.8
Mechanical characteristics
Vickers hardness (load 500g) (GPa) Flexural strength (MPa) Young’s modulus of elasticity (GPa) Poisson’s ratio
Thermal characteristics
Coefficient of linear thermal expansion (×10–6K–1) Thermal conductivity 20 °C (w/(m·k)) Specific heat (J/(kg·K))
Electrical characteristics
Dielectric strength (V/m) Volume resistivity (Ω·cm)
40–400 °C 40–800 °C
Dielectric constant (1 MHz) Dielectric loss angle (1 MHz) (×10–4) Loss factor (×10–4)
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92 3.6
×103
96 3.7
×103
99.7 3.9 × 103
12.7 320 320 0.23
12.3 340 280 0.23
13.7 350 320 0.23
17.2 380 380 0.23
7.3 8.1 12 0.75
6.9 7.8 18 0.78
7.2 7.9 24 0.78
7.2 8.0 32 0.79
12 20 °C 300 °C 500 °C
×103
×103
×106
1011 107 105 9.8 20 190
×103
16 × 106 >1014 1012 1010 9.0 6 54
×103
15 × 106 >1014 1010 108 9.4 4 38
×103
15 × 106 >1014 1013 1010 9.9 1 10
Microjoining and nanojoining
Table 26.1 Characteristics of Al2O3 ceramics3 (for IC packages, semiconductor processing equipment parts)
Ceramic/metal bonding
761
ZrO2 ceramic stabilized with Y2O3 has sensitive electrical properties, being a functional material mainly used as solid electrolyte and the positive electrode of solid oxide fuel cell (SOFC). Table 26.2 shows the typical characteristics of zirconia ceramics3.
26.2.3 Aluminum nitride ceramics Aluminum nitride is a potential candidate for replacing alumina, because it has ten times higher thermal conductivity than alumina. The thermal expansion coefficient of AlN (4.8 · 10–6 K–1 at RT-673 K) closely matches the expansion coefficient of silicon wafers (2.7 · 10–6 K–1), hence a component supported on AlN substrate is less likely to fail due to thermal cycling than that supported on beryllium oxide or alumina substrates. Aluminum nitride also has a low dielectric constant and loss, and is nontoxic in nature with a less temperature sensitive thermal conductivity compared to the thermal conductivity of beryllium oxide. Table 26.3 shows the typical properties of aluminum nitride ceramic3. The drawbacks of aluminum nitride ceramics include easy oxidation at temperatures above 800 °C, instability in humid atmospheres and poor corrosion resistance in strong alkaline conditions.
26.3
General difficulties in joining of ceramic/metal
There exist many differences between ceramic and metallic materials, such as the atom bond configuration, chemical and physical properties, etc. These differences make the joining of ceramics to metals difficult. The main difficulties include: 1. Ionic bonds and covalent bonds are characteristic atomic bond configurations of ceramic materials. The peripheral electrons are extremely stable. Using the general joining method of fusion welding to join ceramics with metals is almost impossible, and the molten metal does not generally wet on ceramic surfaces. When joining ceramics to metals with the brazing method, metallization on the ceramic surface is necessary with general inactive brazing filler metal or the use of active brazing alloys in order to get a reliable joint. 2. The thermal expansion coefficients of ceramics are generally much lower than metals. Stress will be generated in the ceramic/metal joint due to the thermal expansion mismatch and will degrade the mechanical properties of joint and can cause joint cracking immediately after the joining process. The thermal stress in the joint due to the thermal expansion mismatch should be carefully considered when joining ceramic with metal. 3. Many ceramics have low thermal conductivity and susceptibility to thermal shock. Using the fusion welding method to join ceramics by concentration heating or with a high energy density heat source, cracking in the ceramic
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Bulk density (kg/m3)
Mechanical characteristics Vickers hardness (load 500 g) (GPa) 10.7~13.2
(5.6~6.0)×103
Flexural strength (MPa) 710~1470
Young’s modulus of elasticity (GPa) 200~220
Poisson’s ratio 0.31
Thermal characteristics Thermal conductivity (w/(m·k))
Coefficient of linear thermal expansion (×10–6K–1)
Specific heat (J/(kg·K)
3
40–400°C 10~10.8
(0.46~0.48)×103
40–800°C 10.5~11.4 Electrical characteristics
Volume resistivity (Ω·cm) 20 °C 300 °C 500 °C
108~>1014 106 103~104
Dielectric constant (1 MHz)
Dielectric loss angle (1 MHz) (×10–4)
Loss factor (×10–4)
Dielectric strength (V/m)
28~33
16~17
476~520
(11~13)×106
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Table 26.2 Characteristics of ZrO2 ceramics3
Table 26.3 Characteristics of AlN ceramics (for IC packages)3 Bulk density (kg/m3)
3.4
Mechanical characteristics Vickers hardness (load 500 g) (GPa) 10.4~10.8
× 103
Flexural strength (MPa) 290~310
Young’s modulus of elasticity (GPa) 320
Poisson’s ratio 0.24
Thermal characteristics Thermal conductivity (w/(m·k))
Coefficient of linear thermal expansion (×10–6K–1) 40–400 °C 4.8
150~160
Specific heat (J/(kg·K)) (0.71~0.72)
Electrical characteristics Volume resistivity (Ω·cm) 20 °C 300 °C 500 °C
>1014 109 107
Dielectric constant (1 MHz)
Dielectric loss angle (1 MHz) (×10–4)
Loss factor (×10–4)
8.7
2~3
17~26
Ceramic/metal bonding
× 103
763
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easily occurs. It is necessary to reduce the temperature gradient in and around the fusion zone as much as possible and to carefully control the heating and cooling speed during the joining process. 4. Ceramics have high melting points, high strength (in compression) and a high hardness value. Direct diffusion bonding of ceramics is very difficult, requiring the finished surface to be extremely smooth and clean, the bonding temperature high and a long bonding time. For example, when joining Si3N4 ceramic directly by diffusion bonding, the finish of the ceramic surface should be more than 0.1 µm and bonding temperature should be 1500~1750 °C. Therefore, joining ceramics by diffusion bonding is normally indirect, using soft, ductile metals as an interlayer which can decrease the bonding temperature considerably and the plastic deformation of interlayer metals can lower the requirements of processing the joining surface of the ceramic. 5. Most ceramics have weak or no electrical conductivity. It is hard to join ceramics by using electrical welding methods unless special techniques used.
26.3.1 Metallurgical inconsistencies between ceramic and metal The wettability of molten metal on ceramic is usually characterized by the contact angle formed in the liquid-solid-gas phase interface, as shown in Fig. 26.1. When the system reaches equilibrium, the interface and surface in tension could be expressed by Young’s equation as formula (26.1).
γsv – γsl = γ lv cos θ
(1)
where γsv is the surface tension at solid-vapor interface, γsl is the surface tension at solid-liquid interface, γlv is the surface tension at liquid-vapor interface. Generally, liquid metal has poor wettability on ceramic because the γsl is too high. The requirement of a liquid wetting on a solid surface (contact angle θ < 90°) is wa > γ lv, wa the work of adhesion, which mainly includes γlv
γsv
θ γsl
26.1 Schematic illustration of interface in tension.
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765
physical work and chemical work (coming from the reaction of liquid-solid interface). It has been found that physical work relates little to the atomic type of metal, type of ceramic and the joining temperature, the maximum value not exceeding 600J/m2, while the γlv of liquid metal is generally over 1000 J/m2, so simple physical interactions cannot make the liquid wet the ceramic surface well. The chemical reactions at the interface which could increase wa will promote the wettability of liquid metal on the ceramic surface. For generating chemical reactions at the interface of ceramic/metal, some active elements which could react with ceramic must be added to the liquid metal. Taking Si3N4 ceramics, for example, metals required to be the active elements should meet the following two requirements4,5. First, these metal elements could form stable nitrides. Second, they could react with ceramics. Zr, Ti, Hf, Al, Nb, Ta, V, Cr, Mo and Fe can form stable nitrides. These ten metals can be divided into two classes. For the first class, including Ti, Zr, Nb, V, Hf, Ta and Al, their Gibbs energy ∆G° of the reaction with Si3N4 ceramic is negative, meeting the above two requirements. These elements are defined as the active elements to Si3N4 ceramic. For the second class, including Cr, Fe and Mo, their Gibbs energy ∆G0 of reaction is positive, making them unable to meet the two requirements to be active elements promoting the wetting of liquid metal on Si3N4 ceramic. They could be defined as the inert elements to Si3N4 ceramic. Active elements promote the wetting of filler metal on ceramic surfaces mainly by their selective segregation to the solid–liquid interface and through reactions with ceramics. Whether active elements in the filler metal segregate to the solid–liquid interface not only depends on the interface reactions of elements with ceramics, but also on their existing state in the filler metal, which means it also depends on the interaction of the active elements with other inert elements of the filler metal. For example, among the Sn-base active filler metals, the Sn-Ti filler metal displays the best wettability as fewer Sn-Ti compounds are found in its metallurgical structure and most of Ti in the filler metal segregates to the solid-liquid interface. While with SnV and Sn-Ta filler metals, active elements V and Ta perfectly exist as Sn-V and Sn-Ta compounds and display poor wettability on ceramic surface. In summary, the selective segregation of active elements to the solid– liquid interface is necessary for the active filler metal to wet the ceramic surface by chemical reactions. For ensuring segregation, active elements need to react with ceramics and produce stable compounds, the Gibbs energy ∆G° of reaction should be negative, and a proper interaction force between active element atoms and solvent atoms of the filler metal is also necessary to ensure the active elements easily move from the inside of liquid to the solid–liquid interface. When joining Si3N4 ceramic with active filler metal, there exists a critical
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temperature of wetting. Only when the joining temperature is higher than a specific temperature range will the active filler metal wet the ceramics. Otherwise the filler metal will not wet the ceramics. The reason that the critical temperature exists is because of the critical reaction temperature, above which the active element will react with ceramics.
26.3.2 Mismatch of the physical and mechanical properties between ceramic and metal Another difficulty in joining ceramic to metal is the difference of thermal expansion coefficients between them. The thermal expansion coefficient of ceramics is generally around 10–6K–1, for instance, the thermal expansion coefficient of Si3N4 ceramic is near 3 × 10–6K–1, while the average thermal expansion coefficient of common metals such as carbon steel or stainless steel is about 14–18 × 10–6K–1. Such mismatch of thermal expansion coefficients leads to residual stresses in the joints after cooling to room temperature from the joining temperature, the joint performance is thereby affected. The residual stresses produced in the ceramic metal joint could be estimated according to eqn 26.2 for fully elastic conditions.
σc =
∆α × ∆T × E m × E c ( E m + Ec )
(2)
where σc is the residual stress after the joint cools to room temperature, ∆α is the difference of thermal expansion coefficient between materials, ∆T is the difference between joining temperature and room-temperature, Em is Young’s model of metal, Ec is Young’s model of ceramic. If the thermal stresses in the metal exceed its yield strength, the residual stresses in the joint could be determined by eqn 26.3.
σc = σmy + ∆α · ∆T · Emp
(3)
where Emp is the linear strain hardening coefficient and σmy is the yield strength of the metal (linear elastic-linear plastic conditions are assumed). For reducing the residual stress induced by the mismatch of the thermal expansion coefficient between the materials to be joining, the following methods can be used4,6–10. 1. Using soft filler metals. Soft filler metals have low yield strength and could release the residual stress. 2. Using soft interlayer. Residual stress could be reduced by the elastic and plastic deformation of an interlayer, e.g. when using Al or Cu as interlayer, the residual stress is decreased. According to eqn 26.2, the residual stress will decrease with Young’s model Em decreasing.
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3. Using hard metals of which the thermal expansion coefficient is close to ceramics as the interlayer. Using hard metals such as W, Mo or Invar as the interlayer, could reduce the residual stress. Their validity is not obvious when hard metals with high yield strength are the interlayer. 4. Using composite interlayer. Composite interlayers often constitute hard metals and soft metals, like Cu/Mo-Cu/Nb, have a noticeable effect on reducing residual stress, with a combination of merits of those two kinds of metals. 5. Joining under low temperature. Joining ceramic to metal at a low temperature is good for reducing the joint deformation and effectively decreasing the residual stresses. 6. Heat treatment after joining. Proper heat treatment post-joining sometimes releases the stress and the strength will vary based on the heat treatment. 7. Optimized design of joint configuration. Appropriate configuration of the joint could decrease the stress concentration extent and reduce the residual stress.
26.4
Brazing ceramics to metal
There are many ways to make ceramic-metal joints, such as mechanical bonding, electrostatic bonding, hot isostatic pressing bonding, brazing and diffusion bonding. Each method has its own features and specific applications. Among them, brazing and diffusion bonding are the most commonly used methods, for the high reliability and good repeatability characteristics of joints. Brazing ceramic to metal can be direct or indirect. Direct brazing uses filler metal containing active elements, oxides, fluorides or solder to join ceramic and metal. Indirect brazing is used first to achieve ceramic surface metallization, then conventional brazing filler metals are used to join them.
26.4.1 Metallization of ceramic surface The ways to realize ceramic surface metallization include thick-film metallization (at low or high temperature), thin-film metallization (such as Ti/Pd/Au), chemical plating (such as plating Ni), co-fired, direct copper bonding and active elements metallization. In this chapter, thick-film, thinfilm and direct copper bonding are briefly introduced. Thick-film metallization11 is a method of producing a metal layer for sealing, conduction and resistance to the ceramic substrates by screen printing, forming the layer for brazing, electric circuits and connection points. Thickfilm slurries are generally made by mixing and milling the metal powders at a size of 1.5 microns, adding a percentage of permanent adhesive, organic solvents, thickeners and surfactant, etc.
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Metal slurries used for thick-film could be Cu, Ag and Au. Adhesives are usually Si2O3-B2O3-ZnO, Si2O3-B2O3-BaO-Al2O3, Si2O3-B2O3-BaO-CaO (for the AlN ceramic thick-film metallization). The sintering atmosphere for Cu slurry is nitrogen, with a sintering temperature of 850 °C. For Ag slurry in the air, the sintering temperature is 920 °C, For Au slurry also in air, the sintering temperature is 850 °C. Thin-film metallization, by using PVD methods like sputtering12 or evaporation, is used to produce a ceramic surface coated with thin metal films, such as NiCr/Pd/Au, Ti/Pd/Au, Ti, Ti/Ag, Ti/Ni, etc. The DCB technique is based on the presence of an eutectic between copper and Cu2O which wets alumina. The process can thus be successfully used only if the AlN surface is oxidized in alumina, in air or in an oxygen atmosphere. AlN was held at 1200 °C in air for some minutes to obtain a certain thickness of the alumina layer. Copper was also oxidized at 1000 °C before bonding. The two oxidized materials were pressed together and held at 1070 °C for 1 ± 2 min. Then, the eutectic phase was used to bring the copper in intimate contact with AlN by wetting and reacting with alumina to form the binary oxide CuAlO2. The composition of the bonding atmosphere is very important since the initial Cu2O must be maintained. If the oxygen partial pressure is too high, all copper is converted into the eutectic melt, conversely, if the partial pressure is less than the equilibrium partial pressure over Cu2O at 1065 °C, oxide is reduced and the eutectic phase will not form. The bonding was generally performed in argon with 70 ppm O213.
26.4.2 Vacuum brazing with active filler metals The chemical bonds of ceramic materials are mainly ionic bonds and covalent bonds, which means ceramics are very stable in chemical reactions. Filler metals with metallic bonds can wet ceramic surfaces when chemical reactions occur between filler metal and ceramic. Transition metals such as Ti, Zr, Hf, Nb, Ta, could make the ceramic surface decompounded by reacting with the ceramic, forming the reaction layer. The reaction layer is mainly compounds that have metallic characteristics, so filler metals can wet the ceramic surface easily. In active brazing alloys (filler metals) Ti is usually used as the active element. Among the commercial filler metals, such as Ag-base, Cu-base or Ag-Cu eutectic filler metal, the content of Ti is 1~5wt%. Sometimes an amount of In is added to the brazing alloys to improve the fluidity and increase the activity of active elements. Besides Ag-base and Cu-base filler metal, there are some other active brazing alloys based on Sn or Pb, their melting point below 300 °C. The working temperature of the ceramic-metal joint bonded with Ag-base or Cu-base brazing alloys are generally not above 400 °C, however, with high melting temperature and noble metals such as
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Pt, Pd, Ni, Co, Au, that could reach to about 800 °C. Some common active brazing alloys with their chemical composition and melting points are listed in Table 26.414–15. In active brazing, the protection of active elements is very important. These active elements are easily oxidized and if that were to occur they would be unable to react with ceramics. Active brazing is usually carried out in a vacuum or an inert atmosphere. In vacuum brazing the vacuum under brazing temperature is usually better than 10–4mbar. Foil forms of filler metals with thickness 50–200 µm is convenient for use in the brazing process compared with powder formulations. In this type, the possibility of active elements being oxidized before brazing is very low, that is because active elements distribute in the base metal homogeneously and are protected. There are many factors that may influence the brazed joint quality. These include the properties of ceramic and metal (including the interlayer metal), the brazing alloy system, the amount of brazing filler metal used, the active element type and its content, the surface structure, atmosphere, brazing temperature and holding time, the size of joint and so on. When ceramics have high fracture toughness, the corresponding joint has high tension strength. The properties and thickness of the interlayer metal have a great influence on the joint strength. Soft and easily yielding interlayer metals favor releasing the stress generated in the joint because of the mismatched heat expansion Table 26.4 Some common active brazing alloys14–15 Brazing alloy
Composition (wt%)
Solidus temperature °C
Liquidus temperature °C
Ag–Cu–Ti Ag–Cu–Ti Ag–Cu–Ti Ag–Cu–In–Ti Ag–Cu–In–Ti Ag–In–Ti Ag–Ti Sn–Ag–Ti Ag–Cu–Ti–Li Ag–Cu–Sn–Ti Pd–Ni–Ti Pd–Cu–Pt–Ti Ag–Pd–Ti Ag–Pd–Pt–Ti Pt–Cu–Ti Zr–Cr–Cu Ni–Hf Co–Ti Au–Pd–Ti
70.5–26.5–3 72–26–2 64–34.5 –1.5 72.5–19.5–5–3 59.5–24–15–1.5 98–1–1 96–4 86–10–4 68–28–2–10 60–28–10–2 58.2–38.8–3 51–43–2–4 56–42–2 53–39–5–3 55–43–2 73–12–15 70–30 90–10 90–8–2
780 780 770 730 605 950 970 221 640 620 1204 1099
805 800 810 760 755 960 970 300 720 750 1239 1170
1195 1208
1250 1235
1200 1215 1148
1225 1320 1205
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behavior of ceramics and metals, the joint performance can be improved when the thickness of the interlayer metal is correct (shown in Fig. 26.216). The content of active elements in filler metals affects the joint strength by its influence on the interface reaction and the microstructures of the brazing seam (Fig. 26.316). There are favorable ranges of brazing temperature and holding time for the joint strength (Fig. 26.416). If the brazing temperature is too high or too low and the holding time is too long or too short, high-strength joints are unlikely. The strength of larger joints is usually less than those of smaller size.
80 70
φ 10
60 50
Cu
40
Ni
8
σb /MPa
Ag
Si3N3 /AgCu3Ti/M/AgCu3Ti/Steel Φ10 1153 K × 600 s, interlayer thickness 1.5 mm
φ 18
30 20
Steel Mo
W
Kovar
10 0 Metal (a)
160 140
Si3N4 /AgCu3Ti /Cu/AgCu3Ti/A3 Φ10 1153 K × 600 s
σb /MPa
120 φ 10
100
8
80 φ 18
60 40 0
5
10 15 20 25 Thickness of Cu film (mm) (b)
30
26.2 Effects of interlayer type and thickness on ceramic-steel joint tension strength16 (a) effect of interlayer type on joint strength, (b) effect of interlayer thickness on joint strength.
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40 Si3N/Ag-Cu-Ti/Steel Φ10 1173 K × 600 s
35 30 σb /MPa
25 φ10
20 15
8
10 φ18
5 0 0
1
2
3 4 Ti content/wt. %
5
6
26.3 Effect of Ti content in AgCuTi filler metal on joint strength.16
Brazing temperature (usually above the liquidus temperature about 50~100 °C) and bonding time are key parameters for joint quality. The temperature of brazing ceramics to metals is usually between 800 °C and 1100 °C. Snbase or Pb-base filler metal with low melting points should also be brazed at high temperature to achieve sufficient thermodynamic activity to make chemical reactions. ZrO2, Al2O3, Si 3N4, SiC and AlN are widely used in applications. Commercial filler metals based Ag, Cu or eutectic Ag-Cu could wet those ceramics and aid in joining them reliably. Table 26.5 shows the results of joint strength by brazing ceramics with various active brazing alloys17. Using AgCuTi filler metal to braze common engineering ceramics, the results of joint strength is shown in Table 26.6.
26.5
Vacuum diffusion bonding
The solid state diffusion bonding technique was originally used to join dissimilar materials. It is now one of the general methods used to join ceramics, with direct bonding or with an interlayer. Compared with fusion welding, the advantages of solid state diffusion bonding include high joint strength, small shrinkage and deformation, accurate dimension control and good fit for joining dissimilar materials. Because bonding is usually carried out in a vacuum atmosphere at high temperature and with a long bonding time, the disadvantages of diffusion bonding are expensive equipment, high production costs and limited sample dimension. The factors influencing diffusion bonding include the bonding temperature, the holding time, the pressure applied, the atmosphere, the surface structure,
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Microjoining and nanojoining 35 30 φ10
σb / MPa
8
25 20 15
φ18
10 5
Si3N4 /Ag-Cu-3Ti/Steel Φ10 600 s
0 1100
1150
1200 1250 1300 Brazing temperature/K
1350
(a)
35 30 25 σb / MPa
φ10
20
8
15 10
φ18
Si3N4 /Ag-Cu-Ti/Steel Φ10 1173 K × 600 s
5 0
200
400
600 800 Holding time s
1000
1200
(b)
26.4 Effects of brazing temperature and holding time on ceramicmetal joint strength16 (a) effect of brazing temperature, (b) the effect of holding time.
the chemical reactions and the similarity of physical properties of materials to be joined16.
26.5.1 Effect of bonding temperature Temperature is the most important parameter in diffusion bonding. The temperature of solid state diffusion bonding ceramics to metals generally needs to be over 90% of the melting point of the metal. In solid state diffusion bonding, the inter diffusion of elements allows
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Table 26.5 Four-point bending strength of joints with various filler metals17 Joint
Filler metal
Average bend strength (MPa)
Si3N4–Si3N4 Si3N4–Si3N4 Si3N4–Si3N4 Si3N4–Si3N4 Si3N4–Si3N4 (pre-metalization) SiC–SiC SiC–SiC SiC–SiC SiC–SiC SiC–SiC AlN–AlN AlN–AlN Al2O3– Al2O3 Al2O3–FeNiCo Si3N4–AISI304 Si3N4–AISI304 Si3N4–AISI304 AlN–FeNiCo Si3N4–FeNiCo SiC–FeNiCo
AgCuTi3 AgCuTi5 AgCuZr5 AgCuHf5 PdNiTi3 AgCuTi3 AgCuTi5 AgCuZr3 CuPtNb3 AuPdTi2 CuSiAlTi2.25 AgCuTi3 AgCuTi3 AgCuTi3 AgCuTi3 AgCuZr3 AgCuHf4 AgCuTi5 AgCuTi5 AgCuTi3
225.2 184.2 108.7 130.1 163 85.2 107.3 52 65 167 98.5 168.9 225 182 84 148 36.5 186 35
chemical reactions to occur and form sufficient interface combination. The formation of the reaction layer and its thickness affects the joint strength remarkably. The thickness of the reaction layer could be estimated according to eqn 26.4. –Q
x = K p × t n = k 0 × t n × e RT
(4)
where k0 is a constant, n is the exponential of time, Q is the activation energy, related to the diffusion mechanism and R is the gas constant. The influence of bonding temperature on joint strength has the same tendency with reaction layer thickness. According to the experimental results of tension strength tests, the influence of bonding temperature on joint strength could be expressed by eqn 26.5
BS = B0 e
– Qapp RT
(5)
Where B0 is a constant (MPa), Qapp is the appearance activation energy, the sum of various activation energy. According to the equation, the increase of temperature favors an increase in the joint strength. Figure 26.5 shows the relation between joint strength and temperature, for bonding alumina to steel using an Al film of thickness 0.5 mm as the interlayer. Joint strength may decrease with higher temperature, because the residual
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Base materials
Joint strength (MPa)
Base materials
Joint strength (MPa)
Si3N4–1Cr13
Bend strength 385~415
95%Al2O3–95%Al2O3
Tension strength >90.8
Si3N4–40Cr
Bend strength 400
95%Al2O3–oxyen-free copper
Tension strength 84.9~95.5
PSZ–1Cr13
Bend strength 77~514
95%Al2O3–kovar
Tension strength 76.4~105.7
PSZ–40Cr
Bend strength 407
Al2O3–Fe41Ni
Shear strength 265
SiC–SiC
Bend strength 350
99 %Al2O3–Ti
Shear strength 120
AlN–copper
Tension strength 50~110
99 %Al2O3–Nb
Shear strength 120
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Table 26.6 Strengths of ceramic-ceramic or ceramic-metal joint brazed with Ag-Cu-Ti filler metal17–18
Ceramic/metal bonding
775
70
Tension strength (MPa)
60 50 40 Measured value
30
Fitting line
20 10 0 400
450
500 550 600 Bonding temperature (°C)
650
700
26.5 Effect of bonding temperature on joint strength.19
stress will also increase and reduce the joint strength. So, when the negative effect brought by residual stress is equal to the positive effect brought by increasing temperature, the joint strength could reach the maximum value. Moreover, increasing the temperature may result in a change in ceramic performance or cause the brittle phase to occur and the joint to fail. In addition to these factors, the strength of ceramic-metal joints is also related to the melting point of the metal. Among ceramic-metal joints, the joint strength increases linearly with increasing of melting point of metals (Fig. 26.6).
26.5.2 Effect of holding time The holding time not only affects the thickness of the reaction layer, but it also affects the interface reaction products19. The effect of time on the thickness of reaction layer (X) could be approximately shown by X = k (Dt)1/2(k is a constant and D is the diffusion coefficient). It can be seen that increasing time will make the reaction layer grow. In general, the relation between bonding time (t) and tension strength is BS = B0×t1/2, where B0 is constant. But, under certain temperatures it has been found that there exists an optimum bonding time range, as shown in Fig 26.7, which shows the effect of bonding time on joint strength when bonding Al2O3 to Al metal. When bonding Si3N4 to Invar there exists an optimum bonding time range. Bonding SiC to SUS304 with Nb as the interlayer metal20, the phase NbSi2 that has low strength and a large heat expansion coefficient compared with SiC occurred and made the joint shear strength
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Tension strength (MPa)
40
30 Fe 20 Cu
Pt
Ni
Al
10 Pb
0 0
500
1000 Melting point (°C)
1500
2000
26.6 Effect of the metal melting point on the strength of aluminametal joints.19 70
Tension strength (MPa)
60 50 40 30 20 10 0 0
20
40
60 80 Bonding time (min)
100
120
26.7 Effect of holding time on the joint strength.19
decrease after the long bonding time Fig. 26.8. When using V as the interlayer metal to bond AlN ceramic, the joint shear strength was also decreased for the formation of brittle phase V5Al8 after the long holding time21.
26.5.3 Effect of pressure Pressure applied in the solid phase diffusion bonding process creates plastic deformation, which could reduce the size of surface asperities and provide
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777
250
Shear strength (MPa)
200
150
100
50
0 0
20
40 60 80 Holding time (ks)
100
120
26.8 Effect of holding time on the shear strength of SiC/Nb/SiC joint.20
oxide breakdown, increasing the surface contact area, to offer the best condition for the diffusion of atoms. For preventing the macroscopic deformation of components, the bonding pressure is usually within the range of 0~100 MPa, while under that pressure sufficient deformation cannot be created. Increasing the pressure within certain limits could enhance the joint strength, as shown in Fig. 26.9 (effect of pressure on joint shear strength in bonding Ag to Al2O3, Cu to Al2O3, Al to SiC). As well as the bonding temperature and bond time, there also exists an optimum pressure range for joint strength. The best pressures for bonding Al to Si3N4 and Ni to Al2O3 are 4 MPa and 15–20 MPa respectively. Other factors, like the bonding material type and the surface oxide thickness, will change the effect of pressure on joint strength. As shown in Fig. 26.10, when bonding alumina with the noble metal Pt, joint strength rises to a constant with the increase of pressure.
26.5.4 Effect of chemical reactions Generally, when joining ceramic to metal with solid diffusion bonding either directly or using a metal interlayer, compounds will be formed at the interface between ceramic and metal. Reaction compounds change with the diffusion bonding conditions, resulting in a variation of joint performance. When bonding SiC with Nb interlayer at a temperature of 1790 K20 and a bonding time less than 7.2 ks the joint consists of SiC/Nb5Si3Cx/Nb2C/Nb, while for a bonding time between 7.2–72 ks the joint structure is SiC/NbC/Nb5Si3Cx/ Nb2C/NbC/Nb, At a bonding time longer than 72 ks, Nb disappears, and the
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Microjoining and nanojoining 80
Shear strength (MPa)
70
Al/Al2O3
60 Cu/Al2O3
50 40 30 20
Al/SiC
10 0
2
4 6 Bonding pressure (MPa)
8
10
26.9 Effect of pressure on the shear strength of joints.22
300
Bending strength (MPa)
250
200
150
100
50 0
2
4 6 8 Bonding pressure (MPa)
10
12
26.10 Effect of pressure on the strength of Al2O3-Pt diffusion joints.19
joint structure becomes SiC/NbC/NbSi2/NbC/ NbSi2/NbC/SiC, with the joint strength decreased after the appearance of NbSi2.
26.5.5 Effect of bonding atmosphere Usually, the joint strength bonded in a vacuum is higher than that in an argon atmosphere or in air. In the bonding of Si3N4 using aluminum, the joint
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strength in the vacuum is highest. The fracture locates both in the Al layer and the ceramic. The fracture in the Al layer is plastic and in the ceramic it is brittle. While bonding in air, the fracture occurs at the interface between Al and Si3N 4 because of the oxidation of Al. Although the pressure could break down the surface oxidizing film, the new oxidizing film forms continuously when the partial pressure of oxygen in the atmosphere is higher. If diffusion bonding of ceramic Si3N4 is conducted at high temperature (1500 °C), the ceramic Si3N4 will decompose forming cavities. When bonding in an N2 atmosphere, the decomposition will be limited. The higher N2 pressure is favorable to joint strength. The flexural strength of the joint bonded in 1 MPa N2 is higher by one-third than that in 0.1 MPa N2.
25.5.6 Effect of mismatch in coefficient of thermal expansion Because of the low coefficient of thermal expansion of ceramic, the ceramic will be under pressure and metal will be tensed in the joint of ceramic and metal. A soft interlayer will make the joint stress complex. When bonding alumina and metal using Al as the interlayer, the joint strength decreases as the thermal expansion coefficient of the metal increases. Similar trends have been reported for other ceramics such as SiC and Si3N4 and Sialon.
26.5.7 Effect of interlayer The use of interlayer in diffusion bonding is required for many applications in order to reduce the bonding temperature, bonding pressure and bonding time, to enhance diffusion, to eliminate impurity elements, and to reduce the residual stresses generated at the bond interface if the interlayer is ductile. When bonding stainless steel to alumina, the lower residual stress caused by the interlayer can be seen from Fig. 26.11. The residual stresses decrease with the increase of the interlayer thickness. The effect of interlayer Nb is evidently due to its thermal expansion coefficient being close to that of alumina. The effect of the interlayer is complex if the reaction layer at the interface forms. The effect of the interlayer will differ by the reaction production type and its thickness. Joint performance is highly dependent on the interlayer selection. If the interlayer is inappropriate, the joint strength will decrease because of the possible brittle reaction; the higher residual stress caused by mismatch of the thermal expansion coefficient between ceramic and metal, or the corrosion resistance of the joint can be reduced. The interlayer can be in the form of a powder, foil or metallization film.
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Microjoining and nanojoining 1200 Ti Mo Cemet Nb
Residual stress (MPa)
1000 800 600 400 200 0 0.0
0.5
1.0 1.5 2.0 The interlayer thickness (µm)
2.5
26.11 Effect of the interlayer thickness on the residual stress of Al2O3AISI405 diffusion joints.19
26.5.8 Effect of surface state The joint strength may be significantly influenced by the roughness of the faying surfaces. High surface roughness may introduce high local stress concentrations with subsequent initiation of brittle fracture. The effect of the surface roughness on the joint strength can be seen in Fig. 26.12 for Si3N4Al joints. Joint strength decreases from 470 MPa to 270 MPa when the surface roughness value varies from 0.1 µm to 0.3 µm.
26.5.9 Performance of joints diffusion bonded with various parameters The joint strength of various material combinations bonded by solid state diffusion bonding with the related parameters is shown in Table 26.7. Most of the data in the table comes from reference19.
26.6
Typical applications
26.6.1 Zirconia to Ti-6Al-4V braze joint for implantable biomedical device2 The BION® microstimulator is a wireless network of up to 255 singlechannel stimulators controlled and powered by a radiofrequency (RF) link from a central external controller. A ceramic-to-metal brazed case was designed for the BION® microstimulator, where the ceramic provides a window for radiofrequency transmission. The ceramic-to-metal case is 2.44 mm in diameter
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781
600 Base material
Bend strength (MPa)
500 400
300 800°C, 10 min, 0.05 MPa
200 100
0 0.0
0.1
0.2 0.3 Surface roughness (µm)
0.4
26.12 Effect of ceramic surface roughness on the strength of Si3N4-Al joints.19
and is composed of an 11.28 mm long ceramic tube brazed onto a 0.76 mm long titanium cap (cathode) at one end and a 1.98 mm long titanium ferrule (anode) at the other end. The wall of the ceramic tube and the ferrule is 0.28 mm thick. The ceramic and metal components selected for this design are 3 mol% yttria-stabilized tetragonal zirconia polycrystal (3Y-TZP) and a Ti6Al-4V titanium alloy. A strong, hermetic, reliable, and biocompatible ceramic to metal seal is desirable. With 0.050 mm thick titanium nickel (TiNi)-clad braze filler material (which is a three-layer laminated structure configured as Ni/Ti/Ni, with each of the Ni layers having a thickness of about 0.008 mm), the ceramic-to-metal case was brazed at 1035 °C for five minutes under a vacuum of 1.3 × 10–5 mbar. Laboratory tests show that the ceramic-to-metal seal is hermetic, strong, and resistant to electrochemical corrosion.
26.6.2 Metallized AlN ceramics brazed to metals applied in power amplifier packages1 Figure 26.13 shows the schematic cross-section of an AlN pill brazed to a copper flange and Alloy 42 leads. In this package, ceramic AlN was first metallized with the thin-film metallization method. The Alloy 42 leads and the copper heat sink were brazed with Cu-Ag eutectic directly to the metallized AlN ceramics. The peel strength of the Alloy 42 leads tested was ≥ 7.5 lbs, this value was more than the required peel strength specified for analogous BeO packages (5 lbs). So the AlN packages were found to be highly acceptable.
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Table 26.7 Performance of joints diffusion bonded with various parameters Temperature/°C
Holding time/min
Pressure/MPa
Interlayer and its thickness
Atmosphere
Strength* (MPa)
Al2O3–Ni Al2O3–Pt Al2O3–Al Al2O3–Cu Al2O3–Cu4Ti Al2O3–Fe Al2O3–Carbon steel Al2O3–Carbon steel Al2O3–high alloy steel Al2O3–Cr Al2O3–Pt–Al2O3 Al2O3–Cu–Al2O3 Al2O3–Ni–Al2O3 Al2O3–Fe Al2O3 Al2O3–Ni–Al2O3 Al2O3–Ag–Al2O3 Al2O3–Cu–Al2O3 Si3N4–Invar Si3N4–Si3N4 Si3N4–Si3N4 Si3N4–Si3N4
1350 1550 600 1025–1050 800 1375 1450 1450 625 1100 1650 1025 1350 1375 1250 900 1000 727–877 770–877 1550 1500
20 1.7–20 1.7–5 155 20 1.7–6 120 240 30 15 240 15 30 2 300–600 300 120 7 10 40–60 60
100 0.03–10 7.5–15 1.5–5 50 0.7–10 <1 <1 50 120 0.8 50 50 50 15–20 6 6 0–0.15 0–0.15 0–1.5 21
– – – – – – Co Ni 0.5 mmAl – – – – – – – – 0.5 mmAl 10–20 µmAl ZrO2 –
H2 H2 H2 H2 Vacuum H2 Vacuum Vacuum Vacuum Vacuum Air Vacuum Vacuum Vacuum Vacuum Vacuum Vacuum Air Air Vacuum 1MPa N2
Si3N4–Si3N4
1500
60
21
–
0.1MPa N2
Si3N4–WC/Co
610
30
5
Al
Vacuum
200b (A) 200–250(A) 95 (A) 153b (A) 45b (T) 220–231 (A) 3–4 (S) 0 (S) 41.5b (T) 57–90b (S) 220 (A) 177 (B) 149 (B) 50 (B) 75–80 (S) 68 (S) 50 (S) 110–200 (A) 320–490 (B) 175 (B) 380(A)–RT, 230(A)–1000 °C 220(A)–RT, 135(A)–1000 °C 208b (A)
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Material combination
610 1050–1100 630 1150 1000–1100 1400 1400
30 180–360 300 0–300 90–1440 30 600
SiC–Nb–SUS304 SiC–SUS304 AlN–AlN ZrO2–Si3N4 ZrO2–Cu–ZrO2 ZrO2–ZrO2 ZrO2–ZrO2
1400 800–1517 1300 1000–1100 1000 1100 900
60 30–180 90 90 120 60 60
5 3–5 4 6–10 7–20 1.96
– >14 6 10 10
Al–Si Fe–Ni–Cr – – – –
25 µmV >0.2 mm Ni 0.1 mm Ni 0.1 mm Cu
Vacuum Vacuum Vacuum Vacuum Vacuum Vacuum Vacuum Vacuum Vacuum Vacuum Vacuum Vacuum Vacuum Vacuum
50b (A) >90 (A) 100 (S) 20 (S) 95 (S) 87 (S) 187 (RT) ≥ 100 (800 °C) (S) 125 0–40 120 (S) 57 (S) 97 (S) 150 (A) 240 (A)
* The letter in the parentheses following the strength value represents strength testing method, A represents 4-point bending test, B represents 3-point bending test, T represents tension test, S represents shear test. Superscript b represents maximal value.
Ceramic/metal bonding
Si3N4–WC/Co Si3N4–WC/Co Si3N4–Al–Si3N4 Si3N4–Ni–Si3N4 Si3N4–Invar–AISI316 SiC–Nb SiC–Nb–SiC
783
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Microjoining and nanojoining Air bridge Lead frame Metallized layer Ceramic AIN Copper silver braze
Copper heat sink
26.13 A power amplifier found in a telecommunications application.1
26.7
Summary and future trends
Advanced ceramics are key materials and are widely used in the electronics, fuel cell, sensor and bioengineering fields. Joining of ceramics to metal is necessary and unavoidable in the miniature manufacturing field. Brazing and diffusion bonding of ceramics to metals are essential and reliable joining methods. Joining ceramics and metal at lower temperatures with higher bonding quality will be one of the aspirations of future research work.
26.8
References
1. Richard Koba, Jonathan Harris Robert Youngman, et al. ‘Aluminum-Nitride Packages Provide Consistent Performance’. Microwaves & RF, October 1997, 36(10): 156~166. 2. Guangqiang Jiang, Delta Mishler, Ross Davis, et al. Zirconia to Ti-6Al-4V Braze Joints for Implantable Biomedical Devices. Published online 20 October 2004 in Wiley InterScience, 316–321. 3. //global.Kyocera.com/ 4. Xian Aiping. Mechanisms of the wetting and bonding between ceramics and metals. Ph D Thesis, Shenyang: Institute of Metal Research Chinese Academy of Sciences, 1991 (in Chinese). 5. Passerone A., Muolo M.L. ‘Joining technology in metal-ceramic systems’. Materials and Manufacturing Processes. 2000, 9: 631–648. 6. Peteves Stathis D., Giacomo Ceccone, et al. ‘Joining silicon nitride to itself and to metals’. JOM 1996, 1, 48(1): 48–52, 74–77. 7. Morozumi Shotaro, et al. ‘Joining of silicon nitride with metal foils’. Nippon Kinzoku Gakkaishi/Journal of the Japan Institute of Metals. 1990, 12, 54(12): 1392–1400. 8. Narita Toshio, et al. ‘Measurement of residual thermal stress and its distribution on silicon nitride ceramics joined to metals with scanning acoustic microscopy’. Journal of the Japan Institute of Metals. 1990, 12, 54(10): 1142–1146. 9. Zhou, Y., Bao, F.H., Ren, J.L. North, T.H. Interlayer selection and thermal stresses in brazed Si3N4-steel joints. Materials Science and Technology, 1991, 7(9): 863– 868.
WPNL2204
Ceramic/metal bonding
785
10. Yunhong Zhou, Kenji Ikeuchi, North, Thomas H. Zhirui Wang. ‘Effect of plastic deformation on residual stresses in ceramic/metal interfaces’. Metallurgical Transactions A, 1991, 22A: 2282–2285. 11. Iwase N. ‘Thick films and direct bond copper forming technologies for aluminum nitride substrates’. IEEE Trans CHM T, 1985, 18(2): 253–258. 12. Sheng Zhu, Wladyslaw Wlosinski. ‘Joining of AlN ceramic to metals using sputtered Al or Ti film’. Journal of Materials Processing Technology, 2001, 109, 277–282. 13. Kara-Slimane, A., Juve, D., Leblond, E., Treheux, D. ‘Joining of AlN with metals and alloys’. Journal of the European Ceramic Society 2000, 20, 1892–1836. 14. Chen Dengquan. ‘Progress of filler metals and brazing processes for ceramic to metal brazing’. Precious Metals, 2001, 22(1): 53–56 (in Chinese). 15. Elssner, G. Petzow G. ‘Metal/Ceramic joining’. ISIJ International, 1990, 30(12): 1011–1032. 16. Yang Ping. Vacuum Brazing of Si3N4 ceramics to metals. (Master Thesis) Beijing: Mechanical Engineering Department of Tsinghua University, 1988. 17. Ren Jialie, Wu Aiping. Joining of advanced materials. Beijing: China Machine Press 2000 (in Chinese). 18. Zhao Wenqing, Wu Aiping, Zou Guisheng, Liu Genmao. ‘Brazing between high purity alumina ceramics and titanium’. Transactions of the China Welding Institution, 2006, 27(5): 85–88 (in Chinese). 19. Akselsen O. M. et al. ‘Review: Diffusion Bonding of Ceramics’. Journal of Materials Science. 1992, 27: 569–579. 20. Fen Jicai, Liu Yuli, Naka Masaaki. ‘Diffusion bonding if SiC ceramic with Nb and stainless steel’. Material Science & Technology, 1998, 6(1): 5–7 (in Chinese). 21. EI-Sayed, M. H. et al. ‘Structure and Strength of AlN/V Bonding Interfaces’. Journal of Materials Science. 1998, 33(11): 2869–2874. 22. Treheux, D. et al. ‘Metal Ceramic Solid State Bonding: Mechanisms and Mechanics’. Scripta Metallurgica et Materialia. 1994, 31(8): 1055–1060.
WPNL2204
Index
absorption, in laser microwelding 350, 352, 353 nonlinear and microwelding of transparent materials 356, 357 absorptivity 350, 354–5 accelerating voltage (AV) 435–40, 448–9, 450 accumulated inelastic strain 322, 323 accuracy 161 acid vapour fluxless soldering 329, 330 acrylic based adhesives (acrylates) 506–7 active alignment laser welder for fibre pigtailed TO-can laser diode package 187–200 assembly processes analysis 187–97 active alignment 189–92 laser spot weld 192–5 post-weld shift 195–7 prototype 197–200 active elements active alloying elements and diffusion bonding 244–5 ceramic/metal bonding with active filler metals 765–6, 768–71, 772, 773, 774 metallic brazing of seals for SOFCs 732–4 adaptive control 381–4 additives, for adhesives 510 adhesion 500, 501–4, 633–4 adhesive bonding 500–44 advantages and disadvantages 501 characteristics of adhesive joints 511–12 classification of adhesives 504–10 additives 510
786 WPNL2204
curing mode 506–9 mode of reaction 504–6 defects due to 522 definition 500 future trends 537–40, 541 health and safety 524 hot melts see hot melt adhesives medical devices 536–7, 712–13 microelectronic interconnections and packaging 524–35 anisotropic conductive adhesives 532–3, 535 chip bonding 524–5 dam and fill coating 528, 530–1 flip chip assembly 526–7 glob topping 527–8, 529 isotropic conductive adhesives 529–32, 533, 534 non-conductive adhesives 528, 532 PCB mounting of SMDs 525–6, 527 sealants and sealings 527 thermal conductive adhesives 533–4 underfill of flip chips 528, 532 as microjoining technology 512 other applications 535–7 principle of adhesion and cohesion 501–4 process 512–21 adhesive packages 519 curing and setting 519–21 dispensing techniques 514–18 mixing 518–19 pre-treatment methods 513
Index requirements of surface treatment 512 properties and benefits of adhesive bonds 500–1 quality control 521–2 selection of adhesive technology 523–4 standards 540–3 test methods 522, 542–3 adsorption theory 502 advanced vision system 225 aerodynamic forces 55, 68–9, 72 ageing tests 543 air brazes 734 alternating current (a.c.) power supply 494–5, 496 alumina (aluminium oxide) 16, 758, 759, 760 aluminium 16, 476 bonding using composite silver nanoparticle 255–9 electron beam microwelding 461–2, 463 filler in brazing and soldering 33–4 nanostructured 745–7 aluminium nitride 758, 761, 763 brazing metallised aluminium nitride ceramics to metals in power amplifier packages 781–4 aluminium oxide (alumina) 16, 758, 759, 760 amorphous carbon 556–9 amorphous thermoplastics 665–6, 667 amplitude 676 process control in wedge-wedge bonding 169 anaerobic adhesives 506 anisotropic conductive adhesives (ACAs) 510, 532–3, 535 anisotropic conductive films (ACF) 532–3 annealing diffusion bonding of BSCCO bulks 591–2 direct wafer bonding 636 effect of irradiation-induced defects in CNTs 548–50 anode 719, 721 anode supported SOFCs 718, 719
787
anodic wafer bonding 634, 641–8 glass-to-glass using silicon layer 651–5, 656, 657 low temperature 642–6 mechanisms 646–8 silicon-to-silicon using glass layer 650–1 ANSYS 145–6, 149–54, 155 arc welding 69, 74–5, 127, 234 arc seam welding 84 arc spot welding (percussive arc welding) 69, 74, 84–6, 706 GTAW 138–40, 141 plasma arc welding (PAW) 74 argon 251–3, 254 plasma fluxless soldering 330–1 cleaning and soldering using argon and hydrogen mixture 331–6 solder ball bumping by plasma reflow 336–9 ASM IDEALine 175–8 asperities, surface see surface roughness assembly process automation 174–202, 716 assembly equipment for an in-line process 175–8 control of process parameters 183–7 design of customised assembly equipment 187–200 future trends 200–1 material handling 178–83 auger pump dispensing 516–17 austenite phase 620–1 autohesion 668 automated assembly process see assembly process automation automatic wire bonding 167–9, 176, 177, 224–5 bags, polymer 713–14 balanced production line 175 ball bonding 114, 115, 205, 206–9, 210 laser solder ball bonding 318–22, 323 ball bumping 210, 212 fluxless solder ball bumping by plasma reflow 336–9 ball grid arrays (BGA) 526, 528, 532 balloon catheters 701, 702, 713 barium-containing silicate glasses 725–6, 727, 728
WPNL2204
788
Index
batch processing 174, 536 Bay model 13 bead transition 357–8, 359 beam manipulation electron beam microwelding 420, 421, 445–8 laser welding of plastics 681–2 beam welding 73 see also under individual techniques binary systems 283, 284 biocompatibility 691 BION microstimulator 704, 780–1 bismuth 271 bond head shear test 169 bond strength see joint strength bonding atmosphere 778–9, 782–3 bonding force 185, 214–15 bonding strength see joint strength brass 476 brazing 25–50 basic metallurgical reaction 31–9 ceramic/metal bonding 767–71, 780–4 implantable devices 780–1 metallisation of ceramic surface 767–8 power amplifier packages 781–4 vacuum brazing with active filler metals 768–71, 772, 773, 774 classification of method 42–4 definition 26–31 diffusion brazing see diffusion soldering/brazing factors influencing quality of joints 31 flux brazing 47 fluxless brazing 42, 44, 47 future trends 48, 49 hermetic seals for SOFCs 732–4 materials for 41–2 nithinol 626–8 brazing filler metals see filler metals brazing fluxes 41–2, 43 brazing sheet 41 BSCCO 583 bulks 589–92 diffusion bonding 590–2 fusion welding 589–90 tapes 585–6, 613 diffusion bonding 594–608 soldering 592–3
bulk nanomaterials 741–57 fabrication processes 742–5 electrodeposition 743–4 powder metallurgy 742–3 severe plastic deformation 744–5 joining techniques 747–52 diffusion bonding using a bed of nanoparticles as interlayer 749–52 spark plasma sintering 747–9 welding processes 752, 753–4 metallic glasses 752–4 non-equilibrium state of microstructure 745–7 bulk resistance 474, 488, 493, 494 bump bonding 109–15 buoyancy force 71, 132 burn through 369, 371 C-SAM (C-mode scanning acoustic microscope) 390 calibration 161 camera-based monitoring systems 378–81, 382, 383 capacitor discharge (CD) power supply 495, 496 capillary forces 59, 435 car electronic parts 537 carbon dioxide laser 301 carbon nanotubes (CNTs) 170, 576 electron beam nanowelding 550–62 irradiation 546–50 annealing effect of irradiationinduced defects 548–50 mechanisms 546–8 resistance nanowelding 563–7 ultrasonic nanowelding 568–72, 573 carbon nanowires 559–62 cardiac rhythmic devices 692, 703–4 Carman-Kozeny equation 134 catheters balloon catheters 701, 702, 713 tip forming of catheter wire 706, 707 cathode 719, 720 CCD camera 302 cell controller 175–7 cellular microstructures 79–80, 367–8 ceramic materials, characteristics of 758–61, 762, 763
WPNL2204
Index ceramic/metal bonding 758–85 applications 780–4 brazing 767–71, 772, 773, 774, 780–4 difficulties 761–7 metallurgical inconsistency 764–6 properties mismatch 766–7 vacuum diffusion bonding 771–80, 781, 782–3 ceramic nanoparticles interlayer 752 cerium oxide 720 chain interdiffusion 669–70 chemical bonding 502 chemical etching 333, 335 chemical mechanical polishing (CMP) 657–8 chemical surface treatment 513 chemical vapour deposition (CVD) 546 chip bonding 524–5 chromium-nickel alloy 462–3, 465 clamping devices/systems electron beam microwelding 428, 429, 430 laser welding of plastics 682–3, 684 clamping force 185, 214–15 Clausius-Clapeyron equation 58 coalescence 749–50 coatings/platings 490–2 cochlear implants 692, 703 coefficient of thermal expansion (CTE) mismatch 658 ceramic/metal bonding 761, 766–7, 779 laser soldering 314–15, 320–2 seals for SOFCs glass 723–4 glass composites 730–1, 732 coherence 347 cohesion 500, 501–4 cold collapse 483 cold cracking 369 cold pressure welding 3, 5, 7, 9, 238, 248, 635 cold working 7–8 collet 182 compliant layer 735–6 compliant seals 736, 737 composite interlayer 284, 285 composite silver nanoparticle see nanoparticles
789
compression ratio lead reduction 93–4 wire reduction 105 compressive seals 734–6 computer memory hard disk suspensions 392–4 condensation-cured adhesives 520 conducting planes 585 conduction welding 127 electron beam microwelding 433 laser microwelding 353–7 conduction to keyhole transition welding 354, 357–8, 359 model predictions 364–7 opaque materials 353–6 stainless steel sheets 395–6 transparent materials 356–7 conductive heat transfer 123, 124–8 boundary conditions 124–5 definition of heat source model 127 example 127–8, 129, 130, 131 governing equations 124–6 solution 125–6 constriction resistance 147, 488, 493 contact adhesives 505 contact angle 32–3, 503–4, 764 contact lenses 678 contact radius 150, 151, 152, 155 contact resistance 474, 488, 493 effects of surface conditions 488–92 modelling 146–8 contact wave 636 contaminant displacement/interatomic bonding 7–14, 22 extension of contacting surfaces 7–8, 9, 10, 11 realignment of grain structures 6, 14, 15 separation of contaminated areas 8–14 contaminant layer 6, 327, 328 thermal dissolution 6, 14–18, 22 continuous-wave lasers 186 control 159–73 adaptive 381–4 defining and methods 161–2 process control see process control control voltage 419–20 convective heat transfer 123, 128–42 boundary conditions 135–7
WPNL2204
790
Index
driving forces of fluid flow in the weld pool 131–2 examples 138–42, 143, 144 governing equations 132–7 solutions 137 conventional diffusion bonding (CDB) BSCCO tapes 594–602 design of joint structure and bonding technology 594–5, 596, 597, 598 effects of bonding parameters on joint quality 595–601 with an interlayer 601–2 without an interlayer 594–601 see also diffusion bonding cooling rate 368–9 copper 461, 462, 476 bonding using composite silver nanoparticle 255–9 comparison of copper-to-copper joints with lead-rich solder joints 265 effects of bonding conditions on joint strength 259–60 microstructure and fracture morphology of joints 260–4 laser microwelding 355–6 nanoparticles as interlayer for diffusion bonding 750–1 resistance nanowelding of copper-filled CNTs 566–7 silver-copper eutectic braze alloy 627, 628 wire bonding with copper wire 226–7 copper oxide planes 584, 585 cored wire 700 covalent bonds 761 cracks 79, 319–22, 369 creep 320–2, 323 critical current ratio (CCR) conventional diffusion bonding of BSCCO tapes 595, 598–600, 601–2 direct diffusion bonding of BSCCO tapes 603–6 critical transition temperature 583, 584, 585 cross-wire welding
electron beam microwelding 463, 466 resistance microwelding 479–81, 482, 483 electrode force 487–8 plating 491–2 crystalline melting point 666–7 crystallisation of glasses 724–9 crystallographic matchup 6, 14, 15 cuprates 584–5 curing of adhesives 519–21 classification of adhesives by curing mode 505, 506–9 curing hastening processes 521 current anodic wafer bonding 642, 651–2 welding current in resistance microwelding 482–4, 485 current density 150, 153 customised assembly process automation 201 case study 187–200 cutting, laser 710–12 cyanoacrylates 506 cyclic extrusion-compression (CEC) 744, 745 dam and fill coating 528, 530–1 Darcy resistance 134 decoalescence 62–3, 64 defects due to adhesive bonding 522 fusion microwelding 78–9 in-process monitoring of soldering 166–7 irradiation-induced in CNTs 546–50 annealing effect 548–50 weld defects in laser microwelding 368–77 development process defects 369–70, 371 manufacturing process defects 370–7, 378 materials defects 368–9 defibrillators 703–4 deflection coils 420, 421 deformation diffusion bonding 239–40 overcoming surface oxide problems 241–4
WPNL2204
Index mechanism of solid-state bonding 6, 7–14, 22 mechanism for wire bonding 217 deformation monitoring 167–9 dendritic microstructures 79–80, 367–8 density 71 depth of focus 348 destructive testing 522 development process defects 369–70, 371 die attach bonder 176, 177 dielectric constant 227 differential scanning calorimetry (DSC) 287–8, 511, 621, 622 diffusion hypothesis for solid-state bonding 241 theories of welding plastics 667–8 theory of adhesion 502 diffusion bonding 6, 234–49 advantages 236–7 approaches to overcome surface oxide problem 241–5 bonding parameters and apparatus 238–9, 240 BSCCO bulks 590–2 BSCCO tapes 594–608 bonding parameters and joint quality 595–601, 603–6, 607 conventional with an interlayer 601–2 conventional without an interlayer 594–601 direct with precursors 606–8 direct without an interlayer 602–6, 607 bulk nanostructured materials using a bed of nanoparticles as interlayer 749–52 ceramics 764 compared with other joining methods 236–8 definition 234–5 limitations 237–8 removal of surface oxides using liquid gallium 245–7 theoretical aspects and modelling 239–41, 242 YBCO bulks 608–9 diffusion soldering/brazing 269–98 applications 270–2
791
evaluating joint properties 285–9 measuring diffusion brazing process kinetics 286–8 mechanical testing 289 metallography 286–7 future trends 296 historical development 269–70 modelling 288, 289–96 process 273–8 advantages and disadvantages 274–5 equipment 275–6 procedures 275 process parameters 276–8 process mechanics 278–85 critical stages 282 isothermal solidification in multicomponent systems 282–4 stages 273, 277, 278–82 temperature-gradient TLP bonding 284–5, 286 wide-gap process variation 284, 285 diode lasers 301, 323, 680 direct copper bonding (DCB) 768 direct diffusion bonding (DDB) 602–8 effects of bonding parameters on joint properties 603–6, 607 method 602–3, 604 with precursors 606–8 without an interlayer 602–6, 607 see also diffusion bonding direct laser welding 678, 687 direct resistance nanowelding 563–6 direct wafer bonding 635–41 background 635 mechanism 638–41 silicon direct bonding process 635–8, 639 dislocation hypothesis of solid-state bonding 240–1 dispensing techniques 514–18 mass dispensing 514–15 selective dispensing 514, 515–18 displacement rates 95–6 dissimilar materials laser welding 698–9, 700 wafer bonding 655–8 dissolution
WPNL2204
792
Index
brazing and soldering 35–9 factors influencing 36–8 IMC formation 38–9 kinetics 35–6 dissolution and widening stage of diffusion soldering/brazing 273, 277, 279–80, 294–5 laser soldering 309–12 thermal dissolution 6, 14–18, 22 divergence 347 drilling 369, 370 medical devices 710–12 vs welding 60–2 droplet welding, laser 364, 397–400 droplets, impinging 64–5, 132 dye penetrants 390, 391 dynamic mechanical analysis (DMA) 511 dynamic method for contact angle 503–4 dynamic mixing system 518–19 dynamic resistance 492–4 eddy currents 390, 391 electrical corona discharge 513 electrical signals 162 electrical-thermal-mechanical analysis 115–17, 118, 145–6, 149–54, 155 electrode force 152–4, 155 resistance microwelding 476, 477, 486–8 electrode sticking 477, 492 electrodeposition 743–4 electrodes resistance microwelding 495–7 SOFCs 719, 720, 721 electrolyte 719, 720 electrolyte supported SOFCs 718, 719 electromagnetic force 55, 69–70, 72, 132, 135 electromagnetic heat generation 672–3, 674 electron beam-assisted deposition (EBAD) 454 electron beam microwelding 73, 418–72 apparatus 420, 421 bead transition 357–8, 359 examples and applications 459–71 metals 459–68, 469 plastics 468–71
process 445–59 addition of filler material 453–9, 460, 461 advanced process variations 450–3 basic beam process variations 445–8 parameter choices 448–50 sequence 445, 447 technology 419–45 beam characterisation methods 424–6 beam manipulation 420, 421, 445–8 clamping and guiding devices 428, 429, 430 conversion to energy vs position 440–2 conversion of an SEM into a micro electron beam welder 422–4, 425 electron beam generation 419–20 energy transfer with the workpice and penetration effect 432–5 fast beam controls and multi-beam technique 426–7 modelling beam/workpiece interaction 435–40 thermal effects in microscale parts 442–5, 446 waveform generator and teach-in programming 428–32, 433 welding machine peripherals 420–2 electron beam nanowelding 546–62 direct joining by using SEM 552–3 direct joining by using TEM 550–2 indirect joining by SEM and TEM 553–9 ion beam nanowelding 559–62 electron hypothesis of solid-state bonding 240 electron microscopy 78 see also scanning electron microscope (SEM); transmission electron microscope (TEM) electronics assembly, soldering for 45–7 electrostatic attraction 502 endoscopes 536, 537 energy
WPNL2204
Index conservation in modelling of fusion microwelding 133–4, 135 electron beam microwelding beam penetration and energy distribution calculations 435–40 electron energy absorbed and accelerating voltage 448–50 electron energy contours 440–2 energy transfer 432–5 free energy change see free energy change threshold energy 547–8 energy contours 440–2 energy density 683–4 enthalpy 134–5 epoxy curing oven 176, 177 epoxy resins 507–8, 537 equal-channel angular pressing (EQAP) 744, 745 equivalent stress 92, 98, 99–100 distribution in FCB 112, 113 erosion 36–8 etch holes, patches over 455, 456 etch-stop layer 634 European Union (EU) Restriction of Hazardous Substances (RoHS) directive 28, 166 eutectic soldering 184 eutectic solidification structure 289, 290–1 evaporation 55, 57–9, 72, 132 evaporation pressure 358–60 extension of contacting surfaces 7–8, 9, 10, 11 external heat source 672, 674 external load cell 185 extrinsic applied mechanical force 55–7, 72 extrusion pressure 13, 14 far-field welding 675 Faraday cup 424–6 fast beam controls 426–7 fatigue 320–2, 323 feedback methodologies 159–73 applications 163–9 in-process monitoring 164–7 process control 167–9 definitions and methods 160–2
793
future trends 169–71 signals from joining processes 162–3 see also control; monitoring; sensing/ sensors fibre lasers 680 SMFL 403–5 fibre optics see optical fibres fibre pigtailed TO-can laser diode package 187–200 Fick’s laws/equations 273, 292 filler material (electron beam microwelding) 453–9, 460, 461 filler metals 26–8, 29, 41 vacuum brazing of ceramic/metal with active filler metals 768–71, 772, 773, 774 fillet size 33–5 film hypothesis of solid-state bonding 240 film resistance 147, 488, 489, 493 fine pitch ball bonding 229 finite difference modelling 145 finite element modelling (FEM) 92–3, 122 conductive heat transfer 125–6 laser soldering 320–2, 323 polymer welding 670–1 resistance welding 149–54, 155 solid-state bonding 92–3 gold bump bonding 109–15 inner lead bonding 93–100 lap resistance welding 115–17, 118 wire bonding 101–7 thermal effects in electron beam microwelding 442–5, 446 thin-die pickup process 181 fixed clamp system 682, 683 fixturing problems in micro fusion welding 76–7 wire bonding 224–6 flame-melting method 589–90 flash-butt welding 8, 11 flashing acceleration 8, 11 flat tool, wire bonding using 102, 104, 105, 106, 107, 109 flip chip assembly (FCA) 526–7 underfill 528, 532 flip-chip bonding (FCB) 91, 109–15, 206
WPNL2204
794
Index
flow soldering 45, 46 flow velocity 374, 375 fluid flow modelling GTAW 138–40, 141 modelling laser microwelding 140–2, 143, 144 molten metal flow 361–3 fluorescent dye penetrants 390, 391 flux brazing 47 flux-cored solder wire 46–7 flux residues 328 fluxes brazing 41–2, 43 role of flux in soldering 327 soldering 40–1 fluxless brazing 42, 44, 47 fluxless soldering 40, 327–44 demand for 328 laser soldering 312–14, 323 processes 328–42 acid vapour fluxless soldering 329, 330 gaseous reduction 328–9 PADS 329–30, 331, 342 plasma fluxless soldering 330–9, 342 solder ball bumping by plasma reflow 336–9 ultrasonic 339–41, 342 focused ion beam (FIB) 453–9, 460, 461 focusing coil 420, 421 folding 103–5 force clamping force 185, 214–15 control of applied force 185 electrode force see electrode force forces acting on the pool in fusion microwelding 54–72, 131–2 for separation of two polymer surfaces 668 ultrasonic welding of plastics 677 forced vortex 65–6 forge welding 3–4, 238 formaldehyde resins 509 formic acid vapour soldering 329, 330 fracture surfaces anodic wafer bonding 654–5, 657 bonding with composite silver nanoparticles 256–7, 261–4
direct wafer bonding 637, 638 laser welding of titanium 693, 699 free-air balls 207 copper wire 226 insulated bonding wires 227, 228 free energy change bonding using nanoparticles 257–8, 259 laser soldering 307–8 frequency-based monitoring 381, 382, 383 fretting wear 219 friction welding 238, 248 combined with diffusion welding 245 stir welding 625, 674, 689, 752 frictional slip see slip fusion microwelding 51–90 arc seam welding 84 arc spot welding 69, 74, 84–6, 706 BSCCO bulks 589–90 forces acting on the pool 54–72, 131–2 aerodynamic forces 55, 68–9, 72 electromagnetic force 55, 69–70, 72, 132, 135 evaporation 55, 57–9, 72, 132 gravity 55, 71, 72 momentum/inertial forces 55, 56–7, 63–8, 72 pool stability 60–3, 64 surface tension 55, 59–60, 72, 131–2 thermal expansion/contraction vs extrinsic applied force 55–7, 72 viscosity 55, 68, 72 laser seam welding 83 laser spot welding 81–2, 192–5 modelling see modelling problems in applying 76–80 process variations 73–6 resistance welding see resistance microwelding thermal aspects 52–4 gallium, liquid 245–7 gallium arsenide 658 gas tungsten arc welding (GTAW) 138–40, 141 gaseous reduction 187, 328–9
WPNL2204
Index Gaussian laser beam 304, 347, 348 gears welds to pins 463, 466 welds to shafts 402, 463–4, 467 geometrical (optical) signals 162, 163 Gibbs’ phase rule 284 Gibbs-Thomson effect 749 glass 634, 658 metallic glasses 752–4 seals for SOFCs 723–32 conventional glasses 723–9 glass composites 730–2 invert glasses 729–30 ultrashort pulse fusion welding 405–8, 409 wafer bonding anodic bonding of silicon to glass 641–8 glass-to-glass anodic bonding via silicon layer 651–5, 656, 657 intermediate wafer bonding via glass layer 650–1 glass ceramics 723–9 glass composites 730–2 glass transition temperature 511, 666–7 glass seals for SOFCs 723–4 glob topping 527–8, 529 gold bonding using composite silver nanoparticle 255–9 fluxless soldering using tin-gold solder 341–2 gold-tin IMCs 315–17, 318–20, 321–2 laser nanowelding of gold nanoparticles 573–5 model of gold bump bonding 109–15 plating and resistance microwelding 490–2 grain boundary diffusion 281 grain boundary grooving 281 grain boundary sliding 320, 321 grain structures, realignment of 6, 14, 15 graphitic carbon 557–9 gravity 55, 71, 72 gross sliding 217, 218, 219, 220, 221 guidewire joining 701 guiding devices 428, 429, 430 hard disk suspensions 392–4
795
hardness metal oxides 7, 8, 10 testing 389 health and safety 524 heat 3–4 diffusion soldering/brazing heating rate 278 heating stage 273, 277, 278–9 electron beam microwelding heat input 448–9 surface vs volume heat source 440–5, 446 fusion microwelding heat input 52–4 modelling 122–42 generation for plastics microwelding 671–3, 674 electromagnetism 672–3, 674 external heat source 672, 674 mechanical movement 671–2, 674 laser soldering heat loss 302 thermal process 303–6 thermal characteristics of composite silver nanoparticles 251–5 thermal deformation 394–5 thermal dissolution 6, 14–18, 22 thermal effects and plastic materials 665–7 thermal expansion/contraction 55–7, 72 thermal modelling see thermal modelling thermal signals 162–3 thermal stability 745–7 heat–cured adhesives 520 heat exchangers 26, 28, 29, 41 heat sink 466–8, 469 heat source models 127 heat transfer 123 conductive 123, 124–8, 129, 130, 131 convective 123, 128–42, 143, 144 hermetic cans 703–4 hermetic sealing implantable devices 708–10, 711 SOFCs 718, 722–37 compliant seals 736, 737 compressive seals 734–6 rigid seals 722–34
WPNL2204
796
Index
hexacelsian glass ceramics 725–6 high frequency (HF) inverter power supply 495, 496 high-pressure torsion (HPT) 744, 745 high temperature superconductors (HTS) 583–619 BSCCO bulks 589–92 BSCCO tapes 585–6, 613 diffusion bonding 594–608 soldering 592–3 future trends 613–14 needs for joining 588–9 processing technologies 585–8 superconducting materials 583–5 YBCO bulks 586–8 joining 608–13, 613–14 holding time 37–8 diffusion bonding of BSCCO tapes 595, 598, 599, 603, 604, 605 diffusion soldering/brazing 278 joint strength in ceramic/metal bonding brazing 770–1, 772 vacuum diffusion bonding 775–6, 782–3 strength of copper-to-copper joints using composite silver nanoparticle 259–60 ultrasonic welding of plastics 677 homogeneous heating 41, 48, 49 homogenisation 273–4, 277, 278, 281–2 homologous temperature 745, 746 hot bar welding 674, 689 hot collapse 483 hot cracking 369 hot melt adhesives 504–5, 524, 537 MEMS and MOEMS applications 537–40, 451 reactive 505, 506 hot pressing 747, 749 humping 374–5 hybrid heat transfer 123 hydrogen gaseous reduction 328–9 plasma fluxless soldering 330–1 cleaning and soldering using argon and hydrogen gas mixture 331–6
fluxless solder ball bumping by plasma reflow 336–9 hydrogen bonding 639, 640, 648 hydrostatic pressure 59, 435 hydroxyl groups 648 ice flow theory 333 ‘ideal’ laser microwelding system 385–7 impact force 214 impinging droplets 64–5, 132 implantable devices see medical devices and components impulse welding 674, 689 in-line process operation 174 assembly equipment for 175–8 in-process monitoring 161–2, 164–7 laser microwelding 377–81, 382, 383 indirect resistance nanowelding 566–7 indium 271 indium oxide 591 induction heating 521 induction welding 674, 689 inertia welding 752 inertial forces 55, 56–7, 63–8, 72 infra-red brazing 627–8 infra-red process control 302–3 inner lead bonding 93–100 inspection see visual inspection insulated bonding wires 227–9 intensity of laser beam 348 interatomic bonding see contaminant displacement/interatomic bonding interchangeability of tool sets 180 interconnect 721 interface breakdown of interfacial structure 18–21, 22–3 interfacial extension and modelling solid-state bonding 95, 97, 99, 106–7, 109 too small an interface area 714–15 viscoplastic deformation model and interfacial deformation 91–3 interface liquid phase assisted joining 612–13 interfacial heating 75, 216–17 intergranular cracking 319–20 interlayer
WPNL2204
Index bed of nanoparticles as interlayer for diffusion bonding of bulk nanomaterials 749–52 diffusion bonding of BSCCO tapes with conventional 601–2 direct 606–8 diffusion soldering/brazing 275 composition 276–8 thickness 276 effect on vacuum diffusion bonding of ceramic/metal 779–80, 782–3 intermediate layer wafer bonding 648–55, 657–8 glass-to-glass anodic bonding via silicon layer 651–5, 656, 657 glass layer 650–1 silicates 649 silicides 648–9 solder 649–50 intermetallic compounds (IMCs) 38–9, 49, 369 diffusion soldering/brazing 270, 288, 289 laser soldering 314–17, 318–20, 321–2 interstitial migration 549–50 inverse Bremsstrahlung effect 352 invert glasses 729–30 ion-assisted etching 333, 335 ion beam-assisted deposition (IBAD) 453–9, 460, 461, 588 ion beam nanowelding 559–62 ionic bonds 761 iron 16 nanostructured 745–7 iron oxide 16 irradiation of CNTs 546–8 annealing effect of irradiation-induced defects 548–50 isothermal solidification 269, 289, 290–1 measuring diffusion brazing kinetics 286–8 modelling 291–4 in multi-component systems 282–4 stage in diffusion soldering/brazing 273–4, 277, 280–1, 282 isothermal solidification rate constant 293, 295
797
isotropic conductive adhesives (ICAs) 510, 529–32, 533, 534 isotropically conductive film (ICF) 529–32 jet dispensing 517–18 joint evaluation techniques 387–91 joint strength bonding with composite silver nanoparticles 255–7, 258 effect of bonding conditions 259–60, 265 silicon chip bonding 266–7 ceramic/metal bonding brazing 770–1, 772, 773, 774 vacuum diffusion bonding 773–80, 782–3 development and welding plastics 667–9 electron beam microwelding 459, 460, 461 fluxless solder ball bumping by plasma reflow 337, 338 resistance microwelding 487–8 solid-state bonding 10–13 and annealing temperature 19–20 ultrasonic bonding 340, 341 wafer bonding anodic 644, 646, 653–4, 656 direct 636–7 joule heating 75, 142–4 keyhole welding 127, 353, 354, 358–63 conduction to keyhole transition welding 354, 357–8, 359 keyholing 57–9, 363–4 electron beam microwelding 433–5 kinematic design 180 kitchen equipment 395–6 lack of fusion 369, 370 lap welding modelling lap resistance welding 115–17, 118 non-bridged beads in laser lap welding 375, 376–7, 378 lapped-window length 605, 606 lapped-window step number 594, 595, 596, 600, 606
WPNL2204
798
Index
large-scale resistance spot welding (LSRSW) 121, 150–2, 153, 154, 164, 476–7 laser-assisted chemical vapour deposition 364 laser-assisted ultrasonic wire bonding 364 laser cutting 710–12 laser drilling 369, 370, 710–12 laser droplet welding (LDW) 364, 397–400 laser microwelding 61–2, 73, 345–417 active alignment laser welder for fibre pigtailed TO-can laser diode package 187–200 applications 391–6 examples 392–6, 397 industrial base of microwelding lasers 391–2 evaluation of microweld joints 387–91 fundamentals 347–63 classification of welding mode 353–63 laser beam propagation and focusing 347–9 laser-matter interaction 349–51 laser-plume interaction 351–3 future trends 408 history 345–6 medical devices 708, 709 hermetic sealing 710, 711 tube-to-pin assembly 713 tube-to-tube assembly 712, 713 modelling see modelling novel technologies 396–408 laser droplet welding 364, 397–400 laser spike welding 73, 400–1 SHADOW welding 54, 83, 401–3, 404 ultrashort pulse fusion welding of glass 405–8, 409 plastics 396, 397, 678–88, 689 advantages of transmission laser welding arc 686 equipment and variations 679–83 examples 687–8 process description 678–9 welding parameters 683–4 seam welding 83 SMAs 625, 628–30
spot welding 81–2, 192–5 technologies 377–87 adaptive control 381–4 ‘ideal’ system 385–7 in-process monitoring 377–81 weld defects 368–77 development process defects 369–70, 371 manufacturing process defects 370–7, 378 materials defects 368–9 laser nanowelding 572–5 laser pulse, timing 183–4 laser radiation 186–7 laser soldering 299–326 fluxless 312–14, 323 fundamentals 301–12 dissolution of materials being soldered 309–12 formation of solder joint 306–9 process control 300, 302–3 thermal process 303–6 future trends 322–4 overview 299–300 reliability of solder joint 300, 314–22, 323 laser reflow soldering 315–18 laser solder ball bonding 318–22, 323 layer scanning mode for electron beam microwelding 447–8 layout of processing modules 178–9 lead-free diffusion soldering 270 lead-free solders 28–31, 38, 39, 48, 49, 270 lead-rich islands 317 lead-rich solder joints 265 leak checking 388 LIGA materials 78–9, 368–9 light-based in-process monitoring 378–81, 382, 383 light-cured adhesives 520 line balancing 175 liquid gallium 245–7 liquid phase-assisted diffusion bonding process 245 liquid width 286–8 liquidus temperature 26–8, 28, 30 lithium 244–5
WPNL2204
Index lithium ion batteries 392, 393 load frame 734–5 local beam pattern motion 431, 432 localised melting region 123 Lorentz force 69 Lorentz force vortex 65–6 low-k materials 227 low temperature wafer anodic bonding 642–6 M parameter 93, 99–100, 106 machine vision algorithm 163 magnesium 244–5 magnesium aluminosilicate glasses 728 magnetic particle techniques 390, 391 magnetic pulse welding 712 manual soldering 45, 46–7 manufacturing process defects 370–7, 378 Marangoni effect 59–60, 128 Marangoni stress 131–2, 136 martensite phase 620–1 mass balance 133, 273, 283–4, 292–3 mass dispensing techniques 514–15 material-related forces 54, 55–68, 71 material safety data sheet (MSDS) 524 materials defects and laser microwelding 368–9 implantable devices 692–700 problems in fusion micro welding 78–80 materials handling 76–7 assembly processes 178–83 mechanical amplifier 186 mechanical interlocking 504 mechanical movement, heat generated by 671–2, 674 mechanical signals 162 mechanical surface treatment 513 mechanical testing see testing medical devices and components 691–717, 780–1 adhesive bonding 536–7, 712–13 future trends 715–16 joint design and process selection 705–14 materials challenges 692–700 SMAs 623–4, 625–6 testing and verification 714–15
799
types of devices 700–5 ultrasonic welding of plastics 677 melt-pump-up method 591–2 melt-textured growth 587–8 melting efficiency 366, 367 melting point 271 melting point depressant (MPD) 275 Mercedes-Benz car remote entry keyless module 396, 397 metal/ceramic bonding see ceramic/metal bonding metallic gaskets 735 metallic glasses 752–4 metallic nanoparticles composite silver nanoparticle see nanoparticles as interlayer for diffusion bonding of bulk nanomaterials 749–51 metallisation of ceramic surface 767–8 metallography diffusion brazing 288–9 evaluation of laser microwelded joints 388 metallostatic head (gravity) 55, 71, 72 metallurgical problems 78–80 metals cold welding 633–4, 635 conduction laser welding 354–6 electron beam microwelding 459–68, 469 material properties 696–7 metallic brazing and hermetic seals for SOFCs 732–4 microwelding of metal foil by SMFL 403–5 molten metal flow 361–3 weldability matrix 695 see also under individual metals mica-based compressive seals 735–6 micro-to-nano adapted electron beam welding (MAEBW) 552–3 microelectromechanical systems (MEMS) 323, 678 hot melt adhesives 537–40, 541 packaging 633 SMAs 626 thin film SMA 624 microelectronics adhesive bonding applications 524–35
WPNL2204
800
Index
wire bonding 164–6, 205–33 microfluidic devices 688 microhardness testing 389 micromachining stations 386 micromanipulation 389 microplastic deformation see plastic deformation microslip see slip microslip annulus 218 microwave-cured adhesive 521 microwave-melting joining method 590 microwave welding 674, 689 milling 742–3 mixing of adhesives 518–19 modal analysis 186 modelling diffusion bonding 241, 242 diffusion soldering/brazing 288, 289–96 analytical solution for isothermal solidification 291–4 multi-component systems 295–6 numerical modelling techniques 294–6 electron beam microwelding 435–45 beam penetration and energy distribution calculations 435–40 surface vs volume heat source 440–5, 446 fusion microwelding 121–58 features of thermal process 122–3 future trends 154–6 modelling of conductive heat transfer 124–8, 129, 130, 131 modelling of convective heat transfer 128–42, 143, 144 laser microwelding 363–8 conduction model predictions 364–7 conductive heat flow 127–8, 129, 130, 131 convective heat flow 140–2, 143, 144 levels of complexity 363–4, 365 metallurgical consequences 367–8 polymer welding 669–71 problems at small scales 78 resistance microwelding 115–17, 118, 142–54, 156
analysis procedure 145–6 comparison with large-scale resistance welding 150–2, 153, 154 contact resistance 146–8 examples 150–4, 155 mesh structure and boundary conditions 149–50 principles and features 142–5 solid-state bonding 91–120 gold bump bonding 109–15 inner lead bonding in TAB 93–100 interfacial micro friction slip 101, 110–15 lap resistance welding 115–17, 118 numerical simulations of FCB 112–15 thermocompression bonding 93–107 thermosonic bonding 107–15 viscoplastic deformation model and interfacial deformation 91–3 wire bonding 100–7, 108, 109 thermal process of laser soldering 304–6 MOEMS 537–40, 541 Mohamed and Washburn model 10, 12 moisture-cured adhesives 520 moisture resistance 511–12 momentum conservation 133, 134 forces in fusion microwelding 55, 56–7, 63–8, 72 monitoring 159–73 defining 161 in-process monitoring 161–2, 164–7 laser microwelding 377–81, 382, 383 monocelsian glass ceramics 725–6 Monte Carlo modelling 435–40 motion control 180, 385 motors, small 392, 393 moulding machine 176, 177–8 moving clamp system 682–3 moving parts 67–8 multi-beam electron beam microwelding 426–7, 450–3 multi-component systems 282–4
WPNL2204
Index modelling in 295–6 multiple location welding 394–5 multi-mode optical fibre (MMF) 190–1, 199 multiple scan electron beam microwelding 445–7 multi-row bonding 228 multiwalled nanotubes (MWNTs) 546, 547, 548 electron beam nanowelding 554–6 ion beam nanowelding 559–62 nanocrystals 742 nanohardness testing 389 nanojoining 545–79 electron beam nanowelding 546–62 future trends 576 laser nanowelding 572–5 resistance nanowelding 563–7 ultrasonic nanowelding 568–72, 573 nanoparticles 545 additives for adhesives 510 bed of nanoparticles as interlayer for diffusion bonding of bulk nanomaterials 749–52 composite silver nanoparticle 250–68 bonding mechanism 257–9 bonding of silicon chip 265–7 bonding strength and fracture morphology 255–7, 258 comparison with lead-rich solder joints 265 effects of bonding conditions on joint strength 259–60, 265 future trends 267 microstructure and fracture morphology 260–4 structure 251, 252 thermal characteristics 251–5 laser nanowelding 572–5 nanosensors 170–1 nanostructured powder 742–3 nanotubes 545–6 see also carbon nanotubes (CNTs) nanowires 545 Nd:YAG laser 301, 680 sales trends 391–2 near-diffraction limited lasers 346, 385, 386
801
near-field welding 675 needle-based selective dispensing 515–16 neurostimulators 692, 703–4 Newton’s second law of motion 56, 64 Ni-YSZ cermet 721 nickel added to filler in brazing nithinol 627, 628 bonding using composite silver nanoparticle 255–9 electron beam microwelding 449, 462, 464 modelling 436–45 in glass composites 731–2 nanostructured 745–7 nickel-chromium alloy in electron beam microwelding 462–3, 465 nithinol SMAs 620, 621–2 applications 623–4, 626 brazing 626–8 stents 625–6 welding 625–6 no clean process 40–1, 49 non-bridged beads 375, 376–7, 378 non-conductive adhesives 528, 532 non-corrosive flux 42, 44 non-destructive testing 522 nonlinear absorption 356, 357 nucleating agents 728 nugget growth curves 152–3, 155 nugget shape and size 150–1, 154 numerical modelling diffusion soldering/brazing 294–6 lap resistance welding 115–17, 118 solid-state bonding 97–100, 105–7, 108, 109, 112–15 opaque materials 353–6 optical coupling efficiency 189–90 optical fibres 186 electron beam microwelding 453 laser beam delivery via 349 laser microwelding of optical fibre connectors 394–5 optical linear encoders 180 optical signals (geometrical signals) 162, 163 optoelectronics 323
WPNL2204
802
Index
interconnections and packaging using adhesive bonding 536 orientation, in assembly line 179 orthodontic treatment 626 oscillating tangential displacement 213, 221–2 overwelding 715 oxidation, excessive 369, 371 oxide surface layer 6, 327, 328 extension of contacting surfaces 7–8, 9, 10, 11 hardness of metal oxides 7, 8 overcoming problem in solid-state diffusion bonding 241–5 removal using liquid gallium 245–7 resistance microwelding 489–90 thermal dissolution 6, 14–18, 22 oxygen-silicon reaction, and anodic wafer bonding 647–8 pacemakers 703–4 packaging adhesives 519 MEMS 633 parallel-gap microwelding 481–2, 484 part fit-up 77 partially stabilised zirconia (PSZ) 759 passive structural control 623 patches over etch holes 455, 456 Peclet number 129, 142, 144 peel testing 714, 715 Peltier effect 75–6 penetration distance and laser microwelding 350–1 electron beam microwelding 432–5 beam penetration and energy distribution calculations 435–40 energy contours 440–2 percussive arc welding 69, 74, 84–6, 706 perovskites 720 phase transformations glass seals 724–6 SMAs 620–2 zirconia 759 phosphorus 37–8 physical sputtering 333, 335
pick and place equipment 179, 181–3 piezoelectric sensors 162 piezoresistors 162, 165, 166 pin transfer 514–15 pinch-off dies 8, 11 pins gear-to-pin welds 463, 466 pin/rivet geometry 464, 467 polymer tube to solid pin joints for implantable devices 713 wire-to-pin joints 707–8, 709 piston pump dispensing 517 placement accurate 184–5 thermoplastic adhesive tape placement process 181, 182–3 planar interface 79–80 plasma arc welding (PAW) 74 plasma assisted dry soldering (PADS) 329–30, 331, 342 plasma fluxless soldering 330–9, 342 argon or hydrogen plasma 330–1 cleaning and soldering using argon and hydrogen gas mixture 331–6 fluxless solder ball bumping by plasma reflow 336–9 plasma seam welding 710 plastic deformation 633, 635 diffusion bonding 239–40 enhancing microplastic deformation of surface asperities 243–4 imposing substantial plastic deformation 241–3 severe (SPD) 744–5 plastics (polymers) 665–90 electron beam microwelding 468–71 future trends 689 implantable devices 712–14 bags and enclosures 713–14 tube to solid pin assembly 713 tube-to-tube joining 712–13 laser microwelding 396, 397, 678–88, 689 theory of welding plastics 665–71 effect of welding parameters 669 modelling of plastics welding 669–71
WPNL2204
Index strength development at a joint 667–9 thermal effects 665–7 ultrasonic welding 673–8, 689, 714 weldability matrix 694 welding processes 671–3, 674 electromagnetism 672–3, 674 external heat source 672, 674 heat generated by mechanical movement 671–2, 674 plastisol adhesives 505 platings/coatings 490–2 platinum platinum deposit from FIB in electron beam microwelding 454–9, 460, 461 platinum-titanium alloys 698–9, 700 plume, laser-induced 351–3, 363 polished polysilicon layers 657–8 polyacetal (POM) 468–71 polyaddition 505, 506, 507–9 polycarbonate (PC) dental prosthesis 688 foils 687 polycondensation 505, 506, 509 polycrystalline silicon 464–6, 468 dies 456–9, 460, 461 polyethylene (PE) electron beam microwelding 468–71 film 687 polyethylene terephthalate glycol (PETG) 688 polymerisation 505, 506–7 polymers see plastics (polymers) polymethylmethacrylate (PMMA) 468–71, 687–8 polyurethanes 508–9 pool sloshing 66–7 pool stability 60–3 porosity 79, 369, 372–4 position difference of laser spots 196–7 positioning unit 423–4, 425 post-mould curing oven 176, 178 post-weld shift (PWS) 192, 194, 195–7 powder-in-tube (PIT) process 585–6 powder metallurgy 742–3 power power supplies for resistance microwelding 494–5, 496
803
requirements for laser microwelding 385–6 ultrasonic welding of plastics 676 power amplifier packages 781–4 power density value 367 pre-applied hot melt adhesives 537–40, 541 precision (repeatability) 160–1 fusion microwelding 77 welding of plastics 685 precursors, direct diffusion bonding with 606–8 pressure 185 ceramic/metal bonding 776–7, 778, 782–3 diffusion bonding of BSCCO tapes 595, 598–600, 603–5, 607 diffusion soldering/brazing 278 and ion bombardment energy 333, 335 modelling convective heat transfer 136 saturated vapour pressure 58 strength of copper-to-copper joints using composite silver nanoparticle 259–60 pressure-assisted powder metallurgy processes 743 pressure bellows 710 pressure-sensitive adhesives 505 primers 513 printed circuit boards (PCBs) 525 mounting of SMDs 525–6, 527 process control 167–9 assembly process automation 183–7 fusion microwelding 77 laser soldering 300, 302–3 resistance microwelding 492–4 process-related forces 54, 55, 55–7, 68–70 projection joints 675 pseudoelasticity 621–2, 623, 629 pull testing 222–3, 388–9 medical implantable devices 714 pulse mode lasers 186 pulsed arc welding 69 pulsed laser welding microwelding 361–3 ultrashort pulse fusion welding of glass 405–8, 409
WPNL2204
804
Index
underfilled beads 370–2 pumps 703 pyrometer 302 quality control adhesive bonding 521–2 wire bonding 222–4 radiance of laser beam 348 radiation 123 radiation-cured adhesives 506–7, 520 radio frequency (RF) curing adhesives 521 radio frequency (RF) dielectric welding 714 radio frequency identification device (RFID) 535 radioactive seed implants 704–5 reaction compounds 777–8 reaction layer thickness 773, 775 reaction mode, classification of adhesives by 504–6 reactive brazing 732–4 reactive hot melt adhesives 505, 506 real-time control methods 161–2 recoil force, vapour 58–9, 61–2, 363 recovery 18–19 recrystallisation 19–20, 21, 240 reducing gases 187, 328–9 reference range 164 reflection 360–1 reflectivity 350 reflow soldering 45, 46 laser reflow soldering 315–18 refraction 352, 353 reliability laser soldered joints 300, 314–22, 323 medical devices 691 monitoring 169–70 remote entry keyless module 396, 397 repeatability see precision (repeatability) reptation theory 668–9 residual stresses 766–7 resistance butt welding 14, 15 resistance microwelding 75–6, 86–7, 121, 473–99 compared with large-scale resistance welding 150–2, 153, 154, 476–7
dynamic resistance and process control 492–4 equipment 494–7 electrodes 495–7 power supplies 494–5, 496 fundamentals 473–7 future trends 497 in-process monitoring 164 medical devices 707–8, 709–10 modelling see modelling process variations 477–82 crossed-wire 479–81, 482, 483, 487–8, 491–2 sheet-to-sheet spot welding 478–9, 486–7 wire-to-sheet 481–2, 484 resistance spot microwelding 121, 142–5, 149–54, 155, 473–7 SMAs 625 surface conditions 488–92 oxide film 489–90 plating 490–2 roughness 488–9 thermal expansion/contraction vs extrinsic applied mechanical force 55–7 welding parameters 482–8 electrode force 486–8 weld time 484–6 welding current 482–4, 485 resistance nanowelding 563–7 direct 563–6 indirect 566–7 resistance project welding 3, 5 resistance seam welding 3, 5 for hermetic sealing 709–10 resistance spot welding (large-scale) 121, 150–2, 153, 154, 164, 476–7 resistive temperature detectors (RTDs) 163 Restriction of Hazardous Substances (RoHS) directive 28, 166 Reynolds number 68, 455–6 Richardson’s law 419 rigid seals 722–34 air brazes 734 conventional glasses 723–9 glass composites 730–2 invert glasses 729–30
WPNL2204
Index metallic brazes 732–4 ring-to-ring joints 712 rod radius 34–5 roll-to-roll processing 535 rolled solder fluxless soldering 342 rolling assisted biaxially textured substrate (RABiTS) 588 rotating part/clock plate weld 464, 467 roughness, surface see surface roughness run-to-run control methods 161 safety, and adhesives 524 sampling 161 sapphire 658 saturated vapour pressure 58 saw bands 427 SB2-Jet (Solder Ball Bumper Jet) 312 scanner processing heads 300 scanning electron microscope (SEM) conversion into a micro electron beam welder 422–4 column modifications 422–3 positioning unit modification 423–4, 425 electron beam nanowelding direct joining of CNTs 552–3 indirect joining of CNTs 553–9 scanning laser welding 681, 682 scattering 352–3 screen printing (stencil printing) 515 hot melts 538, 540, 541 screening 161 sealants adhesive bonding 527 hermetic sealing see hermetic sealing selective dispensing techniques 514, 515–18 semi-automatic bonder 225–6 semiconductor laser stripe 378–81 semicrystalline thermoplastics 666–7 sensing/sensors 159–73 medical sensors 703 separation of contaminated areas 8–14 series-mode resistance spot microwelding 478, 482 set-down (embedment) 479–80, 482, 483 setting of adhesives see curing of adhesives severe plastic deformation (SPD) 744–5
805
SHADOW laser welding 54, 83, 401–3, 404 shape memory alloys (SMAs) 620–32 application 623–4 background to joining 624–6 basics 620–2 brazing of nithinol 626–8 future trends 630 laser microwelding 628–30 shape memory effect (SME) 620, 621 shape of laser soldered joint 306–9, 323 shear joints 675 shear testing 223–4, 289, 388–9 cross-wire resistance microwelding 479, 482 sheet-to-sheet resistance microwelding 478–9 electrode force 486–7 shockwave consolidation technique 743 signals from joining processes 162–3 silane primers 513 silicates 649 silicides 648–9 silicon 658 bonding of silicon chip using composite silver nanoparticles 265–7 electron beam microwelding 448–9 melted fibres 466, 468 modelling electron beam microwelding 436–45, 446 oxygen-silicon reaction 647–8 polished polysilicon layers 657–8 polycrystalline silicon 464–6, 468 dies 456–9, 460, 461 wafer bonding 634 anodic glass-to-glass bonding using silicon layer 651–5, 656, 657 anodic silicon-to-glass bonding 641–8 anodic silicon-to-silicon bonding using glass layer 650–1 direct 635–41 silicon carbide 658 silicon rubbers 509 siloxane 509 direct wafer bonding mechanism 639–41
WPNL2204
806
Index
silver bonding using composite silver nanoparticle 255–9 composite silver nanoparticle see nanoparticles in glass composites 731 nanoparticles as interlayer for diffusion bonding 750, 751 silver-copper eutectic braze alloy 627, 628 silver carbonate 250, 253 silver oxide 591 SIMPLE algorithm 137 simultaneous welding 681, 682 single-mode fibre laser (SMFL) 403–5 single scan electron beam microwelding 445, 447 single-walled nanotubes (SWNTs) 546–7, 548 electron beam nanowelding 550–2 ultrasonic nanowelding 568–72, 573 sliding, gross 217, 218, 219, 220, 221 slip modelling solid-state bonding 101, 110–15 wire bonding and wear by microslip 217–22 small-scale resistance spot welding (SSRSW) see resistance microwelding smart materials 623 smart systems for civil structures 623 sol-gel intermediate wafer bonding 650–1 solder ball bumping 336–9 solder paste 46 solder reflow oven 184 soldering 25–50 basic metallurgical reaction 31–9 classification of method 42–4 definition 26–31 diffusion soldering see diffusion soldering/brazing fluxless see fluxless soldering for electronics assembly 45–7 future trends 48, 49 high temperature superconducting BSCCO tapes 592–3 YBCO bulks 609–12 in-process monitoring 166–7
influences on quality of joints 31 laser soldering see laser soldering materials for 39–41 soldering fluxes 40–1, 327 soldering iron 46–7 solders 26, 27, 28–31, 39–40 intermediate wafer bonding via solder 649–50 solid oxide fuel cells (SOFCs) 718–40 approaches to sealing 722–36 compliant seals 736, 737 compressive seals 734–6 rigid seals 722–34 materials involved in fabrication 720–1 solid sintering process 586–7 solid-state bonding 3–24 breakdown of interfacial structure 18–21, 22–3 contaminant displacement/interatomic bonding 7–14, 22 mechanisms 4–7 modelling see modelling thermal dissolution of oxides/ contaminants 6, 14–18, 22 solid-state diffusion bonding see diffusion bonding solidification microstructures diffusion brazing 289 fusion microwelding 79–80 laser microwelding 367–8 solidification cracking 79 solubility products 15–18 solvents 328, 513 spark plasma sintering (SPS) 747–9 spatter 60–2, 369, 370 spin-on glass films 650–1 spin welding 674, 689 spot radius, of laser beam 347–8 sputtering 333, 335 stagnation pressure 69 stainless steel 626 joining nithinol to 628 sheets 395–6 standards for adhesive bonding 540–3 static method for contact angle 503–4 static mixing system 518 statistical process control (SPC) 161 stents 701
WPNL2204
Index bare metal 691 laser cut 701, 702 nithinol 625–6 sticking 76 stigmator 420, 421 stitch bond 227–8 strain gauges 162 stress-induced transformation 759 stress testing see testing stud bumps 528 superconducting materials 583–5 see also high temperature superconductors superconducting prisms 592 superconductor paste 592 superplastic forming 236 superposition 450–3 surface deformation mechanisms 6, 7–14, 22 surface heat sources 440–5, 446 surface mounted devices (SMDs), PCB mounting of 525–6, 527 surface oxide layer see oxide surface layer surface roughness ceramic/metal bonding 780, 781 diffusion bonding 239–40 enhancing microplastic deformation of asperities 243–4 resistance microwelding 488–9 surface superheating joining 613 surface tension 55, 59–60, 72, 131–2 keyhole laser microwelding 358–60 surface treatment adhesive bonding 512–13 diffusion brazing 275 wafer bonding 635 T junctions 551, 552 tack 668 tape automated bonding (TAB) 91 inner lead bonding 93–100 teach-in programming 428–32, 433 technical data sheet 524 temperature ceramic/metal bonding brazing temperature and joint strength 770–1, 772
807
vacuum diffusion bonding 772–5, 776, 782–3 copper-to-copper joints using composite silver nanoparticle fracture morphology 263–4 joint strength 259–60 critical transition temperature 583, 584, 585 crystalline melting point 666–7 diffusion bonding of BSCCO tapes 595, 598, 599, 603, 605 diffusion soldering/brazing 278 distribution in fusion welding 53–4, 124 distribution in laser solder joint 305–6 glass transition temperature 511, 666–7, 723–4 high temperature superconductors see high temperature superconductors (HTS) homologous 745, 746 and joint strength in solid-state bonding 19–20 liquidus temperature 26–8, 28, 30 low temperature wafer anodic bonding 642–6 melting point 271 nanosensors 171 profiling 184 resistive temperature detectors (RTDs) 163 wire bonding 215–16 temperature-gradient TLP bonding (TG–TLP bonding) 284–5, 286 tensile testing see testing ternary systems 283–4 testchips 165 testing adhesive bonding 522, 542–3 diffusion soldering/brazing 289 electron beam microwelding 459, 460, 461 implantable devices 714–15 laser microwelded joints 388–9 pull testing 222–3, 388–9, 714 shear testing see shear testing thermal conductive adhesives (TCAs) 510, 533–4 thermal contours 442–5, 446
WPNL2204
808
Index
thermal deformation 394–5 thermal dissolution 6, 14–18, 22 thermal-electrical-mechanical analysis 115–17, 118, 145–6, 149–54, 155 thermal expansion/contraction 55–7, 72 thermal mechanical analysis (TMA) 511 thermal modelling laser microwelding 363–8 plastics welding 670–1 thermal signals 162–3 thermal stability 745–7 thermistors 163 thermocompression bonding 93–107 inner lead bonding in TAB 93–100 wire bonding 100–7, 108, 109 thermocouples 162–3 electron beam microwelding 463, 465 thermoplastic adhesive tape placement process 181, 182–3 thermoplastic elastomers 507 thermosonic bonding 107–15 ball bonding 114, 115, 205, 206–9, 210 gold bump bonding 109–10, 111 interfacial micro friction slip 110–12 numerical simulations of FCB 112–15 thick-film metallisation 767–8 thickness, sample 448–50 thin-die pickup process 181 thin film metallisation 768 thin film SMA 624 threshold energy 547–8 time bonding time for wire bonding 215 exposure time in air for plasma fluxless soldering 334–6 holding time see holding time pressing and annealing in diffusion bonding of BSCCO tapes 598, 599 weld time for resistance microwelding 484–6 time/pressure dispensing 516 times diffraction limit factor 349 timing control 183–4 tin diffusion soldering 270–1, 272
fluxless soldering using tin–gold solder 341–2 gold-tin IMCs 315–17, 318–20, 321–2 titanium 16 added to filler in brazing nithinol 627, 628 bonding using composite silver nanoparticle 255–9 laser welding platinum-titanium alloys 698–9, 700 titanium with oxygen in argon shielding gas 693, 698, 699 titanium alloy brazed to zirconia in BION microstimulator 780–1 top seeded melt-textured growth (TSMTG) 587–8, 589 transient thermal history 123 transition from conduction to keyhole welding 354, 357–8, 359 transmission electron microscope (TEM) direct joining of CNTs 550–2 indirect joining of CNTs 553–9 transmission laser welding 356, 357, 670–1 plastics 678–88 advantages of transmission laser welding arc 686 transparent materials 356–7, 386–7 travel speed 364–7 trim and form machine 176, 178 triode systems 419 tube-to-pin joints 713 tube-to-tube joints 712–13 tungsten arc welding 752 two-component adhesive systems 507, 518–19 ultrashort pulse fusion welding 405–8, 409 ultrasonic cavitation theory 313–14 ultrasonic fluxless soldering 339–41, 342 ultrasonic heating 216–17 ultrasonic modulated laser fluxless soldering 312–14 ultrasonic nanowelding 568–72, 573 ultrasonic sewing 678 ultrasonic softening 213–14
WPNL2204
Index ultrasonic welding 238, 248 plastics 673–8, 689, 714 advantages and disadvantages 677 equipment 675, 676 examples 677–8 process description 673–5 welding parameters 675–7 ultrasonic wire bonding 3, 101, 206, 208, 210–14 mechanisms of bond formation 216–22 ribbon bonding 210–12 wedge bonding 210, 211, 219, 220, 221 ultrasonics application in assembly process automation 185–6 defect techniques 390, 391 thermosonic bonding 107–15 underfill 528, 532 underfilled beads 370–2 upsetting (forging) 3–4, 238 UV-cured adhesives 507, 520, 524 V-grooved tool 102, 104–5, 106, 108, 109 vacancy healing 548–9 vacuum brazing 42, 44 with active filler metals 768–71, 772, 773, 774 vacuum diffusion bonding 771–80 bonding atmosphere 778–9, 782–3 bonding temperature 772–5, 776, 782–3 chemical reactions 777–8 CTE mismatch 779 holding time 775–6, 782–3 interlayer 779–80, 782–3 joint strength 773–80, 782–3 pressure 776–7, 778, 782–3 surface state 780, 781 vacuum fluxless soldering 341 vacuum wafer bonding 636, 637–8, 639 vapour recoil force 58–9, 61–2, 363 vascular devices 701–2 viscoplastic deformation model 91–3 viscosity 55, 68, 72 glass composites 730, 731 vision alignment technique 179–80
809
visual inspection diffusion brazing 286–7, 289 fusion microwelding 77–8 laser microwelded joints 387 voids copper-to-copper joints using composite silver nanoparticle 260–1 shrinkage 240 voltage recording devices 160 volume heat sources 440–5, 446 vortex, forced 65–6 wafer bonding 633–64 anodic see anodic wafer bonding direct 635–41 dissimilar materials 655–8 via intermediate layers 648–55, 656, 657–8 watch gear to shaft welds 402 wave soldering 45, 46 waveform generator 428–32, 433 weak boundary layer theory 502–3 wear 217–22 weld lobe 475 welding bulk nanomaterials 752, 753–4 plastics see plastics SMAs 624–6 see also under individual welding techniques welding current 482–4, 485 welding function 669 wetting 25, 31–5, 503–4 ceramic/metal bonding 764–6 wide-gap diffusion soldering/brazing 284, 285 Wiedemann-Franz-Lorentz law 148 wire bonding 91, 205–33 accurate timing 184 automatic wire bonders 167–9, 176, 177, 224–5 equipment and fixturing 224–6 future trends 226–9 in-process monitoring 164–6 mechanisms of bond formation 216–22 deformation 217 ultrasonic heating 216–17 wear by microslip 217–22
WPNL2204
810
WPNL2204
Index
modelling 100–7, 108, 109 process 206–12 basic description 206–9 process variants 209–12 process parameters 212–16 bonding force 214–15 bonding temperature 215–16 bonding time 215 ultrasound 212–14 quality control 222–4 wire leads 703–4 wire-to-pin joints 707–8, 709 wire-to-sheet resistance microwelding 481–2, 484 wire-to-wire joints electron beam microwelding 428, 429, 430 implantable devices 705–7 work of adhesion 764–5 work life 518 workpiece manipulation 681–2 workpiece motion system 421
workpiece viewing system 421–2 Wright model 10–13 X junctions 550–1 X-ray testing for defects 389–90, 391 Y junctions 551, 552 YBCO 583 bulks 586–8 coated tapes 588 joining of YBCO bulks 608–14 diffusion bonding 608–9 other joining processes 612–13 soldering 609–12 Young’s equation 32–3, 503, 764 yttria stabilised zirconia (YSZ) 720 zinc-based solder 40 zirconia (zirconium oxide) 731, 758, 759–61, 762 brazed joint with titanium alloy for BION microstimulator 780–1
WPNL2204
