Non-equilibrium Processing of Materials (Pergamon Materials Series)

PERGAMON MATERIALS SERIES VOLUME 2

Non-equilibrium Processing of Materials

PERGAMON

MATERIALS

SERIES

Series Editor: Robert W. Cahn F R S Department of Materials Science and Metallurgy, University of Cambridge, UK

Vol. 1 CALPHAD (Calculation of Phase Diagrams): A Comprehensive Guide

by N. Saunders and A. P. Miodownik Vol. 2 Non-equilibrium Processing of Materials edited by C. Suryanarayana A selection offorthcoming titles in this series:

Phase Transformations in Titanium- and Zirconium-based Alloys

by S. Banerjee and P. Mukhopadhyay Wettability at High Temperatures by N. Eustathopoulos, M. G. Nicholas and B. Drevet Ostwald Ripening by S. Marsh Nucleation by A. L. Greer and K. F. Kelton Underneath the Bragg Peaks: Structural Analysis of Complex Materials by T. Egami and S. J. L. Billinge The Coming of Materials Science by R. W. Cahn Thermally Activated Mechanisms in Crystal Plasticity by D. Caillard and J. L. Martin

PERGAMON MATERIALS SERIES

Non-equilibrium Processing of Materials edited by C. Suryanarayana The George S. Ansell Department of Metallurgical and Materials Engineering Colorado School of Mines Golden, Colorado, USA

1999

PERGAMON An Imprint of Elsevier Science Amsterdam- Lausanne- New Y o r k - Oxford- Shannon- Singapore- Tokyo

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First edition 1999 Library of Congress Cataloging in Publication Data A catalog record from the Library of Congress has been applied for. British Library Cataloguing in Publication Data A catalogue record from the British Library has been applied for.

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Contents xv xvii xix

Series Preface Preface List of Contributors

CHAPTER 1 INTRODUCTION

1

R. W Cahn

CHAPTER 2 THERMODYNAMICS AND KINETICS OF METASTABLE PHASE FORMATION

5

K. N. Ishihara 2.1. Introduction 2.2. Thermodynamics of Metastable Phase Formation 2.2.1 Free Energy of Elements 2.2.2 Free Energy of Alloy Phases 2.2.3 Determination of Free Energy of Metastable Phases 2.2.4 Lattice Parameter of the Supersaturated Phase 2.3. Kinetics of Metastable Phase Formation 2.3.1 Nucleation of Metastable Phases 2.3.2 Nucleation of Alloy Phases 2.3.3 Crystal Growth Rate of Metastable Phases 2.4. Summary 2.5. List of Symbols References

5 5 5 7 9 11 12 13 15 18 19 19 20

CHAPTER 3 RAPID SOLIDIFICATION

23

H. Jones

3.1. Introduction 3.2. Methods of Rapid Solidification 3.2.1 Droplet Methods 3.2.2 Spinning Methods 3.2.3 Surface-Melting Methods 3.3. Constitution and Microstructure Formation by Rapid Solidification

V

23 24 24 26 27 28

3.3.1 Non-Equilibrium Constitution 3.3.2 Microstructure Formation 3.4. Properties, Performance and Applications of Rapidly Solidified Materials 3.4.1 Properties and Performance 3.4.2 Applications of Rapid Solidification References Selected Bibliography

CHAPTER 4 MECHANICAL ALLOYING

29 32 34 34 41 43 45

49

C. Suryanarayana 4.1. Introduction 4.2. Nomenclature 4.3. The Process of Mechanical Alloying 4.3.1 Raw Materials 4.3.2 Process Control Agents 4.3.3 Type of Mills 4.3.3.1 SPEX Shaker Mills 4.3.3.2 The Planetary Ball-Mill 4.3.3.3 Attritor Mills 4.3.3.4 Commercial Mills 4.3.3.5 New Designs 4.4. Mechanism of Alloying 4.5. Consolidation 4.6. Synthesis of Non-Equilibrium Phases 4.6.1 Solid Solubility Extensions 4.6.2 Synthesis of Intermetallics 4.6.3 Disordering of Ordered Intermetallics 4.6.4 Nanocrystalline Materials 4.6.5 Solid-state Amorphization 4.6.6 Displacement Reactions 4.7. Powder Contamination 4.8. Modeling 4.9. Industrial Applications 4.9.1 Nickel-Base Alloys 4.9.2 Iron-Base Alloys 4.9.3 Aluminum-Base Alloys 4.9.4 Magnesium-Base Alloys 4.10. Concluding Remarks References

vi

49

50 51 51 52 52

53 53 53 53 54 55 56 58 59 62 64 66 67 71 74 75 76 77 79 79 80 80 81

CHAPTER 5 LASER PROCESSING

89

K. F: Kobayashi 5.1. Principles of Lasers 5.2. Classifications of Laser Processing 5.3. Analysis of the Laser Melting and Quenching Process 5.3.1 Heat Transfer 5.3.1.1 Absorption Phenomena of Laser 5.3.1.2 Pulsed Laser 5.3.1.3 Continuous Wave (CW) Laser 5.3.2 Kinetic Conditions for Solidification 5.3.2.1 Kinetic Conditions for Amorphous Phase Formation 5.3.2.2 Solidification Modes in the Laser Processing of Materials 5.3.3 Microstructure Selection Maps (MSM) for Alloys 5.4. Laser-Quenching 5.4.1 Amorphous Phase Formation 5.4.2 Formation of Crystalline Phases 5.5. Laser Surface-Alloying and Cladding 5.5.1 Ferrous Alloys 5.5.2 Non-Ferrous Alloys 5.6. Laser-Annealing 5.7. Laser-Beam Joining 5.7.1 Welding of Structural Materials 5.7.1.1 Ferrous-Based Alloys 5.7.1.2 Non-Ferrous Alloys 5.7.2 Microjoining 5.8. Conclusions References

89 89 92 92 92 93 94 96 96 97 97 98 98 99 100 100 101 102 106 106 106 108 112 115 115

CHAPTER 6 THERMAL PLASMA PROCESSING

121

I? I!Ananthapadmanabhan and N. Venkatramani 6.1. Introduction 6.1.1 Advantages of Plasma Processing 6.2. Thermal Plasmas 6.2.1 Principles of Plasma Generation 6.2.1.1 DC Plasma Torches 6.2.1.2 AC Plasma Torches 6.2.1.3 F W Plasma Torches 6.2.2 Plasmagen Gases 6.2.3 Plasma-Particle Interaction 6.2.4 Plasma Processing Systems vii

121 121 123 123 124 126 126 127 128 129

6.3.

Processing of Materials 6.3.1 Plasma-Spraying 6.3.1.1 Structure of Sprayed Deposits 6.3.1.2 Plasma Spray-Deposited Materials 6.3.1.3 Low-Pressure Plasma-Spraying 6.3.1.4 Reactive Plasma-Spray Forming 6.3.1.5 Plasma Spheroidization 6.3.2 Plasma Reactors and Furnaces 6.3.2.1 Plasma Decomposition 6.3.2.2 Plasma Metallurgy 6.3.2.3 Processing of Ceramics 6.3.2.4 Treatment of Hazardous Wastes 6.3.3 Processing of Metastable Phases 6.3.3.1 Plasma Deposition of Diamonds 6.3.3.2 Thermal Plasma Synthesis of Ultrafine Alumina Powder 6.4. Summary and Conclusions Acknowledgments References

129 129 131 132 134 135 135 137 137 138 140 141 141 141 142 147 148 148

CHAPTER 7 153

SPRAY-FORMING Bing Li and E. J. Lavernia 7.1.

7.2. 7.3.

7.4.

7.5.

Introduction Principles Variations and Distinctions 7.3.1 Variations 7.3.1.1 Crucible: Induction Skull Melting/Spray-Forming 7.3.1.2 Atomizer: Circular vs. Linear Spray-Forming 7.3.1.3 Atomizer: Close-Coupled vs. Free-Fall Spray-Forming 7.3.1.4 Atomization Gas: Reactive Spray-Forming 7.3.1.5 Substrate: Near Net Shape Spray-Forming 7.3.1.6 Spray-Forming and Co-Injection 7.3.2 Nomenclature 7.3.2.1 Spray Atomization 7.3.2.2 Related Processing Techniques Applicability Non-Equilibrium Phenomena in Spray-Forming 7.5.1 Non-Equilibrium Nature 7.5.1.1 Rapid Solidification in Atomized Droplets 7.5.1.2 Transient Semi-Solid Layer in Deposition Stage 7.5.2 Non-Equilibrium Related Features in the Deposit 7.5.2.1 Metastable Phases 7.5.2.2 Extended Solid Solubility 7.5.2.3 Absence of Macrosegregation/Minimized Microsegregation viii

153 154 157 157 158 158 159 160 163 163 166 166 167 168 172 173 173 176 178 178 178 179

7.5.2.4 Refinement Effects of Non-Equilibrium Features on Mechanical/Physical Properties 7.5.3.1 Direct Effects 7.5.3.2 Indirect Effects 7.6. Concluding Remarks Acknowledgments References

180

7.5.3

181 183 185 189 189 189

CHAPTER 8 ION-MIXING

197

B. X. Liu

8.1. 8.2. 8.3.

Introduction Brief Description of Underlying Physics in Ion Mixing Thermodynamics of Alloy Phase Formation 8.3.1 Miedema's Theory and Alonso's Method 8.3.2 Interfacial Free Energy in the Multilayers 8.4. Experimentation of Ion-Mixing 8.4.1 Sample Design 8.4.2 Ion-Mixing Parameters 8.4.3 Characterization Methods 8.5. Amorphous Phase Formation 8.5.1 Glass-Forming Ability of Systems with a Negative Heat of Formation 8.5.2 Glass-Forming Ability of Systems with a Positive Heat of Formation 8.5.2.1 Amorphous Alloys Formed Within Restricted Compositions 8.5.2.2 Amorphous Alloys Formed in a Broad Composition Range 8.5.2.3 Nominal and Intrinsic Glass-Forming Ability 8.6. Formation of Metastable Crystalline Alloys 8.6.1 Structural Classification of the Metastable Crystalline Phases 8.6.1.1 Solid Solutions 8.6.1.2 h.c.p.-I and f.c.c.-I Phases 8.6.1.3 h.c.p.-II and f.c.c.-II Phases Based on b.c.c. Metals 8.6.2 Free Energy Calculation of the MX Phases 8.7. Interface-Generated Solid-State Vitrification in Systems with a Positive Heat of Formation 8.8. Concluding Remarks Acknowledgments References

197 199 200 200 201 203 203 205 205 205 206 207 207 207 211 213 213 213 213 215 216 218 219 220 220

CHAPTER 9 PHYSICAL VAPOR DEPOSITION

225

J. S. Colligon

ix

9.1. 9.2. 9.3.

Introduction Development of PVD Deposition Methods 9.3.1 Evaporation 9.3.2 Molecular Beam Epitaxy (MBE) 9.3.3 Sputtering 9.3.4 Ion-Assisted Deposition 9.3.5 Magnetron Sputtering 9.4. Influence of Energy on Coatings 9.4.1 Film Morphology and Density 9.4.2 Nucleation and Adhesion 9.4.3 Metastable Phases 9.4.4 Microstructure 9.5. Applications of PVD Coatings 9.5.1 Optical Coatings 9.5.2 Corrosion Protective Coatings 9.5.3 Hard Coatings 9.6. Future Trends Acknowledgment References

225 225 228 228 229 229 231 235 238 238 239 240 240 243 243 245 246 248 249 249

C H A P T E R 10 CHEMICAL VAPOR DEPOSITION

257

E Teyssandier and A. Dollet

10.1. Introduction 10.1.1 Presentation 10.2. Gas-Phase Transport and Reactivity 10.2.1 Non-Reactive Fluid Flow 10.2.1.1 Geometrical Effects 10.2.1.2 Influence of Operating Conditions 10.2.2 Reactive Flows 10.2.2.1 Thermodynamic Approach 10.2.2.2 Gas-Phase Mechanisms and Kinetics 10.2.2.3 Simulations of the Gas-Phase Composition 10.2.2.4 Nucleation and Growth of Solid Particles from the Gas Phase 10.3. Solid Phase Formation 10.3.1 Solid-Gas Thermodynamic Equilibrium Approach 10.3.2 The Structure of CVD Layers 10.3.3 The Driving Force for Crystal Growth from the Vapor Phase 10.3.4 Surface Mechanisms and Kinetics 10.3.4.1 Rate of Elementary Surface Processes 10.3.4.2 Estimation of the Kinetics of Surface Processes in CVD 10.3.5 Heterogeneous Nucleation 10.3.5.1 Nucleation Theories X

257 258 259 260 260 261 261 263 264 264 265 267 267 268 270 271 271 272 273 274

10.3.5.2 10.3.5.3 10.3.5.4 10.3.5.5 10.3.5.6 10.3.5.7

Two- and Three-Dimensional Nucleation Epitaxial Nucleation and Growth Multi-Component Nucleation and Chemical Reactions Nucleation Enhancement Selective Vapor Deposition Theoretical and Experimental Studies of Nucleation in CVD 10.3.6 Crystal Growth 10.3.6.1 Mechanisms of Crystal Growth 10.3.6.2 Classification of Crystal Faces 10.3.6.3 Crystalline Morphology 10.3.6.4 Morphology and Normal Growth Rate of Crystal Faces 10.3.6.5 Relation between Preferred Orientation and Morphology of Polycrystalline Films 10.3.6.6 Theoretical Studies of Crystal Growth from the Vapor Phase 10.4. Conclusions References

275 275 276 276 277 277 278 278 278 279 279 281 281 281 282

C H A P T E R 11 COMBUSTION

SYNTHESIS

289

S. B. Bhaduri and S. Bhaduri

11.1. Introduction 11.2. Thermodynamic Considerations 11.3. Kinetic Considerations 11.4. Field-Activated Combustion Synthesis 11.5. The "Azide" Process 11.6. SHS Reactions in Synthesizing Ti3SiC2 11.7. Controlled Reactions in the Ti-B Binary System 11.8. Auto-Ignition Synthesis of Nanocrystalline Oxides 11.9. Non-Equilibrium Effects 11.10. Concluding Remarks Acknowledgments References

289 291 294 295 297 299 301 304 307 307 308 308

C H A P T E R 12 NANOSTRUCTURED

MATERIALS

313

c. Suryanarayana and C. C. Koch

12.1. 12.2. 12.3.

Introduction Classification Preparation 12.3.1 Inert Gas Condensation

313 313 314 314 xi

12.3.2 Rapid Solidification 12.3.3 Electrodeposition 12.3.4 Chemical Reactions 12.3.5 Mechanical Attrition 12.3.6 Devitrification 12.4. Structure 12.4.1 Microstructure 12.4.2 Atomic Structure of the Crystal Lattice 12.4.3 Atomic Structure of the Grain Boundaries 12.4.4 Triple Junctions and Higher-Order Grain Junctions 12.5. Stability 12.5.1 Grain Growth in Nanocrystalline Materials 12.5.2 Grain Growth at Ambient Temperature 12.5.3 Examples of Grain Growth Inhibition 12.5.4 Isothermal Grain Growth Kinetics 12.6. Particulate Consolidation 12.7. Properties 12.7.1 Diffusion and Sinterability 12.7.2 Mechanical Properties 12.7.2.1 Elastic Properties 12.7.2.2 Hardness and Strength 12.7.2.3 Ductility and Toughness 12.7.2.4 Superplastic Behavior 12.7.2.5 Deformation Mechanisms in Nanoscale Materials 12.7.3 Magnetic Properties 12.7.3.1 Fundamental Properties 12.7.3.2 Soft Magnetic Materials 12.7.3.3 Hard Magnetic Materials 12.7.3.4 Other Ferromagnetic Nanocrystalline Materials 12.7.4 Chemical Properties 12.7.4.1 Corrosion Behavior 12.7.4.2 Catalytic Properties 12.8. Applications--Present and Potential 12.8.1 Structural Applications 12.8.1.1 Cutting Tools 12.8.1.2 Nanocomposites 12.8.1.3 Superplastic Materials 12.8.1.4 Coatings 12.8.2 Magnetic Applications 12.8.3 Catalysts and Hydrogen Storage Materials 12.8.4 Functional NanostructuresmElectronic Applications 12.9. Concluding Remarks References

xii

316 317 318 318 319 319 320 321 321 322 323 323 324 324 325 326 327 327 328 328 329 329 331 331 333 333 333 333 335 335 335 335 337 337 337 338 339 339 340 340 341 341 342

CHAPTER 13 POWDER CONSOLIDATION

347

J. R. Groza

13.1. 13.2. 13.3.

Introduction Metastability in Powder Consolidation Consolidation of Metastable Powders 13.3.1 Sintering Mechanisms 13.3.2 Scaling Laws 13.3.3 Powder Contamination 13.4. Consolidation Methods 13.4.1 Conventional Sintering 13.4.1.1 Cold Compaction 13.4.1.2 Warm Compaction 13.4.1.3 Conventional Sintering of Nanopowders 13.4.1.4 Nano-Composite Densification 13.4.1.5 Grain-Size Control 13.4.2 Pressure Consolidation of Metastable Powders 13.4.3 Pressure-Assisted Consolidation Methods 13.4.4 Non-Conventional Sintering Methods 13.4.4.1 Microwave Sintering 13.4.4.2 Field-Assisted Sintering 13.4.4.3 Shockwave Consolidation 13.5. Concluding Remarks Acknowledgments References

347 348 349 352 353 355 356 356 356 358 359 361 361 363 364 366 366 366 368 368 369 369

CHAPTER 14 BULK AMORPHOUS ALLOYS

375

A. Inoue

14.1. 14.2. 14.3. 14.4. 14.5. 14.6.

14.7. 14.8. 14.9.

History of Bulk Amorphous Alloys Dominant Factors for High Glass-Forming Ability Continuous Cooling Transformation of Alloys with High Glass-Forming Ability Preparation Methods and Elemental Effect Structural Relaxation and Glass Transition Physical Properties 14.6.1 Density 14.6.2 Electrical Resistivity 14.6.3 Thermal Expansion Coefficient Mechanical Properties Viscoelasticity Soft Magnetic Properties xiii

375 376 380 383 387 392 392 393 394 395 399 403

14.9.1 Formation and Soft Magnetic Properties of Bulk Amorphous Alloys 14.10. Viscous Flow and Micro-Formability of Supercooled Liquid 14.10.1 Feature of Phase Transition of Bulk Amorphous Alloys 14.10.2 Deformation Behavior of Supercooled Liquid 14.10.3 Micro-Forming of Supercooled Liquid 14.11. Applications and Future Prospects References

403 406 409 409 412 413 413

Author Index

417

Subject Index

419

xiv

Series Preface

My editorial objective in this new series is to present to the scientific public a collection of texts that satisfies one of two criteria: the systematic presentation of a specialised but important topic within materials science or engineering that has not previously (or recently) been the subject of full-length treatment and is in rapid development; or the systematic account of a broad theme in materials science or engineering. The books are not, in general, designed as undergraduate texts, but rather are intended for use at graduate level and by established research workers. However, teaching methods are in such rapid evolution that some of the books may well find use at an earlier stage in university education. I have long editorial experience both in covering the whole of a huge field- physical metallurgy or materials science and t e c h n o l o g y - and in arranging for specialised subsidiary topics to be presented in monographs. My intention is to apply the lessons learnt in more than 35 years of editing to the objectives stated above. Authors (and, in some instances, editors) have been invited for their up-to-date expertise and also for their ability to see their subjects in a wider perspective. I am grateful to Elsevier Science Ltd., who own the Pergamon Press imprint, and equally to my authors, for their confidence.

ROBERT W. CAHN, FRS

(Cambridge University, UK)

XV

This Page Intentionally Left Blank

Preface

Metals and materials have been used for a long time. But the rapid technological developments during the later half of the twentieth century demanded materials that are stronger, capable of use at much higher temperature, more corrosion-resistant, and much less expensive than those currently in use. These demands become even more significant when we are on the threshold of the new century and the millenium. Significant improvements in properties can only be achieved by processing the materials under far-from-equilibrium (or non-equilibrium) conditions. Accordingly, several new processing technologies have been developed during the past few decades, including rapid solidification, spray-forming, mechanical alloying, ion-mixing, vapor-deposition, laser processing, plasma-processing, etc. New types of materials such as amorphous alloys and nanostructured materials have also been developed. Professor Robert Cahn has put these developments in a proper perspective in his introductory chapter, which also sets the tone for the chapters that follow. Remarkable advances have been made in recent years in the science and technology of these processes used to synthesize, characterize, and apply materials processed under non-equilibrium conditions. There have been many success stories. Some of these techniques have evolved from laboratory curiosities to commercial-scale manufacturing in a few short years. In other cases, industrial necessity prompted development of the technology, and the science followed later. The very large number of technical papers published in these different areas have been summarized in reviews, monographs, or conference proceedings, but have almost always been devoted to a single processing technique. I felt that a book of this sort was needed to provide a critical and comparative study of the different processes in a single volume. The different chapters in this book have been written by people who are worldrecognized experts in their respective fields. Each chapter describes the principles, processing techniques, special features of the materials produced and their applications. An extensive list of references is provided at the end of each chapter which will facilitate location of additional information on specific aspects of any technique. Even though great care has been taken to provide coherency and uniformity amongst the different chapters, the individual style of the authors has not been completely sacrificed. After all, variety is the spice of life! The book is primarily intended for use by graduate-level students of materials science and engineering. However, scientists beginning their research careers in any of the processes covered in this book can use it to get a fairly good knowledge and overview of the state of the technique and its potential. Scientists who wish to enter this area of research can also use this book to get a comparative idea of the potential and limitations of the different non-equilibrium processing techniques.

xvii

I wish to thank all the contributing authors for their considerable efforts in summarizing the voluminous literature of their specialties, and their dedication, diligence, and patience in working with me (and tolerating the many editorial iterations). The encouragement provided by Robert Cahn (Cambridge, UK) and John Moore (Golden, CO, USA) is greatly appreciated. I would like to express my appreciation to Phil Mestecky and his staff at Elsevier for bringing this book out in a timely and attractive fashion. Finally, I wish to thank my wife Meena for all her support during the preparation of this volume.

Golden, CO, USA August 31, 1998

C. SURYANARAYANA

xviii

List of Contributors

E V. ANANTHAPADMANABHAN Laser and Plasma Technology Division, Bhabha Atomic Research Centre, Mumbai400 085, India S. BHADURI Department of Metallurgical and Mining Engineering, University of Idaho, Moscow, ID 83844-3024, U.S.A. S. B. BHADURI Department of Metallurgical and Mining Engineering, University of Idaho, Moscow, ID 83844-3024, U.S.A. R. W. CAHN Department of Materials Science and Metallurgy, University of Cambridge, Pembroke Street, Cambridge CB2 3QZ, U.K. J. S. COLLIGON Manchester Metropolitan University, Chester Street, Manchester M1 5GD, U.K. A. DOLLET Institut de Science et de Grnie des Matrriaux et Procrdrs, CNRS, UPR8521, Universitr, Avenue de Villeneuve, F-66860 Perpignan Cedex, France J. R. GROZA Department of Chemical Engineering and Materials Science, University of California at Davis, One Shields Ave., Davis, CA 95616, U.S.A. A. INOUE Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan K. N. ISHIHARA Department of Socio-Environmental Energy Science, Kyoto University, Kyoto 6068501, Japan H. JONES Department of Engineering Materials, University of Sheffield, Mappin Street, Sheffield S1 3JD, U.K.

xix

K. E KOBAYASHI Department of Manufacturing Science, Graduate School of Engineering, Osaka University, 2-1 Yamadaoka, Suita, Osaka 565-0871, Japan C. C. KOCH Department of Materials Science and Engineering, North Carolina State University, Raleigh, NC 27695-7907, U.S.A. E. J. LAVERNIA Department of Chemical & Biochemical Engineering and Materials Science, University of California, Irvine, CA 92697-2575, U.S.A. B. LI

Department of Chemical & Biochemical Engineering and Materials Science, University of California, Irvine, CA 92697-2575, U.S.A. B. X. LIU

Department of Materials Science and Engineering, Tsinghua University, Beijing 100084, China C. SURYANARAYANA The George S. Ansell Department of Metallurgical and Materials Engineering, Colorado School of Mines, Golden, CO 80401-1887, USA E TEYSSANDIER Institut de Science et de G6nie des Mat6riaux et Proc6d6s, CNRS, UPR8521, Universit6, Avenue de Villeneuve, F-66860 Perpignan Cedex, France N. VENKATRAMANI Laser and Plasma Technology Division, Bhabha Atomic Research Centre, Mumbai400 085, India

XX

Chapter 1 Introduction ROBERT W. CAHN

Man has used some of his materials in non-equilibrium form for a long time: The Japanese samurai sword of the thirteenth century, with its martensitic edge; age-hardened aircraft alloys, introduced early in this century; vapor-quenched superconducting alloys in the 1930s; melt-spun transformer laminations consisting of metallic glass in the 1970s. It was the chance discovery of age-hardening by Wilm at the start of this century, supplemented by Merica's recognition that the hardening stemmed in some way from a supersaturated solid solution, that really drew metallurgists' attention to the practical importance of non-equilibrium states. From the mid-1930s onwards, crystallographers like Guinier and Preston, physical chemists like Turnbull and physical metallurgists like Nutting, devoted their skills to unravelling the nature of age-hardening, at the same time as other metallurgists such as Kurdyumov in Russia and Cohen in the USA began the long task of understanding the details of the martensitic transformation in steels. Metastability had become a way of life for some metallurgists. The attitude of metallurgists to metastability underwent a sudden sea-change when Pol Duwez in California, in 1960, invented what at first came to be called splat-quenching (a term which, incidentally, he hated) and later, rapid solidification processing. This novel technique, centred on rapid quenching of hot molten alloys, introduced to the materials community things like hugely supersaturated solid solutions and, particularly, metallic glasses. Like all important innovations, this one did not spring fully formed from its progenitor's fertile imagination. Some years earlier, Duwez had studied the cooling rates accessible with a thin solid foil and a fast-moving stream of helium, and found them limited to around 10,000 K/s. A number of engineers had used rapid freezing, for instance of wires, purely as a means of manufacturing convenient shapes cheaply. Duwez realized that, to make progress, he had to quench, not a solid but a melt. What was new about Duwez's researches was that, for him, rapid freezing was a means to enhance accessible cooling rates by a large factor, and thereby to generate new states of matter: it was a scientific study in the first instance, and it led to technological advances as a by-product. What went on at California Institute of Technology at the beginning of the 1960s was a textbook illustration of the principle that technological innovation often (by no means always) springs from new scientific insights, and from experiments designed to create such insights.

2

R.W. Cahn

Once it had become clear that strength, in particular, could be greatly enhanced by supersaturating solid solutions and then heat-treating such non-equilibrium structures, metallurgists began to examine other ways of achieving such structures. Mechanical alloying began in the nickel industry as a way of making useful nickel-base alloys, and was then adopted by academically inclined researchers as a way of producing alloys in non-equilibrium states: in particular, alloys such as Ni3A1 which in equilibrium are fully ordered could be produced in disordered form, and some of these artificially disordered alloys then turned into glasses; this was a kind of two-stage metastability. Physical vapor deposition, which as mentioned above had been studied since the 1930s, was looked at far more systematically, and new devices such as lasers (that notorious family of solutions looking for a problem) were pressed into service both for this and for other processing strategies. The microelectronic industry put great effort into chemical vapor deposition and into ion implantation and ion mixing, and some of this work gave new ideas to classical metallurgists. Some years ago, an influential committee of the National Academies of Science and Engineering in the USA pointed to the overwhelming importance of improvements in processing to the national economy. Non-equilibrium processing of the types mentioned can be said to have responded to this need. Research was not limited to the academic (as in much recent work on mechanical alloying and milling) or to the tiny, as in microelectronics. New major processing techniques were improved: plasma-processing and spray-forming are large-scale techniques which happen to incorporate non-equilibrium states of matter. Another approach, mostly pioneered in Russia, was combustion synthesis (which, like splat-quenching, passed through several names); long ago, this approach was used to weld together streetcar rails in situ, but it has come a long way since then, and recently has turned full circle as a way of improving casting technique. Finally, and in some ways the most exciting, is the development of research on nanostructures, introduced to the world of modem science by Gleiter in Germany. Nanostructural features affect both functional materials (quantum dots, soft magnets) and structural materials. A century earlier, something similar had formed the subject-matter of the science of colloid science, but this never really took off as a wholly respectable science: maybe, in part, this was because early colloid scientists had poured cold water on the hypothesis that giant (macro-) molecules can form from organic monomers, and held back the development of polymer science by a couple of decades at least. It is always hard to forgive willful obscurantists! Now, many years later, colloid scientists (they still exist) have made fascinating discoveries about "colloidal crystals", consisting of equisized colloidal particles rather than of atoms. There have been books devoted to rapid solidification, to nanostructures, to metallic glasses (before they found application in that most cogent of industries, that of golf), to laser-processing and to other techniques toombut this, I do believe, is the first book to cover these variegated techniques in a critical and comparative way. I have long known Prof. Suryanarayana as a masterly editor, and here he has excelled himself.

Chapter 2 Thermodynamics and Kinetics of Metastable Phase Formation 2.1. 2.2.

Introduction Thermodynamics of Metastable Phase Formation 2.2.1 Free Energy of Elements 2.2.2 Free Energy of Alloy Phases 2.2.3 Determination of Free Energy of Metastable Phases 2.2.4 Lattice Parameter of the Supersaturated Phase 2.3. Kinetics of Metastable Phase Formation 2.3.1 Nucleation of Metastable Phases 2.3.2 Nucleation of Alloy Phases 2.3.3 Crystal Growth Rate of Metastable Phases 2.4. Summary 2.5. List of Symbols References

5 5 5 7 9 11 12 13 15 18 19 19 20

This Page Intentionally Left Blank

Chapter 2 Thermodynamics and Kinetics of Metastable Phase Formation KEIICHI N. ISHIHARA

2.1. I N T R O D U C T I O N

Metastable phases are commonly seen in nature. Typical examples are the occurrence of martensite in steel and diamond form of carbon. Recently, non-equilibrium processes such as rapid solidification, mechanical alloying and vapor deposition have been developed enabling formation of a variety of metastable phases. The fundamentals of some of these processing techniques have been reviewed (Turnbull 1981, Shingu 1992, Boettinger and Perepezko 1993, Herlach 1994, Baricco 1995). In this chapter, the thermodynamics and kinetics of metastable phase formation will be treated. The thermodynamics of competitive phases will be first discussed followed by the nucleation problem of metastable phase formation. Thermodynamically, reaction or phase transformation can occur only when the free energy of the system decreases. However, there is the possibility that it may not occur. For example, the free energy of diamond is higher than that of graphite at ambient pressure and room temperature. But diamond is quite "stable" at room temperature and pressure without any noticeable change occurring. Put differently, thermodynamics tells us that a reaction cannot occur when the free energy of the system increases. In the following paragraphs, we will show how the formation of a metastable phase can be expected to occur. Then we will describe how a phase can be formed under a wide range of thermodynamical possibilities. This is kinetics in contrast to thermodynamics, which only tells us whether, or not, a phase can form. In this part, the classical nucleation theory will play an important role. Using this theory, Ostwald's step rule will show that it is the metastable and not the stable phase that is transformed.

2.2. T H E R M O D Y N A M I C S O F M E T A S T A B L E P H A S E F O R M A T I O N

2.2.1 Free energy of elements Figure 2.1 shows schematically the free energies of the or,/~, y and 3 solid phases and the liquid phase of a hypothetical element. In this case, the/3, V and 6 solid phases are metastable with respect to the ot phase since the free energies of these phases are higher

6

K.N. Ishihara

0

6 solid

2 solid solid Liquid

0

Tg

T~m

T~

Tm

Temperature Figure 2.1. Schematic free energy hierarchies.

than that of the u phase at all temperatures. Tm ~ represents the melting point of the ot phase, where the free energy of the u phase is the same as that of the liquid phase. Similarly, Tm ~ and Tmy represent the melting temperatures of the fl and y phases, respectively. The Gibbs free energies of the different competing phases can be compared, at fixed temperature and pressure, to determine the relative stabilities of the phases under equilibrium. At a temperature higher than Tm ~ the free energy of the liquid phase is the lowest, and those of the c~,/3, y and 6 phases are higher in that order. This means that while liquid is the most stable phase under these conditions, the or, /3, y and 6 phases are less stable. We call this the hierarchy of stability. The hierarchy of stability of the different phases changes through a crossover of the free energy curves during a change in temperature and/or pressure. We will, however, ignore the pressure effect in this chapter. In this case, for the range of undercooling from Tm ~ to Tm ~, ot is the stable phase, the free energy of which is less than that of the liquid phase. Only the u phase can be formed from the undercooled liquid in this temperature region. Below Tm ~ it is possible to solidify the fl phase directly from the liquid. Therefore, there are two different ways in which the liquid can solidify. Thermodynamics does not tell us which phase will form preferentially from the liquid state. This will be discussed in the section dealing with the kinetics. In Fig. 2.1, the free energy of the 6 phase is higher than that of the liquid phase at all temperatures. Therefore, the 6 phase cannot form from the undercooled liquid. However, it may form from the vapor phase, depending on the relative position of the free energy of the vapor and 3 phases. The position of the amorphous phase is also indicated in this figure. The amorphous phase is the frozen state of an undercooled liquid. Since the free energy of the amorphous

Thermodynamics and Kinetics of Metastable Phase Formation phase depends on the frozen temperature, called the glass transition temperature Tg, the free energy cannot be precisely determined. The free energy of the amorphous phase is always higher than that of the liquid phase. Experimental determination of the free energy hierarchy has been performed for various elements. When the difference of specific heat between the liquid and solid phases is known, the difference in free energy, A G can be calculated using the equation:

A G - - AHf(Tm Tm - T) - fT Tm ACpdT + T fTTm ACP T dT

(2.1)

where AHu is the enthalpy of fusion, Tm is the melting temperature, and ACp is the difference of specific heat between the liquid and solid phases as a function of temperature. If no experimental data are available, the best approximation is to set A Cp = 0. The various expressions for this can be found in Kelton (1991). The free energies of different metallic elements are summarized by SGTE (Scientific Group Thermodata Europe) and can be found in CALPHAD (1991, 15, 317).

2.2.2 Free energy of alloy phases The minimum degree of undercooling which is thermodynamically necessary for the diffusionless transformation of a liquid alloy into a solid solution for a given alloy composition, is expressed by the To curve in the phase diagram. Figure 2.2 shows part of a hypothetical binary phase diagram and the free energy vs. composition curves at a temperature 7"1. Since the common tangent from A to B is lower than the free energy of either the single phase liquid or solid phases, the most stable state at the concentration from CA to C B is a mixture of the liquid and solid phases. This is indicated as the two-phase region in the phase diagram. However, in the case of partitionless solidification that is often observed during rapid solidification, the position of the free energy curve becomes important. Thus, at 7'1, in the concentration range below Co, the solid phase is more stable than the liquid phase and above Co the liquid phase is more stable than the solid phase. The trace of Co at which the solid and liquid phases have the same free energy can be expressed in the phase diagram; it is called To curve indicated by the dashed line in Fig. 2.2. If the To curve plunges to very low temperatures, single-phase crystals cannot form from the melt. In this case, phase separation from the liquid phase depresses the solidification temperature to the glass transition temperature, Tg, where the liquid transforms to the amorphous phase. The system with the To curves, only slightly depressed below the stable liquidus, is a good candidate for partitionless solidification in the entire composition range (Herlach 1994). It is very important to know the position of the To curves in the phase diagram. Since the equilibrium phase diagram is a way of expressing thermodynamic data, there might be a simpler way of finding the approximate position of the To curve from the general features of the solidus and liquidus lines. The simplest guess, which often works, is to draw the To curve connecting the midpoints between the liquidus and solidus lines at a given temperature.

K. N. Ishihara

uid

~

0

0

i i i i

I

CA

CO

CB

i

(D .4,..a

c~ ~D T1

iquidus Solidus ~

Concentration Figure 2.2. Part of a hypothetical binary phase diagram and the free energy curves of the solid and liquid phases at temperature T 1.

If we know the free energy as a function of temperature and composition, the metastable phase diagrams including To curves and extrapolation of phase boundaries can be drawn. The free energy function of the binary system is now available for the purpose of calculating the phase diagram. This calculation of the phase diagram using the free energy function is called the CALPHAD (CALculating the PHAse Diagram) method (Saunders and Miodownik 1998). Thermocalc program module that contains the database of these functions is useful for this purpose. The recent edition of the ASM handbook on phase diagrams includes reference to the calculation of the phase diagrams (Massalski et al. 1990). The specific journals available for the assessment of free energy functions in connection with phase diagrams are "CALPHAD" and Journal of Phase Equilibria.

Thermodynamics and Kinetics of Metastable Phase Formation

9

2.2.3 Determination o f free energy o f metastable phases

Direct measurement of thermodynamic data of metastable phases is difficult and only a limited data is available. This has been done by measuring either the heat released during transformation of the metastable phase to the stable form or by measuring the electromotive force (emf). Measurement of the heat released on transforming to the stable phase is one of the simplest methods. Usually this measurement has been performed by DSC (Differential Scanning Calorimetry) (Michaelsen et al. 1997), especially to determine the heats of crystallization of the amorphous phases. The specific heat can also be measured by DSC. The specific heat and enthalpy of transformation for a metastable solid phase have been measured and the free energy is calculated using Eq. (2.1) for the pure metal gallium. An example of the experimentally determined free-energy hierarchy for gallium is shown in Fig. 2.3 (Bosio et al. 1973). ot is the stable phase and fl and y are metastable phases in Fig. 2.3. Measurements of emf for the stable and metastable phases have also been reported (Ishihara et al. 1993, G~xtner et al. 1996). In this method, only the partial quantity of free energy, activity, can be directly measured. Such a measurement for the Ag-Cu system is shown in Fig. 2.4. When any thermodynamic value is not available, the enthalpy of mixing based on Miedema's macroscopic atom model can be used (de Boer et al. 1988). This approach gives us not only the enthalpies of mixing for the liquid and solid phases for a binary alloy, but also the interfacial energy between solid phases. Since in the macroscopic atom model, atoms are conceived as 'blocks' of the element, assuming the size of the block as the grain size the enthalpy of the structure including the grain boundaries can be obtained. The enthalpy of nanostructures, frequently obtained by mechanical alloying and other

a 1500

~,

-

~ 10oo 500

0

100

200

300N~

T e m p e r a t u r e (K) Figure 2.3. Experimentallydeterminedfree energyhierarchyfor Gallium (Bosioet aL 1973).

10

K. N. I s h i h a r a

10000

.

.

.

.

.

.

.

.

.

I 9

1

@

1000 I:- ~

't

0 1O0

> r @

10

1

0.0

0.2

0.4

0.6

0.8

1.0

Cu Figure 2.4. Activity of Cu in Ag-Cu solid solution phase measured by the emf method. Solid circles and diamonds show the values measured using the CuSO4 solution and the solid electrolyte RbCulCI, respectively.

(a)

(b)

~xo

9

A

B

Concentration

A

B

Concentration

Figure 2.5. Changes of enthalpy by grain refinement for the case of systems with (a)

positive heat of mixing and (b) negative heat of mixing.

Thermodynamics and Kinetics of Metastable Phase Formation

11

methods, can be evaluated by this method. In this case, the enthalpy of the nanostructure should be close to that of the liquid or amorphous state so that the grain size reaches nanometer dimensions (Bakker 1998). The enthalpy for the simple mixing of powders can be indicated by a horizontal line in Fig. 2.5. When the grain size becomes small, the total enthalpy of grain boundary, AHgb changes to closely resemble the enthalpy of mixing of the liquid phase. In the case of a system which has a positive heat of mixing, A Hgb should raise its value, while in the case of the negative heat of mixing, A Hgb should increase at both low and high concentration ranges but may decrease in the middle concentration range, as shown in Fig. 2.5 (a) and (b). It means that during mechanical alloying, the enthalpy of mixing of the solid phase does not always increase and approach that of the liquid. 2.2.4 Lattice parameter o f the supersaturated phase

The lattice parameter of a solid solution phase varies linearly with composition; this is known as Vegard's law. If the crystal structures of the terminal solid solutions are different, the average atomic volume varies linearly with the concentration. A positive deviation from Vegard's law is often seen in nature (Najafabadi et al. 1993). Extension of solid solubility is commonly observed during non-equilibrium processing and the deviation from Vegard's law has been confirmed for many systems and for wide composition ranges. Several attempts have been made to explain this deviation on the basis of the theory of elasticity (Gschneidner and Vineyard 1962). However, this deviation is attributable to the excess energy of the solid and can be thermodynamically explained using the enthalpy of mixing. Here, a simple idea will be presented. One must first consider how much energy can be stored by the deviation from the ideal volume which is shown by the Vegard's law. This energy E can be expressed as (Ziman 1972): E = (B/v)(AV/V)

(2.2)

where B is the bulk modulus, 7/is the Griineisen constant and A V is the change in the unit cell volume. In the case of a system involving metals with a cubic structure, A V can be expressed in terms of the lattice parameter a, as A V~ V ~- 3 A a / a

(2.3)

where V and a are the ideal unit cell volume and lattice parameter, respectively, according to Vegard's law and Aa is the deviation of the lattice parameter from the ideal value. This energy should be equal to the excess enthalpy of mixing. In the case of a regular solution model, it can be shown that E

= CACB~']

(2.4)

where f2 is the interaction parameter. Then, A a / a = (v/3B)CACB~2

(2.5)

K. N. Ishihara

12 0.006

!

!

!

i

0.004

<~ 0.002

0.000

,

0.0

I

0.2

,

I

I

0.4

0.6

,

I

0.8

1.0

Cu Figure 2.6. Deviation of the lattice parameter form Vegard's law, Aa for the AgCu system. The open and filled circles show the reported results of ball-milled and rapidly-quenched samples, respectively (Najafabadi et al. 1993).

In this model, the deviation from Vegard's law for the regular solution model is parabolic, which is consistent with the reported experimental data. Moreover, the sign of the deviation is the same as the sign of the heat of mixing. This is the reason why a positive deviation from Vegard's law is often seen in nature. This is because, in the case of the system that has a negative heat of mixing, intermetallic compounds are formed rather than a supersaturated solid solution. In the Ag-Cu system, the calculated value of Aa calculated using the heat of mixing by the CALPHAD method (Murray 1984) matches moderately well with the experimental results shown in Fig. 2.6 for y = 2 and B calculated as the average of 1.0 x 10 ll N/m 2 for Ag and 1.5 x 10 ll N/m 2 for Cu.

2.3. K I N E T I C S OF M E T A S T A B L E P H A S E F O R M A T I O N

Many metastable phases are found in nature and have also been produced by the recently developed novel materials processing techniques. According to thermodynamics, as mentioned above, there is a possibility for both stable and metastable phases to form. So, which is more probable? This question can be answered by the kinetics of the process. There are two main reasons why metastable phases form. The first reason is that a metastable phase can nucleate before the stable phase can nucleate. The second reason is that the crystal growth rate of the metastable phase exceeds that of the stable phase. There have been attempts to draw kinetic phase diagrams (Chvoj 1983). However, no systematic study in this field has been performed as compared to the thermodynamic in-

Thermodynamics and Kinetics of Metastable Phase Formation

13

vestigations. In this section, the nucleation and the growth rates of the metastable and stable phases will be discussed.

2.3.1 Nucleation of metastable phases The free energy hierarchy, the relative magnitude of the driving forces for the possible phase transformations, does not tell us which phase is more likely to form under a given set of conditions. It is well known that according to the Ostwald step rule (Ostwald 1897) the first phase to appear will be the metastable phase which has a higher free energy than the stable phase. Here, the proof of the rule using the classical nucleation theory will be presented (Ishihara et al. 1985). The relative magnitude of the homogeneous nucleation frequencies for various phases from undercooled liquid metals and alloys is considered. The volume nucleation frequency, Iv, at a temperature T given by the classical nucleation theory can be expressed as follows:

Iv--Kvexp

-

AG*f(O)) kT

(2.6)

where Kv is a kinetic parameter, 0 is the wetting angle of the nucleus on the surface of a substrate material, and k is the Boltzmann's constant. A G* is the excess free energy of the critical nucleus and is expressed as: 16zr o'3l) 2 AG* -- -~- ( A G 2 )

(2.7)

where o" is the surface free energy between the liquid and solid, v is the molar volume, and A G is the free energy difference between the liquid and solid phases, o" may be empirically expressed as a function of the latent heat and molar volume as: o- -- 0.065AHfv -2/3

(2.8)

where A H f is the enthalpy of fusion. AG may be simply expressed by Eq. (2.1) when

ACp = O ,

as

AG = A H f ( T m - T)/Tm

(2.9)

where Tm is the melting temperature. From these equations, the following expression can be obtained: AG* = m167r(0.065)3AHf

3

T2

(2 10)

(Tin - T) 2"

The magnitude of Kv depends on the atomic transport rate in the liquid across the solidliquid interface. The value of Kv for the nucleation of the metastable and the stable phases may be considered equal since both the phases appear from the same liquid phase. So, in the case of homogeneous or fixed wetting angle nucleation, the nucleation frequency

K. N. Ishihara

14

expressed by Eq. (2.6) may be compared by calculating the differences between the value of A G* and A G .f (f refers to the metastable phase). The critical temperature, where the nucleation frequency of the metastable phase is equal to that of the stable phase, can be given as a solution to the equation AG* = AG .I

(2.11)

This equation has a significant solution under the condition that the ratio of enthalpies of fusion for stable and metastable phases, D, is less than unity:

D-

AH~f/AHf < 1.

(2.12)

When the above condition is satisfied, the solution of Eq. (2.11) is to --

A _ ~/-DA

(2.13)

Tr - x/~A where A is the ratio of the melting temperatures (T~m/Tm) and Tr is the reduced temperature (T/Tm). Since Ttm/Tm < 1, the critical temperature, to, should exist between Tm ~ and OK. A schematic illustration of AG* and A G .I vs. temperature is shown in Fig. 2.7. The condition (B < 1) should usually be satisfied because A Hf approaches A G at low temperatures and A G .I for the metastable phase is always less than that of the stable phase. Therefore, it may generally be expected that the nucleation frequency of the metastable phase is larger than that of the stable phase if the undercooling is large enough. Thus,

]! Meta~ table

o

I

0

to

A

Reduced Temperature (T/I'm) Figure 2.7. Example of the relationship of A G* for stable and metastable phases when A = 0.8 and B = 0.5. For the temperature region below t0, the nucleation of a metastable phase is more likely than that of the stable phase.

Thermodynamics and Kinetics o f Metastable Phase Formation

15

300

250 liquid --~ 3' o

200

[.~ 150 100

0

1

2

time (s) Figure 2.8. The trace of temperature during rapid cooling of Ga droplet. It is clearly shown that the sample is crystallized first into the metastable y phase (melting temperature 238 K). Then, it transforms into the metastable/3 phase (melting temperature 257 K) and finally it transforms to the stable ot phase. The free energies of these phases are shown in Fig. 2.3.

the Ostwald step rule tends to be valid for large undercooling or supersaturation. Experimental confirmation of this rule can be found in several alloy systems (Mackintosh et al. 1983, Levi et al. 1988, Cortella et al. 1993, Abe et al. 1994, Yoshioka et al. 1994, Yang et al. 1996). The sequence of phase changes during rapid cooling of Ga is shown in Fig. 2.8 (Mackintosh et al. 1983), confirming the Ostwald step rule.

2.3.2 Nucleation o f alloy phases

The question to be answered here is what would be the composition of the solid phase that nucleates first from the undercooled liquid of a given initial composition. Figure 2.9 shows schematically the possible nucleus composition at a given undercooling. For the undercooled liquid of composition Co, the formation of a solid phase with a composition between a and b reduces the free energy of the system. The magnitude of reduction in the free energy is the largest for the nucleus with the composition at c. But, the most probable nucleus composition can be explained by the relative value of the nucleation frequency. The theory of nucleation in a binary system has been presented earlier (Kamenetskaya 1967, Kelly and van der Sande 1982, Gressin et al. 1983, Ishihara, et al. 1985, Lin et al. 1988, Evans and Greer 1990, Thompson and Spaepen 1993, Ghosh 1994). According to Kamenetskaya (1967), when a solid nucleus of radius r with a concentration Cs forms

16

K. N. I s h i h a r a

e~t) Solid ~D

O Liquid

"""

i

Co

i

a Concentration

c

Figure 2.9. Schematic illustration of the free energy-concentration diagram explaining the magnitude of the driving force for nucleation of solid phase in a solid-solutionforming binary alloy system.

from a liquid of initial composition Co, the total free energy change of the system is given as"

A G --

1

3Nv

Gt(CI) + \ 3Nv

Gs(Cs) q

N

(2.14)

where N is the number of atoms in a system, Gl and Gs are the free energies of the liquid and solid phases and Ct is the liquid composition after nucleation. The mass conservation law gives the relation"

1

3--lqv- CI +

3Nv

(2.15)

The nucleus should have the composition that gives the smallest free energy increase on nucleation. Such a composition should satisfy the condition when the free energy surface shows a saddle point. Considering a binary system in which both the liquid and solid phases are ideal solutions, the surface energy can be expressed as: cr -- aA(1 -- Cs) -k- ffBCs

(2.16)

Thermodynamics and Kinetics of Metastable Phase Formation

1200

17

Co To

1000

800

e~

M 600 N

~' 400 To 3

200

O

0.0

0.2

I

i

i

0.4

0.6

0.8

P 1.0

Concentration Figure 2.10. Example of a calculated phase diagram with the ideal liquid and solid solutions. The curves SP and SQ show the most probable nucleus concentration for composition-independent surface energy and for composition-dependent surface energy, respectively.

where O"A and cr8 are the surface energies between the pure metals and the liquid. The results of calculation of the composition of the most probable nucleus as a function of temperature are shown in Fig. 2.10. This estimation agrees well with the experimental results (Fecht 1991). The area below MOSN, which is the trace of the compositions a and b in Fig. 2.9, refers to the region where the driving force for solidification from the initial liquid of composition Co is positive. The line SQ shows the trace of the saddle points. The temperature at which partitionless solidification is most probable, the intersection point of SQ and C = 0.2 in Fig. 2.10, is exactly To~3.

18

K. N. Ishihara

2.3.3 Crystal growth rate of metastable phases There have been only a few systematic studies of the growth rate. In general, the growth rate S can be expressed as a function of the undercooling by the equation (Jackson 1984)"

S-

So(AT) f

(2.17)

with So and fl (1 < fl < 2) as constants, and A T is the degree of undercooling. So and fi have no physical meaning. However, when fl of the metastable phase is larger than that of the stable phase, the growth rate of the metastable phase may become larger than that of the stable phase at a large undercooling, as shown in Fig. 2.11. Such a situation can occur when the stable phase has a faceted solid-liquid interface while the metastable phase has a molecularly rough interface (Tuckerman and Bechhoefer 1992). The experimentally investigated growth rates of the stable and metastable phase in the F e - C - X system are schematically shown (Magnin and Kurz 1988).

metastable phase

~D

o

stable phase

Tm(metastable) Tm(stable) Temperature Figure 2.11. Hypothetical relation of the growth rate for both stable and metastable phases. In this figure, fl of stable and metastable phases is 1.4 and 2.0, respectively. For SO,the value for the metastable phase is three times larger than that of the stable phase.

Thermodynamics and Kinetics of Metastable Phase Formation

19

2.4. S U M M A R Y

The thermodynamic nature of the metastable phases was discussed. The possibility of the formation of metastable phases was also discussed from the kinetic point of view-nucleation and growth of metastable phases. Furthermore, emphasis was laid on the relationship between the excess heat of mixing and the deviation of lattice parameter from Vegard's law.

2.5. L I S T O F S Y M B O L S

A a Aa B C

A Cp D E G AG A G*

AHf A Hgb

Iv Kv k N r S So T

Tm Tr V AV v fl y cr 0 f2

ratio of melting temperatures lattice constant difference of lattice constant bulk modulus concentration difference of specific heat ratio of enthalpy of fusion stored energy free energy difference of free energy excess free energy enthalpy of fusion enthalpy of grain boundary nucleation frequency kinetic parameter Boltzmann's constant number of atoms in a system radius of nucleus growth rate constant temperature melting temperature reduced temperature unit cell volume difference of unit cell volume molar volume constant Grtineisen constant surface energy wetting angle interaction parameter

20

K. N. I s h i h a r a

REFERENCES

Abe, T., Akiyama, A. and Onodera, H. (1994) ISIJ International, 34, 429. Bakker, H. (1998) Enthalpies in Alloys (TransTech Publications, Ztirich, Switzerland), p. 31. Baricco, M. (1995) Key Eng. Mater, 103, 1. Boettinger, W. J. and Perepezko, J. H. (1993) Rapidly Solidified Alloys: Processes, Structures, Properties, Applications, ed. Liebermann, H. H. (Marcel Dekker, New York), p. 17. Bosio, L., Cortes, R. and Defrain, A. (1973) J. Chem. Phys., 70, 357. Chvoj, Z. (1983) J. Non-Equilib. Thermodyn., 18, 201. Cortella, L., Vinet, B., Desr6, E J., Pasturel, A., Paxton, A. T. and van Schifgaarde, M. (1993) Phys. Rev. Lett., 70, 1469. de Boer, E R., Boom, R., Mattens, W. C. M., Miedema, A. R. and Niessen, A. K. (1988) Cohesion in Metals: Transition Metal Alloys (North-Holland, Amsterdam). Evans, E V. and Greer, A. L. (1990) in Principles of Solidification and Materials Processing, Vol. 2, eds. Trivedi, R., Sekhar, J. A. and Mazumder, J. (TransTech Publications, ZUrich, Switzerland), p. 741. Fecht, H. J. (1991) Acta Metall. Mater, 39, 1003. Gartner, E, Busch, R., Haasen, E and Bormann, R. (1996) in Metastable Phases and Microstructures, Vol. 400, eds. Bormann, R., Mazzone, G., Shull, R. D., Averback, R. S. and Ziolo, R. E (Mater. Res. Soc., Pittsburgh, PA), p. 119. Ghosh, G. (1994) Mater Sci. Eng., A189, 277. Gressin, E, Eustathopoulos, N. and Desr6, E (1983) Scripta Metall., 17, 711. Gschneidner Jr., K. A. and Vineyard, G. H. (1962) J. Appl. Phys., 33, 3444. Herlach, D. M. (1994) Mater Sci. Eng. Rept., R 12, 177. Ishihara, K. N., Maeda, M. and Shingu, P. H. (1985) Acta Metall., 33, 2113. Ishihara, K. N., Diaz de la Torre, S., Hosoe, T., Yamamoto, A. and Shingu, E H. (1993) in Processing Materials for Properties, eds. Heinen, H. and Oki, T. (TMS, Warrendale, PA), p. 667. Jackson, K. A. (1984) Mater Sci. Eng., 65, 7. Kamenetskaya, S. S. (1967) Sov. Phys. Crystallogr., 12, 71. Kelly, T. E and van der Sande, J.B. (1982) in Chemistry and Physics of Rapidly Solidified Materials, eds. Berkowitz, B. J. and Scattergood, R. O. (TMS, Warrendale, PA), p. 35. Kelton, K. E (1991 ) Solid State Phys., 45, 75. Levi, C. G., Jayaram, V., Valencia, J. J. and Mehrabian, R. (1988)J. Mater. Res., 3, 969. Lin, H. M., Kim, Y. W. and Kelly, T. E (1988) Acta Metall., 36, 2537. Mackintosh, A. J., Ishihara, K. N., and Shingu, P. H. (1983) Scripta Metall., 17, 1441. Magnin, E and Kurz, W. (1988) Metall. Trans., 19A, 1955, 1965. Massalski, T. B., Okamoto, H., Subramanian, E R. and Kacprzak, L. (1990) Binary Alloy Phase Diagrams (ASM Intemational, Materials Park, OH). Michaelsen, C., Gente, C. and Bormann, R. (1997) J. Mater. Res., 12, 1463. Murray, J. L. (1984) Metall. Trans., A15, 261. Najafabadi, R., Srolovitz, D. J., Ma, E. and Atzmon, M. (1993) J. Appl. Phys., 74, 3144. Ostwald, W. (1897) Z. Phys. Chem., 22, 22. Saunders, N. and Miodownik, A. P. (1998) CALPHAD: A Comprehensive Guide (Elsevier, Oxford). Shingu, P. H. (ed) (1992) Mechanical Alloying (TransTech Publications, Ztirich, Switzerland). Thompson, C. V. and Spaepen, E (1993) Acta Metall., 41, 2021. Tuckerman, L. S. and Bechhoefer J., (1992) Phys. Rev., A46, 3178. Tumbull, D. (1981) Metall. Trans., 12A, 695. Yang, G., Zhuang, H. and Biswas, P. (1996) NanoStructured Mater., 7, 675. Yoshioka, M., Hancock, B. C. and Zografi, G. (1994) J. Pharmac. Sci., 83, 1700. Ziman, J. M. (1972) Principles of the Theory of Solids, second edition (Cambridge University Press, Cambridge), p. 67.

Chapter 3 Rapid Solidification 3.1. 3.2.

Introduction Methods of Rapid Solidification 3.2.1 Droplet Methods 3.2.2 Spinning Methods 3.2.3 Surface-Melting Methods 3.3. Constitution and Microstructure Formation by Rapid Solidification 3.3.1 Non-Equilibrium Constitution 3.3.2 Microstructure Formation 3.4. Properties, Performance and Applications of Rapidly Solidified Materials 3.4.1 Properties and Performance 3.4.2 Applications of Rapid Solidification References Selected Bibliography

23 24 24 26 27 28 29 32 34 34 41 43 45

This Page Intentionally Left Blank

Chapter 3 Rapid Solidification H. JONES

3.1. I N T R O D U C T I O N

The term solidification simply means formation of solid so the parent material(s) can be gaseous (vapor) or liquid (e.g., melt). Thus, solidification can involve continuous or discontinuous deposition with or without a chemical reaction or passage of a solidification front through a volume of melt. The term rapid solidification is normally applied to solidification from the melt. The description "rapid" can be taken to imply a short-time interval between initiation and completion of solidification and a high velocity of propagation of the advancing solidification front. Such rapid solidification is most readily achieved by imposing a high cooling rate during solidification. This requires a thin layer, filament or droplet(s) of melt to be generated in good contact with an effective heat sink. Rapid solidification then occurs either directly, as a result of coupling between external heat extraction and the transport of latent and specific heat required to propagate the solidification front, or indirectly during the recalescence that follows nucleation of solidification at large undercooling. Rapid solidification processes such as melt-atomization, spray-deposition, splat-quenching, melt-spinning and planar flow-casting and their derivatives all generate high cooling rates during solidification and their products contain structural features formed at different levels of prior undercooling. More control over the nucleation stages of rapid solidification can be obtained by slow cooling of a volume of melt (which can be quite large) that has been freed of agents than would otherwise catalyze nucleation at low undercooling. Nucleation of rapid solidification can then be triggered by contact with a nucleating phase at predetermined levels of undercooling, though in this case the slow cool that follows completion of solidification may be insufficient to preserve to ambient temperature some of the non-equilibrium features generated by the rapid solidification step. Correspondingly, growth velocity during rapid solidification can be coupled to the velocity of a heat source (such as a laser or electron beam) traversing the surface of a block of material. In this case the heat source generates a stable melt-pool which melts new material at its leading edge as fast as previously melted material refreezes at its trailing edge. Melting and freezing front velocities increase from zero at the bottom of the melt-pool to a value related to the traverse velocity at the sample surface. The present short survey is in three sections. The first section reviews methods of rapid solidification from the melt and their characteristics. The second section reviews current knowledge of the effects of 23

H. Jones

24

rapid solidification on constitution and microstructure formation. The third section gives examples of resulting property changes and of current or potential applications.

3.2. M E T H O D S

OF R A P I D S O L I D I F I C A T I O N

These are most conveniently discussed under the headings droplet, spinning and surface melting technologies. Droplet technologies all involve generation of multiple or single droplets of melt which solidify either individually or in a state of partial or complete coalescence on a deposition surface. Spinning technologies involve stabilization of a flowing melt-stream so that it freezes as a continuous filament, ribbon or sheet in contact with a moving chill surface or heat extracting fluid. Surface melting technologies involve rapid melting at a surface followed by rapid freezing sustained by rapid heat extraction into the unmelted bulk. Cooling rates during solidification can reach 101~ K/s in cross sections as small as ~ 0 . 1 / z m with undercoolings that can attain hundreds of K, front velocities as high as 100 m/s and solidification times as short as nanoseconds.

3.2.1 Droplet methods These are all variants of the long-established technology of making lead shot by teeming molten lead through a preheated steel dish perforated by an array of equal sized holes (Johnson et al. 1976). The fundamental principle is that a pendant drop or free falling melt-stream tends to break up into droplets as a result of surface tension. This process can be intensified, for example, by impingement of high velocity jets of a second fluid, by centrifugal force at the tip of a rotating cup or disc or by an applied electric field. The standard method, known as atomization, uses high velocity gas or water jets to rapidly fragment a volume of melt into large numbers of small droplets (Fig. 3.1) (Schmitt 1979). The jets form, propel and cool the droplets that may freeze completely in flight to form powder or Metal

Metal

Connifed~ II1 Figure 3.1. Principle of spray droplet (atomization)by impingementof high velocity gas jets on to a free falling or emergent melt-stream (Schmitt 1979).

Rapid Solidification

25

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Figure 3.2. Particle size distribution for Sn-5wt%Pb alloy as a function of atomizing gas pressure (Anderson 1991).

may form individual splats or a thick deposit by impact on to a suitable substrate. Powder or shot particulate generated by such techniques displays a range of particle sizes and shapes even for a single set of operating conditions. Increasing atomizing gas pressure, for example, then increases the yield of finer particle sizes (Fig. 3.2) (Anderson 1991) while the more effective quench of water atomization leads to more irregular shapes. Although particles more or less identical in size and shape from the same powder sample can show quite different microstructures (because of different local conditions), smaller particles tend to cool more rapidly and/or undercool more prior to solidification so tend to solidify more rapidly (i.e., at higher front velocity). Splats formed from droplets of given size tend to solidify even more rapidly because of more effective heat extraction from the larger surface area they offer, especially when at least one of their surfaces is in good contact with an efficient heat sink, such as a water-cooled rotating copper drum. A spray deposit can maintain the same microstructure as the equivalent splats provided that their solidification time is sufficiently less than the time interval between deposition of successive splats at a given location on the substrate (Oguchi et al. 1990). Atomizers range in size from laboratory units with capacities of less than 1 kg per run to commercial scale facilities with capacities as large as 50,000 tonnes per year (Anon 1989). Powders destined for high performance applications tend to be atomized with inert gases or in vacuum to minimize formation of oxides or other potentially damaging inclusions. The variability of powder particle sizes, shapes and microstructures within bulk atomized powder samples, splats and spray deposits (which is acceptable for many purposes), and the need to consolidate the materials for typical engineering applications, has stimulated work on droplet methods that employ lower cooling rates and allow measurement or control of the temperature at which nucleation occurs. Such methods include emulsification techniques in which a distribution of droplets is produced within a carrier liquid which can be inert, which can act as a flux for impurities in the droplets that otherwise would activate nucleation at low undercoolings or which can trigger nucleation at a spe-

26

H. Jones

cific undercooling (Perepezko 1984). A variant uses rapid solidification to produce a fine secondary network of eutectic in a single- phase matrix. This network is then spheroidized by heat treatment above the solidus to produce a distribution of liquid droplets in the solid matrix (Prasad et al. 1984). Information regarding melting and nucleation temperatures of the droplets can then be obtained by controlled heating and cooling coupled with differential calorimetry, and the results related to microstructures formed in the solidified droplets (Kim and Cantor 1994). Undercoolings of hundreds of degrees can be generated in relatively large volumes of melt by prior removal of impurities e.g., by fluxing (Kui et al. 1984), by levitation (containerless) melting (Weber et al. 1991), or by induction remelting in contact with a water-cooled copper hearth. A recent triumph of this latter approach has been to form substantial pieces of zirconium-based metallic glasses that display ultra-high strengths and good fatigue characteristics along with extensive plasticity in the supercooled liquid range (Johnson 1996). Benefits of the related levitation melting approach have included much fundamental information on the formation of solidification microstructures at very high undercoolings and growth velocities that cannot be accessed by other rapid solidification techniques (Herlach 1994, Herlach et al. 1993). 3.2.2 Spinning methods

These derive from the simplest system in which a single melt-stream emerging from an orifice is stabilized by surface film formation or solidification before it can break up into droplets (so-called melt-extrusion or melt-spinning), to the most sophisticated in which

gas pressure

eit

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./~

~ /orifice

/ribbon

rotot/ng drum Figure 3.3. Schematicof free-jet chill-block melt-spinning (Overshott 1979).

Rapid Solidification

27

one or more melt-streams are used to make wide or composite ribbons by impingement on single or twin chill roll surfaces. In standard free-jet chill-block melt-spinning (Fig. 3.3) (Overshott 1979) the melt-stream forms a single ribbon typically 10 to 50 # m thick and a few millimeters wide by impingement on to a single chill roll rotating at a surface speed of some tens of meters per second. In planar flow-casting the distance from the nozzle orifice to the chill surface is reduced to less than ~0.5 mm to eliminate formation of a melt-pool and associated instabilities. Thin strips up to 500 mm wide have been cast in tonnage quantities by this method, either for direct application or pulverization into platelet particulate for consolidation into other product forms. Melt-extraction (Maringer and Mobley 1974), in which the rotating chill roll forms the products by direct contact with the surface of a crucible of melt or the melted extremity of the solid electrode, and its derivative melt-overflow (Boulby and Wood 1986a, b), do not involve generating a melt-stream. Continuous rapidly solidified filament with round section can be produced by the rotating water bath process in which a melt-stream emerging from a cylindrical orifice is solidified directly on entry into a rotating annulus of water contained within the lip of a rotating water bath (Ohnaka 1985). This process can also be used to generate powder particulate when the conditions result in the break-up of the melt-stream on entry into the water bath (Raman et al. 1982). Free-jet chill-block melt-spinning and meltextraction can be used to generate a flake product directly if continuity is interrupted by a series of regularly-spaced notches on the chill-block. Melt-spun ribbon or planar flow cast sheet can be pulverised, for example, by means of a blade-cutter mill (Gelinas et al. 1988, Pelletier et al. 1990) into platelets (_~0.3 mm across). These provide a more suitable feedstock for consolidation by powder metallurgy than continuous ribbon, for example, which exhibits a much lower packing density.

3.2.3 Surface-melting methods These derive from spot or traverse welding techniques and differ from them only in that the depth melted is limited to ensure that the ensuing solidification will be sufficiently rapid. In its simplest form, a single pulse or continuous traversing heat source is used to rapidly melt the surface of a block material, the unmelted bulk acting as the heat sink during the subsequent rapid solidification (Fig. 3.4) (Lux and Hiller 1972). The resulting rapidly solidified material has the same composition as the underlying parent material, although the rapid solidification may produce a different microstructure and much improved properties. A second possibility is to preplace or inject alloy or dispersoid additions at its surface so they are incorporated into the melt-zone to form a surface region of composition different from the underlying bulk. The third possibility is to melt a different material preplaced on the surface so that mixing with the underlying material is limited to the minimum required for effective bonding. All three variants offer the practical possibility of generating a more durable surface on an underlying material that is in all other respects entirely adequate for the application in view. Both nanosecond and picosecond power laser sources have been used to generate some quite spectacular non-equilibrium effects in surface melt-zones as shallow as 0.1 /xm in which cooling rates during solidification have been estimated to reach 10 l~ K/s (von Allmen et al. 1984) or more and solidification times to be as short as 10 -9 s (Spaepen 1987). Both traversing laser and electron beams

28

H. J o n e s

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Figure

have been used to treat entire surfaces via repeated incremental lateral displacement of the beam by the width of the melt-zone at the start of each new traverse (Lux and Hiller 1972). Samples with crack-free, relatively smooth treated surfaces can now be produced by appropriate control of the process parameters. The technique can also be used to develop coupling between the traversing beam and the solidification front so that the effects on resulting microstructure of systematic variations in front velocity up to -~ 1 m/s can be determined experimentally (Boettinger et al. 1984, Zimmermann et al. 1989). (See also Chapter 5 on "Laser Processing" by K.E Kobayashi in this book.)

3.3. C O N S T I T U T I O N AND M I C R O S T R U C T U R E F O R M A T I O N BY RAPID SOLIDIFICATION

Rapid solidification produces changes in constitution because the large undercoolings and front velocities involved promote formation of non-equilibrium phases and extensions in composition range of surviving equilibrium phases. Microstructural differences include changes in mode of growth and size refinement as a result of the short diffusion distances imposed. These aspects will be dealt with separately.

29

Rapid Solidification

3.3.1 Non-equilibrium constitution Non-equilibrium phases formed by rapid solidification include glasses, quasicrystalline phases, new crystalline phases and equilibrium phases with extended composition ranges. Glass formation requires that the entire or residual melt is undercooled below the glass-forming temperature before the advance of crystal formation can complete the solidification process. Glass formation is highly sensitive to composition, with pure metallic melts in particular requiring very high cooling rates during solidification to suppress the rapid advance of crystal formation. Alloy systems showing deep eutectics, however, form glasses rather more readily in the composition ranges near to these eutectics. Critical cooling rates for glass formation can be predicted by modelling the nucleation and growth of the competing crystalline (or quasicrystalline) phases. The melt will become a glass if cooling is rapid enough not to intersect the knee of the time-temperature-transformation (TTT) curve of the competing crystal forming process (Fig. 3.5) (Davies et al. 1974). In the absence of such diagrams, easy or bulk glass formation has been variously associated with ratio of glass formation to melting temperatures, large intervals between glass formation and crystallization temperatures (the so-called "supercooled liquid" range), low thermodynamic driving force for crystallization, large negative heats of mixing, simultaneous precipitation of more than one phase, large atomic size differences between the components, as well as deep eutectic formation (Peker and Johnson 1993, Massalski et al. 1981, Inoue et al. 1993). (See also Chapter 14 on "Bulk Amorphous Alloys" by A. Inoue in 2000

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30

H. Jones 9

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Figure 3.6. Glass forming ability map for Ni-based alloys under conditions of rapid solidification at ,~106 K/s. Coordinates are enthalpy of mixing A H~size to enthalpy of formation derived from the difference in atomic volume between solute and solvent (Alonso et al. 1988). ( 9 glass forming and (O) not glass forming.

this book.) Figure 3.6 (Alonso et al. 1988) shows a glass formability map for binary solute additions to nickel in which high enthalpy of formation arising from atomic size difference and large negative enthalpy of mixing are seen to correlate well with glass formation under melt-spinning conditions (cooling rate ~ 106 K/s). Formation of non-equilibrium crystalline phases by rapid solidification occurs when their nucleation and growth kinetics results in a higher rate of formation than the competing equilibrium phase. Formation of metastable (cementite) cast irons rather than graphitic cast irons at temperatures below the austenite-cementite eutectic temperature is a well-known example for relatively modest cooling rates. Possibilities for such metastable phase formation multiply manyfold at the high undercoolings and cooling rates that typify processes such as melt-spinning and splat quenching. The TTT diagram in Fig. 3.7 (Herlach et al. 1992) illustrates the competitive kinetics of formation of three contending intermetallic phases and or-A1 in undercooled Al-12 at %Mn alloy. Based on assumptions made, equilibrium A16Mn should form at cooling rates > 5 • 103 K/s, displaced by T-

Rapid Solidification

31

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32

H. Jones

phase at 5 x 103 to 2 x 105 K/s and an icosahedral phase at 2 x 105 to 2 x 106 K/s, in accord with experimental results. Extension of solid solubility by rapid solidification, again, has been widely observed, and can amount to orders of magnitude in some cases where equilibrium solid solubilities are very low (Fig. 3.8) (Midson and Jones 1982). While limited atomic size and electrochemical differences between solvent and solute appear to be necessary conditions for such solubility extension, the actual outcome is expected to be determined by conditions at the advancing solidification front. Fig. 3.9 (Boettinger et al. 1987, Kurz et al. 1988) shows the continuous increase in solute concentration with increasing solidification front velocity for ot-Ag solid solution in the Ag-15 wt%Cu alloy. The experimental data match predictions of dendrite growth modeling and approach 15 wt%Cu (segregation-free solidification, SFS) at 0.4 m/s above which plane front solidification is expected to be stabilized. The phenomenon of solute trapping is also expected to take effect in this velocity range in which the time for incremental advance of the solidification front is less than the time required for diffusional jumping of solute atoms so that instead they are engulfed by the advancing solid. Fig. 3.10 (Boettinger et al. 1984) shows the corresponding effect of varying alloy composition as well as growth velocity. The increase in velocity required for SFS with increasing Cu alloying content up to 15 wt% can be associated with incidence of plane front solidification while the decrease in this velocity above 15 wt%Cu can be attributed to control by solute trapping (Jones 1988).

3.3.2 Microstructure formation Figure 3.10 (Boettinger et al. 1984) is also an example of a solidification microstructure selection diagram (SMSD) in which dominant solidification mode is mapped on a plot which has a material parameter, such as alloy composition, as one axis and a processing parameter, such as front velocity, as the other axis. Figure 3.11 (Boettinger 1982) shows a further example in which the competition is between glass formation, formation of primary Pd3Si, y and x phases and formation of the Pd3Si-K eutectic. Modelling of the competitive growth situation implicit in such diagrams presents a challenge for solidification theory. The fundamental premise is that the solidification mode with the highest growth temperature will dominate under given conditions. Figure 3.12 (Gilgien et al. 1995) shows a comparison between experimental fact and predictions of such competitive growth modelling for Al-rich A1-Fe alloys, which form the basis of several high performance rapidly solidified aluminium alloys. The level of agreement is encouraging. While growth temperature can be crucial in determining which constituent and which mode of solidification prevails, an equally important consideration for performance and properties is the size-scale of the resulting microstructural constituent. Measurements of grain size ~.0 vs. cooling rate T (Fig. 3.13) (Jones 1991), suggest a relationship of the form ~.0 - K l J ' - n with n -- 1 and K1 ~ 107 # m K/s for Ni-5wt%A1 and Ti6wt%A1-4wt%V and ~ 106 # m K/s for Fe-6.3wt%Si and type 316 stainless steel. A prediction for pure aluminium assuming homogeneous nucleation and linear growth kinetics gave n = 0.9 and K ~ 2 x 107 (K/s) n (Boswell and Chadwick 1977). The recent observations of nanoscale dispersions of extended and non-equilibrium phases often adjacent or related to glass-forming composition ranges has also been attributed to

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operation of homogeneous nucleation (Greer 1994). Spacings of cellular or-A1 in A1Fe and A1-Cu alloys grown at 0.03 to 1 m/s show good agreement with predictions of the tip stability model of Kurz et al. (1986) and a newer array model of Hunt and Lu (1996). An example is shown in Fig. 3.14 (Lu et al. 1994). The tip stability model predicts ,k = K 2 / V 1/2, where K2 is 4Jr(DF/ATo) 1/2, in which D is the solute diffusivity in the melt, 1-" is the ratio of interfacial energy to entropy of fusion and AT0 is the alloy freezing range. Approximate analytical models predict primary dendrite spacing )~1 = K3G1/2V 1/4 with K3 = K4(AToDI') 1/4, where G is temperature gradient, K4 is 2.83k 1/4 (Hunt 1979) or 4.3/k 1/4 (Kurz and Fisher 1981) with k as a solute partition coefficient, and experimental data lend some support to this relation (Moir and Jones 1991). Secondary dendrite arm spacing )~2 is predicted to follow )~2 -- K5 i?-n with K5 "~ 50/xm (K/s) n and n = 1/3, as exemplified for or-A1 in dilute A1-Cu and AI-Si alloys in Fig. 3.15, although the experimental data for other systems show 0.25 < n < 0.5 and 16 < K5 < 160 # m K n s -n (Jones 1984). Predictions for the interphase spacing of a regular eutectic give )~ = K 6 / V 1/2 (Jackson and Hunt 1996), where 7 < K6 < 50 # m 3/2 s 1/2 for Al-based eutectics (Juarez-Islas and Jones 1987). Figure 3.16 shows eutectic spacing A vs. V for the A1-Cu eutectic, for which K6 ~ 9.4 # m 2/3 s-1/2 right up to the velocity for extinction of stable eutectic separation at V ~ 0.5 rrds (Jones 1984, Zimmermann et al. 1989).

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20

,& l

!

25

~l 30

COMPOSITION (wt % Cu)

Figure 3.10. Conditions for microsegregation free solidification of a-Ag as a function of growth velocity and alloy composition for Ag-rich Ag-Cu alloys (Boettinger et al. 1984).

3.4. P R O P E R T I E S ~ P E R F O R M A N C E

AND A P P L I C A T I O N S O F R A P I D L Y S O L I D I F I E D

MATERIALS

These will be discussed under the separate headings of Properties and Performance, and Applications (Sections 3.4.1 and 3.4.2).

3.4.1 Properties and performance Metallic glasses can exhibit high flow stresses and hardness with fracture stress approaching E/50, the theoretical limit, where E is elastic modulus. This can be combined with ductile behavior in bending (Fig. 3.17) (Davis 1976), shear and compression. Young's and shear moduli are typically 20-40% lower than in the crystalline state because the atoms in the glass are not constrained by a lattice to make self-similar displacements. Bulk modulus is however only a few percent lower in the glassy state reflecting the dense random packing which gives a density only 1-2% less than for the crystalline state. Further increments in hardness and strength along with increased stiffness can be obtained by precipitation of nanocrystals within a metallic glass matrix so as to suppress the inhomogeneous shear process (Fig. 3.17) that characterizes flow and fracture of metallic glass.

35

Rapid Solidification ,

-

25 A

E

E

I0

>~-

5

m

1

I

J

~

I

o ~, ~ ~ \ 0

'

9 Gloss

\\

/

~o~._ "\ - ~ ~ \

N

~

I

t

i

i

'i

I

l

I 15

!

a

/

\',

%~

~

,/X

\~

-

_

i '

/

g

/

/

Gio,=+x z~

9

A

/k

~2.5

> ::E I'--

tw (..9

_

D Pd 3 s i + X + ~

"

/

eut~tic

///

0.5

0.25 -

"~' J

~ I

I

l

0

r 1 5

\~

\~

o //

\ \

l

/

K + euttctie

o

/

1 ~ ~ AT. % Cu

I l0

~

I

A

I

Figure 3.11. Constituent selection diagram for Pd-Cu-Si alloys for solidification front velocities of 0.25-50 mm/s (Boettinger 1982).

1000

AI plane front '

. . . . . . .

'

- ~,~,-,~ ,-,i ,-,2\",'~ . . . . . . . .

100

.=.

10

AI cell / dend. + AI-AIsFe eul.

E u

,~TITITIT11]Trr[~

L---,

>

I-l~Tr'-'-

r

/! IIIIW%

,/111 L V

.01

.'"

-- "

AI3F e dend- +

AI-AI,Fe eat

AI den~. +" - - ~ - J / ~ " "-s "-A/F~ d e n d ~ 1 2 1 ; e ;ut" ~l-Al,,Feeut,l ~ t ~ 3 -, " 3 " ". / , ~ P ~ ' I . AI-AI,'}FeeutI,,, ,

.001 0

1

2

3

" 4

Fe-content [at%] Figure 3.12. Comparison between experimental (full lines) and predicted (broken lines) solidification microstmcture selection map for Al-rich A1-Fe alloys (Gilgien et al. 1995).

H. Jones

36

lOO

104

10 ~

106

10"

"i" K/s Figure 3.13. Grain size vs. cooling rate during solidification of [[5]] Ti-6wt%Al4wt%V, (A) Ni-5wt%A1, (~) Fe-6.3wt%Si, and ( 9 type 316 stainless steel (Jones 1991). Full line represents a prediction (Boswell and Chadwick 1977) for aluminium.

This is illustrated in Fig. 3.18 (Kim et al. 1990) for an Al-based A1-Y-Ni-M (M = Mn or Fe) metallic glass. Metallic glasses exhibit a well defined fatigue limit corresponding to ,~0.35 of the flow stress. Fatigue life above this limit tends to be shorter than for crystalline materials because absence of work hardening allows cracks to initiate at relatively low stresses. Their wear resistance can be superior to that of crystalline materials of the same hardness under sliding conditions, but tends to be inferior under abrasive wear conditions. Metallic glasses undergo considerable creep deformation at temperatures close to but below their glass transition temperature Tg compared with crystalline metals and ceramics, Newtonian viscous flow is observed at much lower temperatures and up to higher stresses. Correlations between minimum creep rate and rupture time and between applied stress, minimum creep rate and temperature are shown in Fig. 3.19 (Rama Rao and Radhakrishnan 1984, Gibeling and Nix 1978) for an F e - N i - P - B glass. The homogeneous flow behavior of metallic glasses at temperatures just below Tg can be applied to hot work them into useful shapes. The increased scattering of electrons by the disordered structure and mixed population of atomic species in a metallic glass gives rise to high electrical resistivity, two or three times that of the same composition in crystalline form. In addition metallic glasses based on ferromagnetics such as iron, cobalt or nickel exhibit excellent soft ferromagnetic behavior. Nanocrystalline counterparts can exhibit superior soft or hard ferromagnetic

Rapid Solidification

37 0.5 w t . %

Dendritic

10 ~

Cellular 10-' 10 -2

~L

2wt. % 10 ~

,~

10"

t., 10 2 ~=~ e,~

10~

$=I

101

~ o~t ~g 02

4wt.

%

!

10"2

l . . . . . . . . . . . . . . . . . . ]

10 ~

..... ~ ~, . .

6wt.%

.

10"1

"~

10"2

"..................

1 8wt.

10 ~

%

10-1

10 .2 -

10 ~

10"1

,

10 .2

Velocity (cm/sec) Figure 3 . 1 4 . a - A 1 cell spacing vs. velocity of solidification front for five A1-Fe alloy compositions compared with predictions of tip stability (dashed lines) and array model (shaded regions) (Lu e t al. 1 9 9 4 ) .

behavior depending on their constitution and alloy compositions. Excellent corrosion resistance can also be developed by appropriate levels of alloying with chromium along with additions such as Mo, W, P or B. Catalytic effectiveness can also be enhanced for the glassy state compared with the crystalline condition. Crystalline materials made by rapid solidification can also show high hardness and strength by combinations of fine grain size, presence of fine disperoids/precipitation and increased solid solution hardening, without the reduced stiffness associated with the glassy state. The Hall-Petch plot shown in Fig. 3.20 (Nagahuma et al. 1993) for rapidly solidified A1-Ni-Mn(-Zr) alloys compared with wrought unalloyed aluminium illustrates both the increased direct contribution of particle and solid solution hardening ('~500 MPa) and the indirect effect of stabilizing small grain size with an increased Hall-Petch slope

38

H. J o n e s

104

I

i V

103

i

w

i

I

i

i

,

~

I

I

I

I

i

i

=

I

I

1,

i

~'.

~2 102 l.Lm 10 I

0.1 I

0.01

10"s

I

I

.

I

!

I

I

100

"1" K/s

10 s

t

101~

Figure 3.15. Dendrite cell size ~-2 vs. cooling rate J" during solidification for A1-4 to 5wt%Cu (open points) and A1-7 to 11 wt%Si (filled points). Full line represents ),2 j"l/3 = 50 # m (K/s) 1/3 (Jones 1984).

10

'"~1

o . Eote0t,0 ~m

x

Chadwick

196

9 Cooksey .tal

01

-

A

Elliott

+ L i v i n g s t o n et al

_

o

Clark&Elliott

0.011 0.001

0.01

r~__

. 1967

1968

13 B u r d e n & Jones 1970 -

9 Zimmermann I r ill

J

1964

Ibaraki&Okamoto M9ore&

~

9

1970 1976

et al 1988 I ~ I,f

0.1

,

,

,

,! 1

V mm/s

10

100

1000

Figure 3.16. Eutectic interphase spacing ,~ vs. V for the regular (lamellar) or a-A1A12Cu eutectic (Jones 1984, Zimmermann et al. 1989).

39

Rapid Solidification

..ii.iiii.i.!.ii!iiii!i!i!!!i!!ii!i!!!ii!;ii!i!!!i!!!i!i!iiii!i!i!i!!ii!i!i!i:!!i i~ii~iiii~i!i!ii!ii~i~ii~i~ii!~i!i!i~!ii~ii~i!i~i~i~i~i~ii~ii~ii~i ;i~!i~i~!ii!ii~i!~i~ii~i~ii~ii i .....................................................................................................................................

~!i!iii!i!iiiiiiiiiiiiiiiiiiii!iiiii!iiiii!iiiiiiiiiiiiiiii iiiiliiiiiiiii!iiiiiiiiiiiii!iiiiiiiiiiiiiiiiiiiiiiiiiii iiiiii!i!iiiiiiiiiiiiiiiii!iiiiiiiiiiiiiiiiiiiiii!i

%!~i (,,.~~ ,;

5:.i% 9,,,,,r

Figure 3.17. Shear band offsets on the outer surface of a bent metallic glass ribbon (Davis 1976).

6 0 0 _ AleeY 2 N i 9 M 1 "r" 4 0 0

~

1~

ol21 Mn

@11 Fe * brittle

200 EL (..9 uJ

60, 50 I

,

I

I

I

I

1400

d3

1200

EL

-

1000

.....

3~ 800

w

2

6

~b

2'o V~

3'o

4'0

(%)

Figure 3.18. Hardness Hv, Young's modulus E, fracture strength af and elongation to fracture Ef of A1-2at%Y-9at%Ni-lat%M (M = Mn or Fe) metallic glass ribbon vs. volume fraction Vf of fcc o~-A1 nanocrystals introduced into the metallic glass matrix (Kim et al. 1990).

40

H. Jones

10%

a)

~i~i I

I

I

"i

I

O

10-5 -

10-6-

eo

10-7 102

1

103

200

b)

I

-

c,, 1 0 0

-

I

10/+ t,ls)

m

E E

' "

105

,

106

9

b

50_.22

,l

'23

,I

,

I

-2z~ -25 P= tog[~s]- O03t+T

1

-26

Figure 3.19. (a) Minimum creep rate ks vs. rupture time tr for Fe-40at%Ni-14at%P6at%B metallic glass for stress range 600-1600 MPa (temperatures: ( 9 523 K, (O) 573 K); (b) applied stress o- vs. parameter P (ks is the strain rate in s -1 , T is the test temperature in K) for the same glass. From (Rama Rao and Radhakrishnan 1984) with data from (Gibeling and Nix 1978).

(3.7 MPa mm 1/2 compared with 2.3 MPa mm 1/2 for the wrought unalloyed condition). The new alloy chemistries afforded by rapid solidification have also led to improved thermal stability, such as in the A1-Fe-V-Si (8009) alloy developed by Allied Signal. Figure 3.21 (Gilman 1990) shows that 8009 retains relatively high strength in the range 150450~ where conventional wrought A1-Cu (2219) and A1-Zn-Mg-Cu (7075) lose most of their high ambient temperature strength as a result of precipitate coarsening. The re-

Rapid Solidification 1200

1000

(3.

8O0

:

O

1

41

i

AI-Ni-Mm-Zr

(~

AI-Ni-Mm

[-]

Fujila and Taba/a

/~

Wyrzykowski and Grabski

1

Hansen

-

t-"

600

-o

400

200

0

20

I

I

I

40

60

80

d 1/2

t

100

,,,

[

120

140

/ mm

Figure 3.20. Proof strength o-0.2 vs. inverse square root of grain size d for rapidly solidified A1-Ni-Mn(-Zr) alloys ((30) compared with results for wrought unalloyed aluminium ([]Am) (Nagahuma et al. 1993).

markable resistance to coarsening of the high volume fraction of the nanoscale silicide dispersoid in 8009 thus raises maximum service temperature by some 200 K. Creep resistance and corrosion resistance are also notably improved relative to established wrought ingot alloys of comparable room temperature strength. The limit on alloying set by incidence of unacceptable ingot macrosegregation in wrought ingot processing has been relaxed by rapid solidification processing for a number of other classes of material such as silicon irons for transformers, tool and bearing steels, nickel based superalloys and copper alloys for conductor applications. It is also a route of choice for processing the new class of high performance Nd-Fe-B permanent magnet materials based on the remarkable properties of the Nd2Fe14B compound.

3.4.2 Applications of rapid solidification It is worth recalling that modem interest in rapid solidification processing stems from curiosity-driven research in the 1960s aimed at establishing its capabilities for generating new alloy phases (Duwez 1967) and size-refined microstructures (Grant 1970). There is no doubt that it has contributed much to our present understanding of what governs formation of alloy phases, their mode of occurrence in solidification microstructures and their effect on properties and performance. It has also demonstrated the development potential of a number of previously underemployed production processes such as melt-atomization and spray forming, containerless processing, melt-spinning and thin-strip casting, and surface melt-traversing, all of which are undergoing intensive appraisal. A number of new

H. Jones

42

Tem~rature (*F) 0 o,. 5OO v

100

200

300

400

500

600

700

800

900

"

70

~,~oo

50

q,)

ta,,,. W

80O9

cO 3OO

,,.'K

c

co

40 N tll

3O

N. 200 2219-T851 ~ A

7075-T651

0

9

20

.

~o

9

0

~: 113 ~

O ~

9O

6O0

8O "~

/rr

soo

~' 4oo w

o~

t~

2219-T851

7075-T651

E lOO

o.,.., ...,.

!

80

r

"-

1

I

i

"

!

'

9

9

"

1

'"I

9

"

!

9

1

I

"'

' I

.

!

' " '"

.

I

"

i

I --

|

9

1_

"

I[

"

0 80

7O 7075-T651

-~ 6 o 9

j

C:~ " c ~0 -

c: ~ I-.

l

/

30 20

60

/

c

V)

J

5O

2219-T851

40 3O 8009

20 -

10

0

----

9 !

70 -

IJJ

CD

2O N E 10

0

so

c

o

~

.~

60

c ID 3O J--.

cll I..- 2OO

-

.~

5O m

8OO9

30O

70

50

10

100 150 200 250 300 350 400 450

Temperature (*C) Figure 3.21. Tensile properties of rapidly solidified 8009 (A1-Fe-V-Si) alloy vs. test temperature compared with high strength wrought alloys 2219-T851 and 7075-T651 (Gilman 1990).

R a p i d Solidification

43

materials have been developed. These include a range of new soft (Smith 1993) and hard (Guthrie 1993) magnetic materials for application in power distribution as well as in magnetic and electronic devices. New high performance light alloys have been established for both wear resistance (Amano et al. 1989) and high temperature (Gilman 1990) performance. Materials with enhanced catalytic performance (Hashimoto 1997) and for fuel cells (Kawashima et al. 1996) have been developed. Powder metallurgy tool steels and superalloys made from prealloyed atomized powder have an established market for inserts, engineering parts and hard facings (Abraham 1991). Direct spray deposition has been employed to produce mill rolls with two or three times the life of conventionally produced rolls (Ikawa et al. 1990). Rapid solidification has been used to develop new alloys for medical implants (Wang et al. 1988) and to make improved Ag-based feedstock particulate for producing dental amalgams (Vero et al. 1992). In addition rapid solidification has provided feedstocks for other developmental technologies such as mechanical alloying, plasma and high velocity spraying, and manufacture of metal matrix composites. It has also provided cost-effective routes to make steel fibres for reinforcement of concrete and castable refractories (Edgington 1977) and aluminium flake for imparting conductivity to plastics (Holbrook 1986) as well as to form ductile brazing foils for a wide variety of precision joining applications (Rabinkin and Liebermann 1993). The newest application, that of a metallic glass golf club head (Ashley 1998) with strength and hardness twice that of cast stainless steel or titanium but of lower modulus and intermediate density, made by bulk undercooling prior to solidification, would have been inconceivable 30 years ago.

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44

H. J o n e s

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R a p i d Solidification

45

Weber, J. K. R., Krishnan, S. and Nordine, R C. (1991) J. Metals, 43(7), 8. Zimmermann, M., Carrard, M. and Kurz, W. (1989) Acta Metall., 32, 3305.

SELECTED BIBLIOGRAPHY

Ananthararman, T. R. (ed.) (1984) Metallic Glasses, Production, Properties and Applications (TransTech, Aedermannsdorf, Switzerland), 300 pp. Anantharaman, T. R. and Suryanarayana, C. (1987) Rapidly Solidified Metals (TransTech, Aedermannsdorf, Switzerland), 260 pp. Cahn, R. W. and Greer, A. L. (1996) in Physical Metallurgy, eds. Cahn, R.W. and Haasen, P., fourth edition (Elsevier Science, Oxford), p. 1723. Gilman, J. J. and Leamy, H. J. (1978) Metallic Glasses (American Society for Metals, Metals Park, OH), 350 PP. G~intherodt, H. J. and Beck, H. (1981) Glassy Metals I (Springer-Verlag, Berlin), 270 pp. Gtintherodt, H. J. and Beck, H. (1983) Glassy Metals II (Springer, Berlin, Germany), 400 pp. Herman, H. (ed.) (1981) Ultrarapid Quenching of Metals and Alloys (Academic Press, New York), 448 pp. Jacobson, L. A. and McKittnick, J. (1994) Mater. Sci. Eng. R, R l l , 355. Jones, H. (1982) Rapid Solidification of Metals and Alloys (Institution of Metallurgists, London), 83 pp. Jones, H. (1984) J. Mater. Sci., 19, 1043. Jones, H. (1991) Metals and Materials, 7, 486. Liebermann, H. H. (ed.) (1993) Rapidly Solidified Alloys: Processes, Structures, Properties, Applications (Marcel Dekker, New York), 808 pp. Luborsky, E E. (ed.) (1983) Amorphous Metallic Alloys (Butterworths, Stoneham, MA), 548 pp. Otooni, M. A. (ed.) (1997) Elements of Rapid Solidification (Springer-Verlag, Berlin), 320 pp. Sahm, P. R., Jones, H. and Adam, C. M. (eds.) (1986) Science and Technology ofthe Undercooled Melt (Martinus Nijhoff, Dordrecht, The Netherlands), 304 pp. Srivatsan, T. S. and Sudarshan, T. S. (1993) Rapid Solidification Technology: An Engineering Guide (Technomatic., Lancaster, PA), 737 pp.

These volumes provide extensive reviews of their subject areas that were comprehensive at the time of their compilation. More specialized review articles also continue to appear but are too numerous to list here. In addition, the series of international conferences that began in Brela, Yugoslavia, in 1970 and reached its ninth at Bratislava in 1996 and the tenth in Bangalore, India in 1999 contain numerous useful review articles within their Proceedings (for the ninth conference see Mater. Sci. Eng. A, 15th June 1997 issue, Vols. A 226-228, 1110 pp). The proceedings of two recent NATO workshops are also useful: Undercooled Metallic Melts: Properties, Solidification and Metastable Phases, eds. D. M Herlach et al., Mater. Sci. Eng. A, 1994, A178(1,2), 320 pp, and Science and Technology of Rapid Solidification and Processing, ed. M. A Otooni, Kluwer, Dordrecht, 1995, 390 pp). International Journal of Rapid Solidification, Vol. 1, No. 1 (1984) to Vol. 9, No. 4 (1996), is the only journal dedicated to publication of research papers in the field. It continues with its new name International Journal of Nonequilibrium Processing with Vol. 10, No. 1 in 1997.

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Chapter 4 Mechanical Alloying 4.1.

4.2. 4.3.

Introduction Nomenclature The Process of Mechanical Alloying 4.3.1 Raw Materials 4.3.2 Process Control Agents 4.3.3 Type of Mills

4.3.3.1 SPEX Shaker Mills 4.3.3.2 The Planetary Ball-Mill 4.3.3.3 Attritor Mills 4.3.3.4 Commercial Mills 4.3.3.5 New Designs 4.4. Mechanism of Alloying 4.5. Consolidation 4.6. Synthesis of Non-Equilibrium Phases 4.6.1 Solid Solubility Extensions 4.6.2 Synthesis of Intermetallics 4.6.3 Disordering of Ordered Intermetallics 4.6.4 Nanocrystalline Materials 4.6.5 Solid-State Amorphization 4.6.6 Displacement Reactions 4.7. Powder Contamination 4.8. Modeling 4.9. Industrial Applications 4.9.1 Nickel-Base Alloys 4.9.2 Iron-Base Alloys 4.9.3 Aluminum-Base Alloys 4.9.4 Magnesium-Base Alloys 4.10. Concluding Remarks References

49 5O 51 51 52 52 53 53 53 53 54 55 56 58 59 62 64 66 67 71 74 75 76 77 79 79 80 80 81

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Chapter 4 Mechanical Alloying C. SURYANARAYANA

4.1. I N T R O D U C T I O N

Mechanical alloying is normally a dry, high-energy ball-milling technique and has been employed to produce a variety of commercially useful and scientifically interesting materials. This technique was developed around 1966 by John S. Benjamin and his colleagues at the Paul D. Merica Research Laboratory of the International Nickel Company (INCO). The main goal was essentially to combine the advantages of precipitation-hardening and oxide-dispersion strengthening in a nickel-based superalloy intended for gas turbine applications. Benjamin (1976, 1990)reviewed the historic origins of the process. This process was referred to as "milling/mixing", but a patent attorney for the INCO coined the term m e c h a n i c a l alloying.

The formation of an amorphous phase by mechanical grinding of an Y-Co intermetallic compound in 1981 (Yermakov et al. 1981) and in the Ni-Nb system by ball-milling of blended elemental powder mixtures in 1983 (Koch et al. 1983) brought about the recognition that mechanical alloying is a potential non-equilibrium processing technique. Beginning from the mid-1980s, a number of investigations have been carried out to synthesize a variety of stable and metastable phases including supersaturated solid solutions, crystalline and quasicrystalline intermediate phases, and amorphous alloys (Suryanarayana 1995a). Efforts were also under way since the early 1990s to understand the process fundamentals of mechanical alloying (Maurice and Courtney 1990). Additionally, it has been recognized that this technique can be used to induce chemical (displacement) reactions in powder mixtures at room temperature or at least at much lower temperatures than normally required (Heinicke 1984, McCormick 1995). Because of all these special attributes, this simple, but effective, processing technique has been applied to metals, ceramics, polymers, and composite materials. The important attributes of mechanical alloying are listed in Table 4.1. The different facets of this technology have been reviewed periodically (Gilman and Benjamin 1983, Weeber and Bakker 1988, Koch 1989, Koch 1991, Herr 1995, Suryanarayana 1996, Lu and Lai 1998, Murty and Ranganathan 1998). The literature available till 1994 has been collected together in an annotated Bibliography (Suryanarayana 1995a). Additionally, a number of conferences are devoted to the science and technology of mechanically alloyed materials (Arzt and Schultz 1989, Yavari and Desr6 1990, Froes and de Barbadillo 1990, de Barbadillo et al. 1993). Furthermore, proceedings of the annual International Symposia on 'Metastable, Mechanically 49

50

C. Suryanarayana

Table 4.1. Attributes of mechanical alloying Production of a fine dispersion of second-phase particles Extension of solid solubility limits Refinement of grain sizes down to the nanometer range Synthesis of novel crystalline and quasicrystalline phases Development of amorphous (glassy)phases Disordering of ordered intermetallics Possibility of alloying of difficult to alloy elements Inducement of chemical (displacement) reactions at low temperatures Scaleable process

Alloyed and Nanocrystalline Materials' (ISMANAM) are regularly published in Materials Science Forum published by TransTech Publications, Switzerland (Shingu 1992, Yavari 1995, Schulz 1996, Fiorani and Magini 1997).

4.2. N O M E N C L A T U R E

Two different terms are commonly used in the literature to denote the processing of powder particles in high-energy ball-mills. Mechanical Alloying (MA) describes the process when mixtures of powders (of different metals or alloys/compounds) are milled together. Material transfer is involved in this process to obtain a homogeneous alloy. On the other hand, milling of uniform (often stoichiometric) composition powders, such as pure metals, intermetallics, or prealloyed powders, where material transfer is not required for homogenization, has been termed Mechanical Milling (MM). Some investigators have referred to MM as Mechanical Grinding (MG). Since "grinding" is normally thought of as an abrasive machining process that involves mainly shear stresses and chip formation, the term "milling" is preferred to include the more complex triaxial, perhaps partly hydrostatic, stress states that can occur during ball-milling of powders (Koch 1991). It should be realized that MA is a generic term, and some investigators use this term to include mechanical milling/grinding as well. However, we will distinguish between these two terms by using MA or MM, depending on whether material transfer is involved or not during processing. Reaction-milling (RM) is the process developed by Jangg et al. (1975). In this process the powder is milled without the aid of any process control agent (see later for its function during milling) to produce fine dispersions of oxides and carbides in aluminum (Jangg 1989). This proprietary process of Eckart-Werke in Germany achieves the dispersion of carbides by adding lamp-black or graphite during milling of aluminum. Adjusting the oxygen content via close control of the milling atmosphere (oxygen, argon, nitrogen, air, etc.) produces the oxides. Thus, the final product of milling contains the aluminum phase along with A14C3 and A1203 and these alloys are given the trade name DISPAL. Even without intentional additions of graphite, fine dispersions of A14C3 and A1203 could be produced when a process control agent is added during milling of metal powders. Thus, Singer et al. (1980) observed the presence of 10-nm size dispersoids of A14C3 and A1203

Mechanical Alloying

51

in a hot-extruded sample after milling the aluminum powder in the presence of Nopcowax 22-DSP. Milling of metal powders in the presence of reactive solids/liquids/gases (enabling a chemical reaction to take place) is now regularly employed to synthesize metal oxides, nitrides, and carbides (Avvakumov 1986, Calka et al. 1996). Thus, milling of titanium in a nitrogen atmosphere has produced titanium nitride (Calka 1991, Suryanarayana 1995b); several other compounds have also been produced in a similar way. Milling of tungsten with carbon (graphite) has produced tungsten carbide (Radlinski and Calka 1991). Milling of metal powders with boron has produced borides, e.g., TiB2 (Calka and Radlinski 1990). Another variation of milling adopted is cryomilling (Luton et al. 1989), in which the milling operation is carried out in the presence of liquid nitrogen. This process produces 2-10 nm sized aluminum oxy-nitride particles which strengthen the aluminum powder. 4.3. T H E P R O C E S S OF M E C H A N I C A L

ALLOYING

The actual process of MA starts with mixing of the powders in the right proportion and loading the powder into the mill along with the grinding medium (generally steel balls). This mix is then milled for the desired length of time until a steady state is reached when the composition of every powder particle is the same as the proportion of the elements in the starting powder mix. Sometimes the powder is milled to an intermediate state either to form metastable phases or to achieve certain desired properties. (During MM, however, the objective would be to produce certain desired properties in the material, e.g., formation of nanostructured or amorphous phases, or disordering of ordered intermetallics since no alloying/homogenization is required to occur). The milled powder is then consolidated into a bulk shape and heat treated to obtain the desired microstructure and properties. Let us now look at the individual components mentioned above. 4.3.1 Raw materials

The raw materials used for MA are widely available commercially pure powders that have particle sizes in the range of 1-200/zm. But the powder particle size is not very critical, except that it should be smaller than the grinding ball size. This is because the powder particle size decreases exponentially with time and reaches a small value only after a few minutes of milling. These powders fall into the broad categories of pure metals, master alloys, prealloyed powders and refractory compounds. The oxygen content of the commercially pure metal powders ranges from 0.05-0.2 wt.%. Therefore, it becomes important to choose high-purity powders, especially if one is interested in studying phase formation, since more often than not, the nature and amount of impurities in the system decide the nature of the final phase formed and the chemical constitution of the alloy. Dispersion strengthened materials usually contain additions of carbides, nitrides, and oxides. Oxides are the most common and these alloys are known as oxide-dispersion strengthened (ODS) materials. In the early days the powder charge for MA consisted of at least 15 vol% of a ductile compressibly deformable metal powder to act as a host or a binder. However, in recent years, mixtures of fully brittle materials have been milled successfully, resulting in alloy formation (Koch 1991). Thus, the requirement of having a ductile metal

52

C. Suryanarayana

powder during milling is no longer necessary. Accordingly, ductile-ductile, ductile-brittle, and brittle-brittle powder mixtures are milled now a days to produce novel alloys. Mixtures of solid powder particles and liquids have also been milled in recent times (Ivanov 1992).

4.3.2 Process control agents

A process control agent (PCA) is added to the powder mixture during milling, especially when the powder mix involves a substantial fraction of a ductile component. The PCAs are mostly organic compounds, which act as surface-active agents. The PCA adsorbs on the surface of the powder particles and minimizes cold welding between powder particles and thereby inhibits agglomeration. The surface-active agents adsorbed on particle surfaces interfere with cold welding and lower the surface tension of the solid material. Because the energy required for milling is equal to the product of new surface area generated times the surface tension, a reduction in surface tension results in the use of shorter milling times and/or finer powders. A wide range of PCAs has been used in practice. These are mostly organic compounds used at a level of about 1-4 wt.% of the total powder charge and include stearic acid, hexane, oxalic acid, methanol, ethanol, acetone, heptane, Nopcowax-22DSP, octane, toluene, etc. (Suryanarayana 1998). Majority of these compounds decompose during milling, interact with the powder and form compounds, and these get incorporated in the form of dispersoids inside the powder particles during milling. Thus hydrocarbons containing hydrogen and carbon, and carbohydrates containing hydrogen, carbon, and oxygen are likely to introduce carbon and/or oxygen into the powder particles, resulting in the formation of carbides and oxides which are uniformly dispersed in the matrix. These are not necessarily harmful to the alloy system since they can contribute to dispersion strengthening of the material. The hydrogen subsequently escapes as a gas or is absorbed into the metal lattice on heating or sintering. The presence of air in the milling container or milling of the powders at very low temperatures also has been shown to minimize welding, most probably due to the increased brittleness of the powder particles at such low temperatures. Metal powders (with an fcc structure) milled in a hydrogen atmosphere have been found to become brittle and not stick to themselves or the container. The choice of a PCA for milling depends on the nature of the powder being milled and the purity of the final product desired. The nature and amount of PCA used during milling determine the final powder particle size and powder yield. It should also be realized that there is no universal PCA. One has to decide on a PCA by looking at the possible interactions between the metal and the components in the PCA.

4.3.3 Type of mills Different types of milling equipment are used to produce mechanically alloyed powders. They differ in their capacity, efficiency of milling and additional arrangements for cooling, heating, etc.

Mechanical Alloying

53

4.3.3.1 SPEX shaker mills Shaker mills such as SPEX mills, which mill about 10 g of the powder at a time, are most commonly used to investigate the fundamental aspects of MA and for alloy screening purposes. In this mill, the vial containing the powder to be milled and the grinding balls, is secured in the clamp and swung energetically back and forth several thousand times a minute. The back-and-forth shaking motion is combined with lateral movements of the ends of the vial, so that the vial appears to be describing a figure-8 or infinity sign as it moves. With each swing of the vial the balls impact against the sample and the end of the vial, both milling and mixing the sample. Because of the amplitude (about 5 cm) and speed (about 1,200 rpm) of the clamp motion, the ball velocities are high (on the order of 5 m/s) and consequently the force of the bali's impact is unusually great. Therefore, these mills can be considered as high-energy variety. 4.3.3.2 The planetary ball-mill Another popular mill for conducting MA/MM experiments is the planetary ball-mill (referred to as Pulverisette) in which a few hundred grams of the powder can be milled at a time. The planetary ball-mill owes its name to the planet-like movement of its vials. These are arranged on a rotating support disk and a special drive mechanism causes them to rotate around their own axes. The centrifugal force produced by the vials rotating around their own axes and that produced by the rotating support disk both act on the vial contents, consisting of material to be ground and the grinding balls. Since the vials and the supporting disk rotate in opposite directions, the centrifugal forces alternately act in like and opposite directions. This causes the grinding balls to run down the inside wall of the vialmthe friction effect--followed by the material being ground and grinding balls lifting off and traveling freely through the inner chamber of the vial and colliding against the opposing inside wall--the impact effect. Even though the linear velocity of the balls in this type of mill is higher than that in the SPEX mills, the frequency of impacts is much more in the SPEX mills. Hence, in comparison to SPEX mills, Fritsch Pulverisette can be considered lower energy mills. 4.3.3.3 Attritor mills Attritors are the mills in which large quantities of powder (from a few pounds to 100 lb) can be milled at a time. The velocity of the grinding medium is much lower (about 0.5 m/s) than in Fritsch or SPEX mills and consequently the energy of the attritors is low. Attritors of different sizes and capacities are available. The operation of an attritor is simple. The powder to be milled is placed in a stationary tank with the grinding media. This mixture is then agitated by a shaft with arms, rotating at a high speed of about 250 rpm (Fig. 4.1). This causes the media to exert both shearing and impact forces on the material. Although much less efficient than either the SPEX or Fritsch mills, the laboratory attritor works up to 10 times faster than conventional jar mills. 4.3.3.4 Commercial mills Commercial mills for MA are much larger in size than the mills described above and can process several hundred pounds of powder material at a time. Mechanical alloying for

C. Suryanarayana

54

Water-cooled stationary tank

Gas seal

Steel ball bearings Ball mill Figure ,1.1.

Rotating impeller

Arrangement of rotating arms on a shaft in the attrition ball-mill.

commercial production is carried out in ball-mills of up to 3,000 lb capacity. The milling time decreases with an increase in the energy of the mill. For example, a process that takes only a few minutes in the SPEX mill may take hours in an attritor and a few days in a commercial mill.

4.3.3.5 New designs Several new designs of mills have been developed in recent years for specialized purposes. These include the rod mills, vibrating frame mills, and the equipment available from Dymatron, Inc., Cincinnati, OH; Super Misuni NEV-MA-8 from Nisshin Giken, Co., Ltd., Tokyo, Japan (with the ability to control the temperature of milling from very low temperatures by spraying liquid nitrogen and up to a high temperature of 300~ by electrical heating); Uni-Ball-Mill from Australian Scientific Instruments, Canberra, Australia (it is possible to control the nature and magnitude of impact of the balls in this machine with the help of magnets); HE Mill from M.B.N. srl, Rome, Italy; and Zoz Maschinenbau GmbH,

Mechanical Alloying

55

Kreuztal, Germany. Some special equipment is also designed for specific laboratory applications. The operating details and special features of these mills can be found in the respective brochures available from the manufacturers. 4.4. MECHANISM OF ALLOYING During high-energy milling the powder particles are repeatedly flattened, fractured and rewelded. Whenever two steel balls collide, some amount of powder is trapped in between them. Typically, around 1,000 particles with an aggregate weight of about 0.2 mg are trapped during each collision (Fig. 4.2). The force of the impact plastically deforms the powder particles, creates new surfaces, and enables the particles to weld together. This leads to an increase in particle size. Since in the early stages of milling, the particles are soft (if we are using either ductile-ductile or ductile-brittle material combination), their tendency to weld together and form large particles is high. A broad range of particle sizes develops, with some as large as three times bigger than the starting particles. The composite particles at this stage have a characteristic layered structure consisting of various combinations of the starting constituents. With continued deformation, the particles get work hardened and fracture by a fatigue failure mechanism and/or by the fragmentation of fragile flakes. Fragments generated by this mechanism may continue to reduce in size in the absence of strong agglomerating forces. At this stage, the tendency to fracture predominates over cold welding. Owing to the continued impact of grinding balls, the structure of the particles is steadily refined, but the particle size continues to be the same. Consequently, the interlayer spacing decreases and the number of layers in a particle increases. As mentioned earlier, the rate of refinement of the internal structure (particle size, crystallite size, lamellar spacing, etc.) is roughly logarithmic with processing time and therefore the size of the stating particles is relatively unimportant. In a few minutes to an hour, the lamellar spacing usually becomes small and the crystallite (or grain) size is refined to nanometer (1 n m - 10 -9 m or 10/k) dimensions (Fig. 4.3). The ease with which this can be achieved is one reason why MA has been extensively employed to produce nanocrystalline materials (Koch 1993, Suryanarayana 1995c, Koch 1997). After milling for a certain length of time, steady-state equilibrium is attained when a balance is achieved between the rate of welding, which tends to increase the average particle size, and the rate of fracturing, which tends to decrease the average composite particle size. At this stage each particle contains substantially all of the starting ingredients, in the proportion they were mixed together and the particles reach saturation hardness due to the accumulation of strain energy. The particle size distribution at this stage is narrow, because particles larger than average are reduced in size at the same rate that fragments smaller than average grow through agglomeration of smaller particles. The average particle size obtained at this stage depends on the relative ease with which agglomerates can be formed by cold welding, fatigue and fracture strength of composite particles, and resistance of particles to deformation. From the foregoing it is clear that during MA, heavy deformation is introduced into the particles. This is manifested by the presence of a variety of crystal defects such as dislocations, vacancies, stacking faults, and increased number of grain boundaries. The presence of this defect structure increases the diffusivity of solute elements into the matrix. Further-

C. Suryanarayana

56

Figure 4.2. Ball-powder-ball collision of powder mixture during MA (Gilman and Benjamin 1983).

more, the refined microstructural features decrease the diffusion distances. Additionally, the slight rise in temperature during milling further aids the diffusion behavior, and consequently, true alloying takes place amongst the constituent elements. While this alloying generally takes place nominally at room temperature, sometimes it may be necessary to heat treat the mechanically alloyed powder for alloying to be achieved. This is particularly true when formation of intermetallics is desired.

4.5. CONSOLIDATION

Widespread application of mechanically alloyed powders requires production of the powder in tonnage quantities and also efficient methods of consolidating the powders into bulk shapes. All the consolidation methods generally used in powder metallurgy processes can also be used for mechanically alloyed powders. However, since the powder particles in the mechanically alloyed condition are stronger and smaller (typically a few microns, even though the grain size may be only a few nanometers) in size than those used in conventional powder metallurgy operations, some special precautions need to be taken to minimize their activity and high level of interparticle friction. Conventional consolidation of powder to full density through processes such as hot extrusion and hot isostatic pressing normally requires use of high pressures and high tem-

Mechanical Alloying

57

E r N ......, t,,..

~

N

r ,,.=.=, t~

Decreasing ball-to-powder ratio

t.,.=

13.

Milling time, h Figure 4.3. Schematic showing refinement of particle and grain sizes with milling time. The rate of refinement increases and the minimum particle (grain) size reached is lower at higher ball-to-powder weight ratios.

peratures for extended periods of time. Unfortunately, this results in loss of the benefits achieved due to the metastable effects or nanostructures obtained by MA. Therefore, novel and innovative methods of consolidating the mechanically alloyed powders are required. (See also Chapter 13 on "Powder Consolidation" by Joanna R. Groza in this book.) Mechanically alloyed powders have been consolidated into useful bulk shapes by several processes. Since the mechanically alloyed powders have a high hardness, cold compaction is not an option. Furthermore, ODS alloys have been found not to densify during simple sintering. The most common method of consolidation is hot compaction followed by hot extrusion, or by direct hot extrusion at temperatures greater than half the melting point. This process could be used if one is not concerned with the loss of metastable effects, such as the crystallization of the amorphous phase in the powder. Some of the consolidation methods used for mechanically alloyed powders, both on laboratory and industrial scales, include hot isostatic pressing, powder rolling, Ceracon processing, plasma activated sintering, electrodischarge compaction, and explosive forming (shock or dynamic consolidation methods). Combination of these techniques is also sometimes used (He and Ma 1996). Some investigations were also carried out to compare the efficiencies of the different consolidation techniques in retaining the material in a nanostructured condition (Suryanarayana et al. 1997). Retention of nanostructures seems to be most efficient when consolidation is carried out at very high pressures and when the powder is exposed to these high pressures (and temperatures) for only a short period of time, e.g.,

58

C. Suryanarayana

Table 4.2 Important milestones in the development of mechanical alloying/milling 1966 1981 1983 1988 1989 1990

Development of ODS nickel-base superalloys Amorphization of intermetallics Amorphization of blended elemental powder mixtures Synthesis of nanostructures Occurrence of displacement reactions Observation of disordering of intermetallics

in shock consolidation methods. These operations are usually carried out at temperatures considerably lower than those used for conventional powders. For example, while conventional y-TiA1 powders are hot isostatically pressed at about 1,100~ (Beddoes et al. 1992), mechanically alloyed y-TiA1 powders could be consolidated to full density by hot isostatically pressing them at about 725~ (Suryanarayana et al. 1997).

4.6. S Y N T H E S I S OF N O N - E Q U I L I B R I U M P H A S E S

MA of suitable alloy compositions can result in the formation of a variety of equilibrium and non-equilibrium alloy phases. These include equilibrium and supersaturated solid solutions, stable and metastable crystalline and quasicrystalline intermediate phases, and glassy (amorphous) alloys (Koch 1991, Suryanarayana 1996). This technique is also capable of producing true alloys starting from pure elements that are not either easy to form by conventional means or sometimes even impossible to prepare, e.g., elements which are immiscible under equilibrium conditions. Rapid solidification processing (RSP) (rapid quenching of metallic melts at rates of ~106 K/s) also produces very similar microstructural and constitutional effects as MA (see Chapter 3 on "Rapid Solidification" by H. Jones in this book) and, therefore, comparisons are frequently made of the product phases obtained by these two techniques. However, MA processing is carried out entirely in the solid state. Thus, limitations imposed by phase diagrams, such as immiscibility in the liquid and solid states, do not apply to the MA process and therefore the alloying capabilities are likely to be much more extensive than by conventional ingot techniques or RSP methods. Furthermore, it should also be noted that whereas the technique of rapid solidification started as an academic curiosity in 1960 (Duwez et al. 1960, Duwez 1967) and progressed into an industrially mature technology about 20 years after the demonstration of its capabilities (Anantharaman and Suryanarayana 1987, Liebermann 1993), MA started as an industrial necessity in the late 1960s and only recently the scientific basis of MA is beginning to be understood. Some of the important milestones in the development of MMMM are listed in Table 4.2. Recent advances in understanding the "science" behind MMMM processing will be discussed first under the following headings: (1) Solid Solubility Extensions; (2) Synthesis of Intermetallics;

Mechanical Alloying (3) (4) (5) (6)

59

Disordering of Ordered Intermetallics; Nanostructures; Solid-State Amorphization; and Displacement Reactions.

The industrial applications of mechanically alloyed materials, such as oxidedispersion strengthened (ODS) alloys and others, will then be discussed towards the end of the Chapter.

4.6.1 Solid solubility extensions Solid solubility extensions have been achieved in many alloy systems by non-equilibrium processing methods such as RSP and vapor deposition. Similar results have also been reported in metal powder mixtures processed by MA. However, variation of solid solubility limits with the process parameters has not been systematically investigated. On mechanically alloying the blended elemental powder mixtures, interdiffusion between the two components occurs and solid solutions can form. This solid solubility is expected to increase with milling time as diffusion progresses and reach a saturation level, beyond which no further extension of solubility occurs. These can be considered as the solid solubility levels achieved by MA. Solid solubility levels have been generally determined from the changes in the lattice parameter values calculated from the shifts in peak positions in the X-ray diffraction patterns. There are certain difficulties associated with the solid solubility limits determined this way and, therefore, they may not be accurate for reasons discussed in an earlier article (Suryanarayana 1996). The absence of a second phase in the X-ray diffraction pattern is usually taken as an indication of formation of a homogeneous solid solution. But this may not always be true, since the appearance of the second phase in the X-ray diffraction pattern depends on the amount of the second phase, its particle size, the strain and defect content present in the material (this may shift the peak positions, sometimes asymmetrically). Therefore, caution must be exercised in interpreting the X-ray diffraction patterns from mechanically alloyed materials. Even though many of these factors have not been taken into account in many of the cases, solid solubility extensions, beyond the equilibrium values, have been reported in a number of alloy systems and some of these values are listed in Table 4.3. It should also be noted that in many cases equilibrium solid solutions (e.g., Cu-Ni, A1-Mg, Cu-A1) were synthesized by MA starting from blended elemental powders. Solid solutions (both equilibrium and metastable) form during MA because diffusion distances between particles are considerably reduced. The increased defect density and a slight rise in temperature during MA further aid diffusion. In fact, the temperature during MA seems to have a predominant effect. Increasing the temperature during MA tends to encourage decomposition of the solid solutions formed to reach thermodynamic equilibrium and oppose the effect of alloying that is possible by MA. The balance between these two opposing forces decides the final constitution of the alloy. For example, MA of a Cu-50 at. % Ag powder mixture at room temperature resulted in the formation of a homogeneous fcc solid solution. But MA at 473 K resulted in a fully decomposed two-phase

C. S u r y a n a r a y a n a

60

Table 4.3 Selected list of solid solubility extensions obtained by mechanical alloying Solvent

Solute

Maximum room temperature solid solubility (at.%) (equilibrium

by MA

(Massalski et al. 1990)) Ag A1

Cu

~0.0

100.0

Gd

~0.0

5.0

Cr

~0.0

Fe

Cd Co

Cu

Fe

0.025

5.0 4.5

Mg

1.2

23.0

Mn

0.3

18.5

Nb

--~0.0

15.0

Ru

~0.0

14.0

Ti

~0.0

Zn

3.5

3.0 ~50 6.0

C

~0.0

Cr

4.0

40

Cu

~0.0

100

Ag

~0.0

100

Hg

~0.0

70

Sb

1.2

3.7

Sn

~0.0

9.8

Ti

~0.0

AI

20.8

50

9.4

Cu

0.3

40

Mn

Co

4.0

50

Nb

AI

6.0

60

Ge

3.4

Ni

1.2

10

AI

7.2

27

C

"~0.0

12

Cr

22.5

50

Ga

10.8

50

In

~0.0

13

Nb

3.2

10

Ti

9.7

28

Ni

8.6

Ru

A1

0.7

14

Ti

A1

11.2

60

Cu

~0.0

10

Mg

~0.0

5O

Si

~0.0

8

W

Fe

2.0

40

Zr

A1

1.0

17.5

Ni

~0.0

7.0

Mechanical Alloying

61

mixture. Partially decomposed structures were obtained during milling at intermediate temperatures (Klassen et al. 1997). A similar conclusion was drawn by Ma et al. (1997) to explain the differences in the alloying behavior between two liquid immiscible systems Ag-Fe and Cu-Fe. Another interesting observation that has been made in mechanically alloyed materials is that the solid solubility achieved increases with increasing solute content in the starting blended elemental powder mixture, making it difficult to report a specific solid solubility value (Clark et al. 1995). This has been tentatively explained on the basis that the concentration gradients between the solvent and the solute are steeper at higher solute contents and that these are expected to result in increased diffusion and consequently in higher solubility levels. Neither the actual mechanism(s) for formation nor the limits for the solid solubility extensions in alloy systems obtained by MA have been well investigated. It was suggested (Schwarz et al. 1985) that the increased solid solubility of Ti in Ni in mechanically alloyed Ni-Ti powder mixtures was due to the metastable equilibrium between the ot-Ni solid solution and the Ni-Ti amorphous phase as opposed to the stable equilibrium between the ot-Ni and Ni3Ti phases. That is, the extent of supersaturation is limited to the composition where the amorphous phase starts forming. A similar conclusion was also arrived at by Murty et al. (1993). However, it has been noted in recent years that significant solid solubility extensions can be achieved in alloy systems (e.g., Ti-Mg) even when an amorphous phase did not form in them (Zhou et al. 1995). Thus, alternative explanations need to be sought to explain the formation of solid solutions and their limits of solid solubility. For example, it was suggested (Suryanarayana and Froes 1990) that formation of supersaturated solid solutions is related to the occurrence of nanocrystalline phases. The large volume fraction of grain boundaries in these materials is expected to enhance the diffusion and consequently extend the solid solubility levels in these types of systems. In support of this argument, it has been shown that the solid solubility of Mg in Ti is zero when the grain size of the titanium phase is on the order of microns, while significant solubility is reported when the titanium grain size is on the order of nanometers. Formation of (supersaturated) solid solutions by MA is more noteworthy in liquid immiscible systems. Accordingly, it has been reported that extensive solid solutions form in Cu-Co, Cu-Fe, Cu-Ta, and Cu-V systems. Gente et al. (1993) and Huang et al. (1994, 1997) suggested that the increased diffusivity due to the presence of structural defects and local stresses in nanocrystals is responsible for the formation of supersaturated solid solutions in these systems. Formation of homogeneous solid solutions is favored when the crystallite size of the constituents is reduced to below about 1-2 nm. The absolute limits of solid solubility extension in RSP alloys have been determined from thermodynamic considerations using the concept of To temperature at which, for a given composition, the solid and liquid phases have the same free energy. Supersaturated solid solutions can be obtained only when the liquid could be undercooled to a temperature below To (Boettinger and Perepezko 1993). Even though MA also is a nonequilibrium processing technique, there have not been any attempts so far to rationalize even the limited experimental data of solid solubility extensions; furthermore, the possible limits have not been defined theoretically. However, since no liquid phase is involved

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C. S u r y a n a r a y a n a

during MA, it is doubtful if the same criterion as in RSP can be used to rationalize the solid solubility extensions achieved during MA. Even though the mechanisms by which supersaturated solid solutions form by MA and RSP techniques are different, comparisons have been frequently made, with no specific conclusions. For example, solid solutions have been formed in the whole composition range in the Cu-Ag system by both the techniques. A similar situation was obtained in several other systems by rapid solidification methods, but not by MA. Solid solutions have been obtained in the full composition range in the Cu-Fe, A1Sb-InSb, and Cu-Co systems by MA but not by RSP. The levels of solid solubility achieved are also different in different systems by the two different techniques (Suryanarayana 1999).

4. 6.2 Synthesis of intermetallics Because of the intimate mixing of the constituent metal powders during MA, it has been possible to synthesize a variety of intermetallic phases at appropriate compositions in a number of alloy systems. The intermetallics synthesized include both equilibrium and metastable crystalline phases and non-equilibrium quasicrystalline phases. Only a few reports exist on the synthesis and transformation behavior of quasicrystalline phases. Icosahedral phases have been synthesized by MA in Mg-Zn-A1 (Ivanov et al. 1989a), Mg-Cu-A1, A1-Cu-Mn (Eckert et al. 1989, 1990a), A1-Cu-MnGe (Bahadur et al. 1992), A1-Cu-Fe (Nasu et al. 1992, Asahi et al. 1994), A1-Cu-Ru (Politis et al. 1989), Ti-Ni-Fe-Si, and Mg-A1-Pd systems (Takeuchi et al. 1993, Suryanarayana 1996). These icosahedral phases have been found to be similar to those produced by RSP. It has been reported (Eckert et al. 1990b) that the nature of the phase formed was different depending on the milling intensity. At a very high milling intensity, a crystalline intermetallic phase formed while at a very low intensity, an amorphous phase formed. A quasicrystalline phase formed at intermediate intensities. The formation of the crystalline intermediate phase at the highest milling intensity was attributed to the increased temperature rise during MA resulting in the crystallization of the quasicrystalline phase. In support of this, it may be noted (Bahadur et al. 1996) that an as-milled Ti-Ni-Fe-Si amorphous alloy transformed to an icosahedral quasicrystalline phase on annealing at a temperature of 1025 K. This trend is somewhat contrary to what has been observed in RSP alloys. In RSP, a stable crystalline phase formed when the melt was cooled at low cooling rates and an amorphous phase formed at the highest cooling rates. A quasicrystalline phase formed at intermediate cooling rates (Suryanarayana and Jones 1988). A similar observation of formation of metastable phases at low milling intensities and stable phases at higher milling intensities has been reported in some other alloy systems also (Gerasimov et al., 1991, Guo et al. 1994, Suryanarayana et al. 1999). The equilibrium Hume-Rothery electron compounds--fl-, y- and E-brass phasesm were synthesized by MA of pure Cu and Zn powders mixed in the proper proportions (McDermott and Koch 1986, Martelli et al. 1988, Koch 1991). A number of intermetallics in other alloy systems have also been synthesized (Koch 1991, Suryanarayana 1995a). These include aluminides, silicides, and other intermetallics. Ceramic materials such as borides, carbides, nitrides, and oxides have also been synthesized by reaction milling. In

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63

some cases, formation of the intermetallics seems to have taken place by a combustion synthesis-type reaction. Amongst these cases, there are also examples where the combustion took place only after "interrupted milling", i.e., milling of the powders for a given length of time, aging the powder at room temperature after stopping the milling, and then resuming the milling operation. An NiA1 compound was synthesized this way (Atzmon 1990). Similar results were also reported in other systems (Mukhopadhyay et al. 1996) and in fact these results compare very favorably with those obtained in the regular combustion synthesis process (Moore and Feng 1995). The synthesis of only a few metastable intermetallic phases has been reported by MA. A metastable deformation-induced martensite phase has been reported to form in mechanically alloyed Cu-Zn alloy powders (McDermott and Koch 1986). Similarly, synthesis of a bct-Fe phase on prolonged processing of iron powder in a nitrogen atmosphere has also been reported (Rawers et al. 1995). It is surprising that not many such deformationinduced metastable phases have been reported in mechanically alloyed powder mixtures in view of the fact that there are several instances where metastable phases are formed by cold working; a good example being the formation of deformation-induced metastable E (hcp) martensite in austenitic stainless steels. Recently, a metastable crystalline phase with rhombohedral crystal structure was reported to form in mechanically alloyed A1-Ge powders (Yvon and Schwarz 1993, Chattopadhyay et al. 1996). Formation of a metastable phase with a simple cubic structure was also reported in mechanically alloyed Te-Ag alloys (Chitralekha et al. 1995). Extreme caution should be exercised in identifying metastable crystalline phases produced by MA. It is possible that some of the phases formed are due to interstitial contamination, and these may be misinterpreted as new intermetallic phases. An accurate chemical analysis of the final resultant powder can, in many cases, clear up this confusion (Suryanarayana 1995b). Both disordered and ordered intermetallics (some of them metastable) have been synthesized by MA. In some cases, the intermetallics were synthesized directly by MA; but, in others an additional heat treatment was required after MA to form the intermetallic. For example, ordered intermetallics have been found to form directly on MA in Al-rich Al-transition metal systems. Some of the examples include A13Nb (Suryanarayana et al. 1994a), A15Fe2 (Mukhopadhyay et al. 1994), A13Zr (Suryanarayana et al. 1994b, Li et al. 1995a, Desch et al. 1991, 1996), (A1,X)3Zr, where X = Fe or Ni (Suryanarayana et al. 1994b, Li et al. 1995a), A13Hf, and (A1,X)3Hf where X = Fe or Ni (Li et al. 1995b). Reasons for the formation of ordered intermetallics have not been investigated in detail. It may be assumed that a phase will exist either in the ordered or disordered condition depending upon the balance between atomic disordering introduced by MA and the thermally activated reordering. The reordering is caused by the difference in energy between the ordered and disordered states. Thus, if this difference in energy is small, the alloy will exist in the disordered state whereas if it is large the alloy will be in the ordered state. It has been shown that both MA and MM produced the NiA1 phase in the ordered state while the FeA1 phase was found to be in the disordered state (Schropf et al. 1994). The ordering energy is related and scales up with the enthalpy of formation ( A H f ) and A H f values

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for NiA1 and FeA1 are 72 and 25 kJ/mol, respectively (Schropf et al. 1994), confirming the above argument. A number of intermediate phases have been synthesized in the disordered condition. This is not surprising because MA involves heavy deformation and this destroys ordering in the lattice (Stoloff and Davies 1966). The number of disordered intermetallic phases synthesized is too numerous to list here and can be found in (Koch 1991, Suryanarayana 1995a, Lu and Lai 1998, Murty and Ranganathan 1998).

4.6.3 Disordering o f ordered intermetallics It has been long known that partially ordered phases are stronger than those wholly disordered or fully ordered (because at a certain value of the long-range order parameter, S, superdislocations separate into unlinked singles). Thus, it is of interest to study the mechanical behavior of materials in various states of partial order. Disordering phenomena of ordered alloys have also been studied to understand the mechanism of disordering and also to produce the disordered material that has a higher deformability than the ordered alloy. Alloys can be disordered by irradiation (Russel 1985), rapid solidification (Koch 1988), or heavy plastic deformation (Stoloff and Davies 1966). The present section will deal with the phenomenon of disordering of ordered intermetallics by MM. An exhaustive review on this aspect has been recently published (Bakker et al. 1995a). Mechanical milling introduces high energy into the material being processed. This energy can be stored in the material as atomic disorder and/or grain boundaries, i.e., A G (milling) = A G (disorder) + A G (grain boundaries). The atomic disorder in an intermetallic can be manifested in three different ways (Bakker et al. 1995a). First, the two atomic species involved can occupy the "wrong" sublattices and this is referred to as anti-site disorder. This type of disorder was observed in a number of mechanically milled compounds with the A15 structure, e.g., Ni3A1, Ni3Si, Fe3Ge and those with the B2 structure, e.g., CoGa. Second, triple defects can be generated. In CoGa, for example, vacancies on the Co-sublattice in combination with Co anti-site atoms in a ratio of 2:1 constitute the triple defects. Third, there could also be redistribution of interstitials wherein the interstitial atoms in the octahedral sites are transferred to the tetrahedral sites, e.g., Mn3Sn2, Fe3Ge2. Additionally, grain refinement increases the grain boundary area and this also raises the free energy of the system. The sum of the energy of these two effects (disordering and creation of grain boundaries) will be the total energy introduced into the material during milling. Mechanical milling of ordered intermetallics has been shown to result in one of the three following types of transformation (Bakker et al. 1995a, 1995b): (1) formation of a solid solution of one component in the other, i.e., the terminal solid solution based on the major component, e.g., Nb3A1, V3Ga, Ni3A1, Fe3Ge, Ni2V, and NbAu2; (2) formation of an amorphous phase, e.g., Nb3Sn, NiZr, NiV2, CoZr; and (3) formation of a different phase with a complex crystal structure, e.g., Ni3Sn2, TiSi2.

65

Mechanical Alloying CoZr

1.0

tv W IW


0.5CuTi

d c/ i3AI 0.0

M E C H A N I C A L A L L O Y I N G TIME

Figure 4.4. Schematic representation of variation of long-range order parameter S with milling time.

It has been noted that upon milling, the long-range order parameter (S) in the intermetallic is gradually reduced and, in many cases, the material may become totally disordered (S = 0), e.g., Ni3A1 (Jang and Koch 1990). In other cases, S is reduced with milling time but does not reach S = 0, i.e., partial order and partial disorder co-exist, e.g., CuTi (Seki and Johnson 1990). In other cases, the S value does not decrease at all and is maintained at S = 1; but, with continued milling the material becomes amorphous, e.g., CoZr (Cho and Koch 1993). These three situations are schematically represented in Fig. 4.4. Thus, upon milling, an ordered intermetallic can transform into a disordered crystalline phase (solid solution) or an amorphous phase either with or without complete loss of long-range order. If the product is a crystalline phase, the material has an extremely fine grain size, usually in the nanometer range. Disordering of intermetallics has been studied primarily by X-ray diffraction techniques. By measuring the intensity of the superlattice reflections relative to that of the fundamental reflections, the long-range order parameter S is evaluated. Additionally, it has also been noted that the lattice parameter increases slightly (0.3-0.8%) in the disordered state (Gialanella et al. 1992). This measurement of change in lattice parameter, of course, is not possible if there is a change in the crystal structure due to disordering. Mossbauer spectroscopy techniques, superconducting properties and magnetization methods have also been employed to study the disordering phenomena (Bakker et al. 1995a). Recently, Zhou and Bakker (1996) have demonstrated that magnetic measurements are very powerful in determining the nature of disordering if one of the atoms involved has magnetic moments.

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During disordering of CoZr, Cho and Koch (1993) did not observe any decrease in the intensity of superlattice reflections (relative to fundamental lines) with milling time before amorphization occurred. It was thus concluded that the grain boundary energy in this system is high enough to drive the crystal-to-amorphous transition without any disordering occurring. Zhou and Bakker (1996) observed that magnetization in the mechanically milled Co-Zr alloys increased continuously with milling time up to about 40 h (no reduction in the relative intensity of superlattice reflections was observed) and also that the lattice parameter of the disordered nanocrystalline phase increased with milling time. From these two observations, they concluded that disordering of CoZr occurs by atomic disorder and grain refinement (i.e., creation of additional grain boundaries). Thus, a combination of techniques can provide confirmatory evidence and also an accurate account of the mechanism of disordering. Whether an intermetallic transforms to a solid solution or an amorphous phase on milling is determined by the relative free energy values of the amorphous and crystalline phases with respect to the energy stored in the intermetallic by mechanical milling. For example, the intermetallic will transform to the solid solution if the enthalpy of the amorphous state as estimated by the Miedema analysis (Bakker 1998) is higher than that of the solid solution. Thus, if AG (milling)> AG a-c, where AG a-c represents the difference in free energy between the ordered crystalline and amorphous phases, complete amorphization occurs. On the other hand, if AG (milling)< AG a-c, no amorphization occurs; instead, a solid solution would form. Similar conclusions can also be reached by observing the nature of the phase diagrams. If the phase diagram shows that the intermetallic transforms to a (disordered) solid solution before it melts, then milling of the intermetallic will result in the formation of a (disordered) solid solution. Continued milling may then produce the amorphous phase. But if the intermetallic melts congruently (irreversibly ordered or permanently ordered intermetallic), then milling of the intermetallic will produce an amorphous phase directly. This has been shown to be true in a number of instances (Suryanarayana 1995a, Bakker et al. 1995a).

4. 6.4 Nanocrystalline materials Nanocrystalline materials are single-phase or multiphase materials, the crystal size of which is of the order of a few (typically 1-100) nanometers. Because of the extremely small size of the grains, a large fraction of the atoms is located in the grain boundaries and thus the material exhibits enhanced combinations of physical, mechanical, and magnetic properties (compared to material with a more conventional grain size, i.e., > 1 #m). Nanocrystalline materials show increased strength, high hardness, extremely high diffusion rates, and consequently reduced sintering times for powder compaction. The reader is referred to some recent comprehensive reviews for full details of the processing, properties, and applications of these materials (Gleiter 1989, Siegel 1991, Koch 1993, Suryanarayana 1995c and Chapter 12, this volume). Nanocrystalline materials can be synthesized by a number of techniques starting from the vapor phase (e.g., inert gas-condensation), liquid phase (e.g., electrodeposition, rapid solidification), and solid state (e.g., mechanical attrition). The advantage of using MA

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67

for synthesis of nanocrystalline materials lies in its ability to produce bulk quantities of material in the solid state using simple equipment and at room temperature. Plastic deformation generally occurs by slip and twinning at low and moderate strain rates, while at high strain rates it occurs by the formation of shear bands, consisting of dense networks of dislocations. Transmission electron microscopy investigations have shown that plastic deformation by ball-milling, as anticipated, occurs via shear bands which are approximately 0.5 # m wide (Hellstern et al. 1989). The plastic strain in the material increases due to increasing dislocation density in the early stages of ball-milling. At a threshold dislocation density, even at only moderately elevated temperatures, the material relaxes into subgrains separated by low-angle boundaries, leading to a decrease of atomic level strain. During subsequent milling the process of high deformation/subgrain formation is repeated resulting in the subgrains becoming finer and finer, and the relative orientation of the subgrains with respect to each other ultimately becoming completely random. Once the subgrains reach a critical level of refinement, further refinement becomes virtually impossible since the stresses required for dislocation movement are enormously high due to the Hall-Petch strengthening. Thus, nanostructures with a minimum grain size are produced. In the case of fcc metals, this minimum grain size was found to scale with the melting temperature, or bulk modulus, of the given element (Eckert et al. 1992). However, in the case of bcc and hcp metals the minimum grain size attained was essentially independent of the melting temperature of the metals (Koch 1993). Nanostructure processing leads to the formation of supersaturated solid solutions and intermetallic compounds, which are otherwise difficult to obtain. For example, a supersaturated solid solution can be obtained in a system, if the grain size is in the nanometer regime; this is irrespective of the method by which the material is produced in the nanocrystalline state. Accordingly, the solid solubility of Bi in Cu is increased from the equilibrium value of <10 -4 to 4 at.% in the nanocrystalline state (Birringer et al. 1988). Similarly, formation of intermetallics such as Pd3Bi and CuEr has been achieved by nanostructure processing. As mentioned above, the enhanced solid solubility limits in mechanically alloyed materials is mostly due to the increased volume fraction of atoms residing at the grain boundaries in nanostructured materials. Since all the aspects of nanostructures are covered in Chapter 12 on "Nanocrystalline Materials" in this book, this topic will not be further covered in this chapter. 4.6.5 Solid-state amorphization

Amorphous phases are formed in suitable alloy systems by processing techniques such as rapid solidification from the liquid state, vapor-deposition, and laser-processing (Cahn 1991). In these methods, there is a change in the state of matter, i.e., a solid phase is formed from either the liquid or the vapor phase, and the high effective "quenching" rate and associated undercooling have been shown to be responsible for the amorphization process. There are also methods of amorphizing a solid without passing through the liquid (or vapor) state at any stage. Known as solid-state amorphization reactions (SSAR), these include irradiation, interdiffusion of metallic elements, pressure-induced vitrification, and mechanical deformation. Two reports of amorphization have triggered an increased research activity in this area. These are the synthesis of an Ni-Nb amorphous

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phase starting from blended elemental powders of Ni and Nb by MA (Koch et al. 1983) and amorphization of intermetallic compounds (e.g., YCo3, Y2Co7, YCo5) by MM (Yermakov et al. 1981, 1982). Since then, amorphous phases have been reported in a number of alloy systems by both MA and MM. Amorphous phases have also been obtained by MM of quasicrystalline phases produced by RSP (Zhang and Wang 1996). Amorphous phases have been synthesized by MA starting from either elemental (of two or more types) or by MM of intermetallic compound powders. A reasonably exhaustive listing of the alloy systems in which amorphous phases have been found to form by MA or MM can be found in Gaffet et al. (1993) and Suryanarayana (1999). The semimetals Si and Ge have also been partially amorphized by MM (Gaffet et al. 1993). A further point to be considered is whether, or not, there exists a difference in the way an amorphous phase forms by MA from a mixture of blended elemental powders (A and B) and by MM from the intermetallic compound (Am Bn). Two reaction routes are possible: (i) A + B --+ (AmBn)crystalline ~ (AmBn)amorphous (ii) (Am Bn)crystalline ~ (Am Bn)amorphous Instances of occurrence of amorphization by both the routes can be found in the literature (Koch 1991, Suryanarayana 1995a, 1999). In general, the end product is the same irrespective of whether the starting material is a mixture of the elements, a single intermetallic phase, or a mixture of two or more such phases. However, it has been reported that MgToZn30 can be amorphized only by MM of the intermetallic, and not by MA of the blended elemental powder mixture (Calka and Radlinski 1989). Amorphization in ordered alloys seems to follow the sequence (Jang and Koch 1990): Ordered phase ~ disordered phase (loss of long-range order) -+ fine-grained (nanocrystalline) phase ~ amorphous phase. As mentioned earlier, it appears that an intermetallic (either as the starting phase or that formed on MA of blended elemental powders) becomes amorphous when its free energy is raised above that of the amorphous phase. The number of alloy systems that show amorphization by MA or MM is too numerous to list here. Even though it appears that any alloy system can be made amorphous under appropriate conditions of milling, it may not be true. This is because a number of variables, the more important of them being milling energy, milling temperature, and impurity contamination, control the constitution of the final product. Increased milling energy (achieved by a higher ball-to-powder weight ratio, increased speed of rotation, etc.) is normally expected to introduce more strain and increase the defect concentration in the powder and thus lead to easier amorphization. However, higher milling energies also produce more heat (and higher temperatures) and this can result in the crystallization of the amorphous phase. Therefore, a balance between these two effects will determine the nature of the final product phase. Examples of the occurrence of both crystalline and amorphous phases at high milling intensities (energies) are available in the literature (Koch 1996).

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69

A few reports also exist which describe the crystallization of amorphous alloys by MA. Eckert et al. (1988) reported that if the milling intensity were too high, Zr-Ni alloys which were amorphized at a lower milling intensity would exhibit crystalline structures. This was explained as due to the temperature rise of the powder caused by the increased kinetic energy of the milling media. Similar results were also reported in other systems (Mizutani and Lee 1990, Trudeau et al. 1990, Trudeau 1994). Mechanical milling has also resulted in the crystallization of amorphous alloys in some cases. Crystallization of amorphous alloys by MA/MM, whether formed during milling or as the starting phase, has been referred to as mechanical crystallization; this behavior is not well understood. It has also been reported that a crystalline F-TiA1 phase formed on milling beyond the amorphous phase in Ti- A1 alloys (Oleszak et al. 1993). This has been, however, suggested to be due to interstitial contamination of the powder (Suryanarayana 1995b). There have also been conflicting results on the effect of milling temperature on the nature of the phase formed. Generally, a lower milling temperature was found to accelerate the amorphization process (Pathak et al. 1993, Yamada and Koch 1993). See Koch (1994) for a detailed review on the temperature effects during MA/MM. The mechanism of amorphization by MA and/or MM is not clearly understood. The early investigators (Yermakov et al. 1981, 1982) assumed that the powder particles melted because of the very high rate of plastic deformation, and consequent rise in the powder temperature. Subsequent rapid quenching of the liquid by heat conduction into the less deformed, and hence cooler, interior regions of the particles results in the formation of the amorphous phase (like in RSP). However, energy input calculations and temperature measurements suggest that the temperature rise is not large enough for the powder particles to melt. Researchers now believe that amorphization during MA is not purely a mechanical process and that a solid-state reaction similar to that observed in thin films (Schwarz and Johnson 1983) occurs during MA. During MM, destabilization of the crystalline phase is thought to occur by the accumulation of point and lattice defects, and anti-phase boundaries. Schwarz and Johnson (1983) proposed that formation of amorphous phases from two elements in the solid state requires that (i) the two elements have a large negative heat of mixing in the amorphous (liquid) state (alloy systems with a positive heat of mixing have also been amorphized by MA; but, negative heat of mixing is an essential condition for glass formation by RSP) and (ii) the two elements have very different atomic sizes, allowing one of the components to diffuse more rapidly than the other. The negative heat of mixing provides the thermodynamic driving force for the reaction and the large difference in the diffusivities is necessary for favoring the kinetics of the reaction. Thus, owing to these factors, the free energy of the amorphous phase is lower than that of the mixture of the pure crystalline compounds, thus explaining the formation of the amorphous phase by MA of a mixture of pure components. Beke et al. (1991a, b) estimated the elastic mismatch energy stored in an ordered solid solution when its long-range chemical order is destroyed. They showed that if Tc > Tm and the ratio of the elastic mismatch energy to the ordering energy is high enough, amorphization is expected to occur. This has been proved to be a valid criterion in many cases of amorphization observed by irradiation.

70

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Recently, attempts have been made to predict the ability of an alloy to become amorphous under MA/MM conditions. Since both atomic sizes and heats of formation are important in the formation of amorphous phases, Zhang (1993) plotted (RA -- R s ) / R 8 (where RA and R8 are the atomic radii of the components A and B, respectively) against A H, heat of formation of the amorphous phase, and observed that the regions where an amorphous phase forms or does not form can be separated by a straight line given by the equation:

(RA -- RB = A R ) / R B

= 0.068 A H + 0.716.

The proportion of correct predictions for the glass-forming alloys was 89.2% while that for the non-glass-forming alloys was only 71.4%. A theoretical model was also proposed in which amorphization was assumed to be realized through interstitial impurity formation during MA (Chakk et al. 1994). It was assumed that amorphization occurred when impurity atoms penetrated into interstitial sites and distorted the lattice locally. When the local distortions reached some critical value, the long-range order of the lattice was destroyed (destabilization of the crystalline phase) and an amorphous phase formed. Even though amorphization has been observed by both MA and MM, there is no quantitative data available to define the "critical" energy input (or other parameters) which can be equated to the critical cooling rate which should be exceeded to form the amorphous phase in RSP studies. Modeling of MA/MM (currently in progress in different laboratories) can provide additional information in this direction. Significant differences have been observed in the formation of metastable intermediate phases or amorphous alloys by MA/MM and RSP techniques. Formation of an amorphous phase has been reported in the Ti-A1 system by MA (due to contamination?); not possible by RSP methods. An important difference noted between these two methods is that while amorphous phases form in the vicinity of eutectic compositions by rapid solidification, they form near the equiatomic composition by MA. Furthermore, the glass formation range (the compositions range over which glasses can form) is wider in the mechanically alloyed condition than in rapid solidification. Formation of nanostructures has been reported in many alloy systems by mechanical alloying methods, but not by rapid solidification methods. But, devitrification of amorphous phases obtained by rapid solidification has led to the synthesis of nanostructure composites useful for magnetic applications (Suryanarayana 1995c, Lu 1996). Another important difference between mechanically alloyed and rapidly solidified alloys containing dispersoids appears to be in the size and distribution of the dispersoids. Figure 4.5 shows a pair of transmission electron micrographs comparing the matrix grain sizes, and size and distribution of dispersoids, in HIP compacts of Ti3Al-based alloys containing Er203 dispersoids. A Ti3A1 alloy to which 2 wt.% Er was added was rapidly solidified and then HIPed at 850~ (Fig. 4.5a). The mechanically alloyed material was Ti-25Al-10Nb-3V- 1Mo (at.%) to which 2 wt.% Er had been added and the alloy powder was HIPed at 1,000~ (Fig. 4.5b). Even after HIPing at a higher temperature, in comparison to the rapidly solidified alloy, the mechanically alloyed material showed a finer matrix

Mechanical Alloying

71

grain size, more uniform distribution of the dispersoids, absence of large dispersoids at the grain boundaries, and absence of dispersoid-free zones near the grain boundaries. Majority of the above-mentioned and other differences between materials processed by these two techniques could be traced to the fact that rapid solidification is a liquid-tosolid transformation while MA is a fully solid-state transformation.

4. 6.6 Displacement reactions It was recognized in 1989 that MA may be used as the basis of a chemical refining process (Schaffer and McCormick 1989) by demonstrating that the pure metal Cu can be synthesized when CuO and Ca were ball-milled together at room temperature. Simultaneous reduction of CuO and ZnO by Ca has also been shown to result in the formation of flbrass. This method of metal synthesis directly from their oxides or chlorides has now been extended to the synthesis of a number of metals, e.g., Zn, Ti, Zr, Ta, Gd, Er, Sm, V, W and some intermetallics, e.g., SmCos. Results of the Australian group have recently been summarized in a review article (McCormick 1995). Other groups also have now taken up this type of work using MA. Most of the reactions studied to date have been displacement reactions of the type: MO + R ---~ M + R O where a metal oxide (MO) is reduced by a reductant (R) to the metal (M). Metal chlorides also have been reduced to metals this way. The displacement reactions studied by MA are characterized by a large negative free energy change at room temperature and are therefore thermodynamically feasible at room temperature. However, commercial operations by the pyrometallurgical techniques are conducted at elevated temperatures to overcome the kinetic barriers and achieve sufficiently high reaction rates. Mechanical alloying can provide the means to substantially increase the reaction kinetics of the displacement reactions. This is because during MA, the repeated fracturing and welding of powder particles increases the area of contact between the particles due to a reduction in particle size and also allows fresh surfaces to come into contact repeatedly. This allows the reaction to proceed without diffusion through the product layer. Additionally, the high defect densities induced by MA/MM accelerate the diffusion process. As a consequence, these reactions will now occur at room temperature. If a reaction cannot occur at room temperature, the particle refinement and consequent reduction in diffusion distances (due to microstructural refinement) can at least significantly reduce the reaction temperatures. The product of the displacement reactions normally consists of two phasesmthe metal (or a compound) and the oxide or chloride associated with the reductant. The removal of the unwanted reaction by-product can be difficult due to the high reactivity of the metal phase associated with nanocrystalline grain sizes and intermixing of the phases induced by the MA process. The by-product phase can be easily removed if the metal particles are embedded in a continuous matrix of the byproduct phase. Removal of the byproduct is achieved by leaching the product mixture in a dilute acid or hot water, or by vacuum

C. Suryanarayana

72

RS+HIP'ed

850~

,,/!~; ~~

" ,

[ .....

' MA + HIP'ed

1000~

B

b

Figure 4 5 Transmission electron micrographs showing the difference in the m a trix grain size, and size and distribution of dispersoids in rapidly solidified (RS) and mechanically alloyed (MA) materials

t

Mechanical Alloying

TIC14-

Kroll Process +Mg (!) @ 1000" C

73

Ti sponge Melt

_., %.

Ingot Casting PREP

Ti powder Figure 4.6. Schematic illustration showing the route by which titanium alloy powder can be synthesized directly from its tetrachloride by mechanical alloying with a reducing agent.

distillation. The use of carbon or hydrogen as the reductant produces gaseous C0 2 or water vapor as the byproduct and obviates the need for leaching/distillation. The displacement reactions have now been studied to synthesize metals, alloys, intermetallics, and rare-earth permanent magnet alloys. The rare-earth alloys synthesized include Nd2Fel4B, Sm2Fel7N2.6, Sm2Fel4Ga3C2, SmCo5, Sm2Co17, SmFeTi and SmCoFe. The product of MA in these cases is usually an amorphous phase that needs to be heat treated to crystallize the hard-magnetic phase. An important advantage of the MA process is that nanocrystalline microstructures are developed and these result in remanence-enhanced magnetic materials. Since with remanence enhancement there is no need for an alignment process, significant potential exists for achieving cost savings and simplification of the manufacturing process. The mechanically driven displacement reaction offers a number of advantages over conventional metal processing techniques. First, it enables the reduction of a number of oxides and halides to pure metals at room temperature, thus effecting energy savings. Second, if a number of components are reduced simultaneously, then it is possible to produce an alloy without first having to convert the oxides to pure metals and then to the desired alloy. Third, for powder metallurgy applications it allows the direct formation of powder product without first having to manufacture the bulk alloy and then convert it to powder form. Thus, a number of high-temperature processes can be combined into one single room-temperature process with the potential for significant cost savings (Fig. 4.6) (Schaffer and McCormick 1992).

74

C. S u r y a n a r a y a n a

4.7. P O W D E R C O N T A M I N A T I O N

A major concern in the MA process is the impurities that get into the powder and contaminate it. The small size of the powder particles, availability of large surface area, and formation of fresh surfaces during milling all contribute to the contamination of the powder. Thus, it appears as though powder contamination is an inherent drawback of the technique, unless special measures are taken to avoid/minimize it. As mentioned earlier, in the early stages of MA, the metal powder coats the surface of the grinding medium and the inner walls of the container. This was expected to have prevented contamination of the powder and so no attention was paid to the problem of powder contamination. However, when different results were reported by different groups of researchers on the same alloy system, it was recognized that contamination could be a serious problem. This problem appears to be ubiquitous and is now encountered in many investigations, especially when reactive metals such as titanium and zirconium are milled. The magnitude of contamination appears to depend on the time of milling, the intensity of milling, the atmosphere in which the powder is milled, and the difference in strength~ardness of the powder and the milling medium. Whereas, 1-4 wt.% Fe has been found to be normally present in most of the powders milled with the steel grinding medium, amounts as large as 7 wt.% (20 at.%) Fe in a W-C mixture milled for 310 h (Wang et al. 1997), 13 wt.% (33 at.%) Fe in pure W milled for 50 h in a SPEX mill (Yang et al. 1993), and 60 ac% Fe in a W-5 at.% Ni alloy milled for 60 h in a SPEX mill (Courtney and Wang 1992) were also reported. These are very high levels of contamination. Iron contamination has also been observed in other alloy systems. For example, the presence of 16 at.% Fe was reported in a Ta-30 at.% A1 alloy powder milled for 400 h in a ball-mill (E1-Eskandarany et al. 1991). Similarly, large amounts of oxygen (up to 36.5 at.%) and nitrogen (up to 22.6 at.%) have also been reported to be present in AI-Ti powders milled for 400 h in a low energy ball-mill (Saji et al. 1992). Contamination of metal powders can be traced to (i) chemical purity of the starting powders; (ii) milling atmosphere; (iii) milling equipment; and (iv) the process control agents added to the powders. Contamination from source (i) can be either substitutional or interstitial in nature, while contamination from source (ii) is essentially interstitial in nature and that from (iii) is mainly substitutional, even though carbon from the steel equipment can be an interstitial impurity. Contamination from the PCA leads to interstitial contamination. Several attempts have been made in recent years to minimize the powder contamination during MA. One way of minimizing the contamination from the grinding medium and container is to use the same material for the container and grinding medium as the powder being milled. In this case also there will be wear of the grinding medium and the debris gets incorporated into the powder. Thus, even though there is no contamination, the chemistry of the final powder will be quite different from the starting powder; the metallic content (from the container and balls) would be higher than in the initial powder. This can be compensated for if we know how much of the metallic content is increased. The above solution may be possible in some cases; but it is difficult in many cases due to the non-availability of the special grinding medium and containers.

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75

The problem of milling atmosphere is serious and has been found to be the major cause of contamination in many cases. In fact, it has been observed that if the container is not properly sealed, the atmosphere surrounding the container, usually air (containing nitrogen and oxygen), leaks into the container and contaminates the powder. Thus, when reactive metals like titanium and zirconium are milled in improperly sealed containers, the powders are contaminated with nitrogen and oxygen. Attempts have been made to improve the container seal integrity to prevent the outside atmosphere leaking inside. Use of high-purity argon (99.998%) atmosphere and improvements of the seal quality resulted in the processing of high-quality titanium alloy powder with as little as 100 ppm oxygen and 15 ppm nitrogen (Goodwin and Ward-Close 1995). This process, however, may not be economically viable and hence may not be feasible on an industrial scale.

4.8. M O D E L I N G

From the description of the MA process it is easy to realize that MA is a process involving a number of both independent and interdependent variables. Like in any other process, modeling of MA is also carried out to identify the salient factors affecting the process and establish process control instrumentation. By modeling the process effectively, it is possible to bring down the number of actual experiments to be conducted to optimize the process and achieve a particular application. The number of variables involved in the MA process is very large. For a particular alloy system the variables include the type of mill, intensity of milling, type of milling media, ball-to-powder weight ratio, the atmosphere under which the powder is milled, purity of the powders, milling time, milling temperature, and nature and amount of the PCA used. All these have a significant effect on the constitution of the powder. Even on a local scale, the nature of impacts between two balls, the frequency of impacts, the amount of powder trapped between two balls during a collision could vary from point to point. Thus, modeling the MA process is a difficult task. In spite of this, some attempts have been made in recent times (Courtney 1994, Courtney and Maurice 1996, Magini and Iasonna 1995, Magini et al. 1996, Abdellaoui and Gaffet 1996) and moderate success has been achieved in modeling the mechanics of the process. From the actual experiments conducted, attempts have also been made to correlate the phases formed with the process parameters during milling. But, we are far from a situation where we can predict the final chemical constitution of the powder (type and description of phases formed). As Courtney and Maurice (1996) succinctly put it: "It is important to appreciate what can be expected of even the most successful efforts of this kind. 'Absolute' predictions are unlikely; this is exacerbated in the case of MA by its inherently stochastic nature. In addition, models require input 'data'; material properties (often not known to precision, particularly at the high strain levels to which powder is often subjected during MA) and process characteristics (collision frequency and velocity, also seldom known a priori). Thus, realistic goals of process modeling are to correctly predict general trends, and perhaps even to predict resulting properties/dimensions within an order of magnitude. The benefits of successful models lie not in their abilities to predict outcomes, but in that they

76

C. Suryanarayana

Table 4.4 Room temperature and elevated temperature mechanical properties of commercial ODS nickel- and iron-base superalloys Alloy

MA 6000

MA 754

MA956

Test

0.2% YS

UTS

temperature o f (~

MPa

MPa

% El.

% RA

RT

1220

1253

7.2

6.5

871 (1600)

675

701

2.2

4.6

982 (1800)

445

460

2.8

1.9

RT

586

965

21

33

871 (1600)

214

248

31

58

982 (1800)

169

190

18

34

RT

517

655

20

35

97 69

100 72

12

30

1000 (1832) 1100 (2192)

MA 6000:15 Cr, 4.5 A1, 0.5 Ti, 2.0 Mo, 4.0 W, 2.0 Ta, 1.1 Y203, balance Ni. MA 754:20 Cr, 0.3 A1, 0.5 Ti, 0.6 Y203, balance Ni. MA 956:20 Cr, 4.5 AI, 0.5 Ti, 0.5 Y203, balance Fe.

help identify critical process and material variables, and reduce the amount of testing needed for process optimization".

4.9. I N D U S T R I A L A P P L I C A T I O N S

The major industrial applications of mechanically alloyed materials have been in the areas of thermal processing, glass processing, energy production, aerospace, and other industries. These applications are based on the oxide-dispersion strengthening effect achieved in mechanically alloyed nickel-, iron- and aluminum-based alloys. Mechanically alloyed materials are strong both at room and elevated temperatures (Table 4.4). The elevated temperature strength of these materials is derived from more than one mechanism. First, the uniform dispersion (with a spacing of the order of 100 nm) of very fine (5-50 nm) oxide particles (commonly used are Y203 (yttria), ThO2 (thoria), and La203 (lanthana)), which are stable at high temperatures, inhibit dislocation motion in the metal matrix and increase the resistance of the alloy to creep deformation. Another function of the dispersoid particles is to inhibit the recovery and recrystallization processes, because of which a very stable large grain size is obtained; these large grains resist grain rotation during high temperature deformation. A stable large grain size can also be obtained by secondary recrystallization mechanisms. Second, the very homogeneous distribution of alloying elements during MA gives both the solid-solution strengthened and precipitation-hardened alloys more stability at elevated temperatures and overall improvement in properties. Mechanically alloyed materials also have excellent oxidation and hot corrosion resistance. The increased resistance to oxidation-sulfidation attack is due to the homogeneous distribution of the alloying elements and the improved scale adherence due to the dispersoid itself (Hack 1987, Fischer et al. 1991).

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77

Temperature, OF 1600

90 ~E

&

-~

50

--~

40

"-~

30

~

20

:3

0 0

o

i

" DS Mar-M200 + Hf

A SC PWA 454 o TD-Nickel

70 60

2200

9MA 6000

8O '

x:"

1900

=,,.,..

10 0 700

800

900

1000

1100

1200

Temperature, ~

Figure 4.7. Comparison of 1000-h specific rupture strength of INCONEL MA6000 with dispersion strengthened Mar-M200 + Hf, TD-Ni, and single-crystal PWA 454. Courtesy of Inco Alloys International.

4.9.1 Nickel-base alloys The most significant advantage of oxide-dispersion strengthened superalloys is the improvement in the stress rupture properties. Figure 4.7 compares the specific rupture strength (strength/density) for a 1,000-h life as a function of temperature for several nickel-base superalloys used for turbine blade applications. Mar-M200 is a nickel-base alloy containing by weight percent 9.0 Cr, 5.0 A1, 2.0 Ti, 12.0 W, 10.0 Co, 1.0 Nb and 1.8 Hf, while PWA454 is a nickel-base alloy containing 10.0 Cr, 5.0 A1, 1.5 Ti, 12.0 Ta, 4.0 W and 5.0 Co, and TD Ni is nickel containing 2.0 wt.% ThO2. It is clear from this figure that the MA 6000 alloy can maintain a given stress for a much longer time than a conventional alloy for similar vane applications. This is mainly due to the benefits of the combined strengthening modes in the mechanically alloyed material. The mechanically alloyed superalloys are considered mainly for three groups of applications--gas turbine vanes, turbine blades, and sheets for use in oxidizing/corrosive atmospheres. The largest use of MA 754 is as vanes and bands for aircraft gas turbine en-

78

C. Suryanarayana

~

R

Figure 4.8. Some typical gas turbine engine components made out of INCONEL alloy MA 754. Courtesy of Inco Alloys International.

gines (Fig. 4.8). For applications requiring good resistance to thermal fatigue, such as gas turbine vanes, Inconel MA 754 is given a strong texture. The majority of the grains are aligned so that their (100) axes are parallel to the principal working direction and along the length of the bar. Such texture results in low modulus of elasticity (149 GPa) in the longitudinal direction. The low modulus improves resistance to thermal fatigue by lowering stresses for given thermal strains. The alloy MA 758 is used in the glass industry for high-temperature components requiring both elevated temperature strength and resistance to extremely corrosive molten glass. It is also used for internal combustion engine components, mainly in critical fuel injection parts. One novel industrial application of alloy MA 754 is a high-temperature atmosphere-circulation fan in a "floating" furnace design being commercialized in Japan. Large rounds are used for the hub and plate material is used for blades of the fan, which operates at temperatures over 1,100~ The MA 6000 alloy is a more complex alloy developed as a blade material for advanced gas turbines. It is used for first and second stage turbine vanes and blades machined form solid bar. Unlike cast alloys, MA 6000 exhibits nearly flat rupture-life curves at high temperatures due to the combination of oxide dispersion strengthening and high grainto-width ratios (typically > 10 to 1). Because of its composition, MA 6000 has excellent resistance to oxidation and sulfidation. The characteristics of this alloy allow blade cooling to be reduced or eliminated as the metal temperature can be increased by 100~ or more in engines where the stresses are medium or low.

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79

4.9.2 Iron-base alloys The mechanically alloyed iron-base superalloys combine the high-temperature strength and stability of oxide dispersion strengthening with excellent resistance to oxidation, carburization, and hot corrosion. These alloys are suitable for use in gas turbine combustion chambers. Incoloy alloy MA 956 (20 Cr, 4.5 A1, 0.5 Ti, 0.5 Y203, balance Fe) is particularly well suited for use in heat processing applications. For example, vacuum furnace fixtures made of MA 956 have shown excellent durability and are able to compete with wrought molybdenum, which is also used in these applications. In comparison to molybdenum, MA 956 is about 30% lower in density, providing weight savings and cost advantages. Furthermore, since MA 956 has a lower vapor pressure than molybdenum, it will not coat the inside of the vacuum chamber or the parts being heat treated. Thus, MA 956 rods, flats, and sheets are used in numerous atmosphere and vacuum furnace applications including muffles, baskets, trays, and thermowells. Alloy MA 956 in tubing form has also been used for high temperature, severe service applications such as coke injection lance pipes in steel making. The alloy MA 956 is also being used in glass processing industry because of its resistance to attack by molten glass. Because of this corrosion resistance, the alloy is being evaluated for applications such as firing-kiln rollers, muffle tubes, and furnace racks. Other applications include molten-glass resistance heaters, thermocouple protection tubes, glass-processing components used in nuclear waste disposal and the bushings used to make single and multistrand fibers. More recently, MA 957 (14 Cr, 1.0 Ti, 0.3 Mo, 0.25 Y203, balance Fe) has been evaluated for use as the fuel cladding in fast neutron, breeder reactors. Conventional austenitic alloys are unsuitable for this application due to the dimensional swelling phenomenon caused by the high neutron fluxes. The mechanically alloyed materials are also being evaluated for heat exchanger components in high-temperature gas-cooled reactors.

4.9.3 Aluminum-base alloys The success of mechanically alloyed superalloys led to the development of dispersionstrengthened aluminum alloys. Since an aluminum oxide layer is always present either on the surface of the powder particles at the start of processing or during milling, its incorporation into the alloy contributes to significant improvements in the properties of the alloy. Furthermore, since aluminum is a ductile metal, PCAs are added to assist in minimizing cold welding during processing. Aluminum carbides are formed during MA by the decomposition of the PCA. Both the oxide or carbide type dispersions are about 30-50 nm in size and stabilize the ultrafine grain size. This results in 50% increase in strength, higher fracture toughness and improved resistance to stress corrosion cracking and fatigue crack growth of the mechanically alloyed materials. IncoMAP alloy AL-9052 (4.0 Mg, 1.1 C, 0.6 O, balance A1) has a density 5% less than that of conventional agehardenable aluminum alloy of comparable strengths such as 2024. With its combination of lightweight, high strength, and corrosion resistance, IncoMAP alloy AL-9052 is evaluated for aerospace applications where marine corrosion is also a factor. Addition of lithium to mechanically alloyed aluminum alloys has produced an ultra lightweight alloy IncoMAP alloy AL-905XL (4.0 Mg, 1.3 Li, 1.1 C, 0.6 O, balance

80

C. Suryanarayana

A1). Its density is 8% lower and stiffness 10% greater than the age-hardenable conventional alloy 7075-T73 of comparable strength. The excellent combination of the properties makes this alloy very attractive for airframe applications. In particular, the freedom from age-hardening treatments makes it possible to produce forgings and heavy sections with homogeneous metallurgical structures.

4.9.4 Magnesium-base alloys A useful application of the MA technique was in the production of supercorroding magnesium alloys that operate as short-circuited galvanic cells to corrode (react) rapidly and predictably with an electrolyte such as seawater to produce heat and hydrogen gas (Black 1979). Such an alloy system is suitable as a heat source for warming deep-sea divers, as a gas generator to provide gas for buoyancy, or as a fuel in hydrogen engines or fuel cells. Magnesium-base alloys containing Fe, Cu, Co, Cr, or Ti have been evaluated for such applications. The Mg-5 to 20 at.% Fe alloy is ideal because of its extremely fast reaction rate, high power output, and the high percentage of theoretical completion of the actual reaction. For corrodable release links an alloy with a slower reaction rate, such as Mg-5 at.% Ti is useful. There have also been a number of investigations in recent years to use the MA processing technique to produce metal hydrides. This is because metal hydrides are materials for safe storage of hydrogen and they can store hydrogen with a higher volume density than liquid hydrogen. However, these are sensitive to surface oxidation and hence can be a limiting factor in their commercial utilization. Nanocrystalline hydrides have a high density of defects and interfaces that could enhance diffusion, and therefore nanocrystalline intermetallics would not require activation treatments at high temperatures and pressures after exposure to air (Tessier et al. 1996). Several magnesium-base and iron-base intermetallics are being evaluated for this application. Figure 4.9 summarizes the applications of mechanically alloyed materials. The range of these applications is continuing to increase.

4.10. C O N C L U D I N G R E M A R K S

Mechanical alloying is a simple, elegant, and useful processing technique that continues to attract the serious attention of researchers. Even though the technique was originally developed to produce ODS superalloys, the synthesis of novel crystalline and amorphous phases has spurred lots of research investigations in recent years. It is estimated that so far about 4,000 research/review papers have been published in this area, with nearly 500 per year during the last 2-3 years. Both equilibrium and non-equilibrium phases and commercially useful materials could be synthesized starting from elemental powders. This also appears to be an economical process with vast potential. One of the greatest advantages of MA is in the synthesis of novel alloys that cannot be made by any other technique, such as alloying of normally immiscible elements. This is because MA is a completely solid-state processing technique and therefore limitations imposed by the phase diagrams do not apply here.

Mechanical Alloying

81

I~ STRENGTHENEDALLOYS ) ~_Ni-, Fe-, AND AI-BASE . . ~

Figure 4.9. Typical applications of the mechanical alloying process.

On an industrial scale, this is an accepted process. The ODS alloys have a higher temperature operating capability at the same stress or increased load-bearing capability at the same temperature than alloys without an oxide dispersion. The ODS alloys continue to find applications in a wide variety of industries. The two production facilities of Inco Alloys International have a combined annual powder capacity approaching 300,000 kg. The yield of the final product varies greatly with the product form and size, but the final wrought capacity is over 200,000 kg. Mechanical alloying involves powders with very small sizes and these should be handled with caution and care. Because of the large surface area, they are highly reactive and can be pyrophoric and can cause health problems when inhaled. Precautions should be taken not to open the powder to atmosphere immediately after milling since this can lead to oxidation of the powders and in some situations they can even catch fire.

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Avvakumov, E. G. (1986) Mechanical Methods of Activation of Chemical Processes (Nauka, Novosibirsk, Russia). Bahadur, D., Singh, K. and Roy, M. (1992) Mater Sci. Eng., A154, 79. Bahadur, D., Dunlap, R. A. and Foldeaki, M. (1996) J. Alloys & Compounds, 240, 278. Bakker, H. (1998) Mater Sci. Foundations, 1, 1. Bakker, H., Zhou, G. E and Yang, H. (1995a) Prog. Mater Sci., 39, 159. Bakker, H., Zhou, G. E and Yang, H. (1995b) Mater Sci. Forum, 179-181, 47. Beddoes, J. C., Wallace, W. and de Malherbe, M. C. (1992) Inter J. Powder Metall., 28, 313. Beke, D. L., Bakker, H. and Loeff, P. I. (1991 a) Acta Metall. Mater., 39, 1267. Beke, D. L., Loeff, E I. and Bakker, H. (1991b) Acta Metall. Mater., 39, 1259. Benjamin, J. S. (1976) Sci. Amer., 234(5), 40. Benjamin, J. S. (1990) Metal Powder Rept., 45, 122. Birringer, R., Hahn, H., H6fler, H. J., Karch, J. and Gleiter, H. (1988) Defect Diff. Forum, 59, 17. Black, S. A. (1979) Report of Civil Eng. Lab. (Navy, Port Hueneme, CA), 40 pp. Boettinger, W. J. and Perepezko, J. H. (1993) in Rapidly Solidified Alloys: Processes, Structures, Properties, Applications (Marcel Dekker, New York), p. 17. Cahn, R. W. (1991) in Glasses and Amorphous Materials, Vol. 9 of Materials Science and Technology--A Comprehensive Treatment, ed. Zarzycki, J. (VCH Verlagsgesellschaft GmbH, Weinheim, Germany), p. 493. Calka, A. (1991) Appl. Phys. Lett., 59, 1568. Calka, A. and Radlinski, A. P. (1989) Mater Sci. Eng., A l l & 131. Calka, A. and Radlinski, A. P. (1990) J. Less-Common Metals, 161, L23. Calka, A., Nikolov, J. J. and Williams, J. S. (1996) Mater Sci. Forum, 225-227, 527. Chakk, Y., Berger, S., Weiss, B.-Z. and Brook-Levinson, E. (1994) Acta Metall. Mater., 42, 3679. Chattopadhyay, K., Wang, X.-M., Aoki, K. and Masumoto, T. (1996) J. Alloys & Compounds, 232, 224. Chitralekha, J., Raviprasad, K., Gopal, E. S. R. and Chattopadhyay, K. (1995) J. Mater Res., 10, 1897. Cho, Y. S. and Koch, C. C. (1993) J. Alloys & Compounds, 194, 287. Clark, C. R., Suryanarayana, C. and Froes, E H. (1995) in Advances in Powder Metallurgy & Particulate Materials - 1995, Compiled by M. Phillips and J. Porter (MPIE Princeton, NJ), Part I, pp. 1-135. Courtney, T. H. (1994) in Reviews in Particulate Materials, Vol. 2, eds. Bose, A., German, R. M. and Lawley, A. (Metal Powder Industries Federation, Princeton, NJ), p. 63. Courtney, T. H. and Maurice, D. R. (1996) Scripta Mater., 34, 5. Courtney, T. H. and Wang, Z. (1992) Scripta Metall. Mater., 27, 777. deBarbadillo, J. J., Froes, E H. and Schwarz, R. B. (1993) Mechanical Alloying for Structural Applications (ASM International, Materials Park, OH). Desch, E B., Schwarz, R. B. and Nash, E (1991) J. Less-Common Metals, 168, 69. Desch, E B., Schwarz, R. B. and Nash, P. (1996) Scripta Mater., 34, 37. Duwez, E (1967) Trans. Amer Soc. Met. Quart., 60, 607. Duwez, E, Willens, R. H. and Klement, W. (1960) J. Appl. Phys., 31, 1136. Eckert, J., Schultz, L., Hellstern, E. and Urban, K. (1988) J. Appl. Phys., 64, 3224. Eckert, J., Schultz, L. and Urban, K. (1989) Appl. Phys. Lett., 55, 117. Eckert, J., Schultz, L. and Urban, K. (1990a) J. Less-Common Metals, 167, 143. Eckert, J., Schultz, L. and Urban, K. (1990b) Z. Metallkde., 81,862. Eckert, J., Holzer, J. C., Krill, III, C. E. and Johnson, W. L. (1992) J. Mater Res., 7, 1751. E1-Eskandarany, M. S., Aoki, K. and Masumoto, T. (1991 ) J. Japan Soc. Powder Powder Metall., 38(1), 59. Fiorani, D. and Magini, M. (1997), Materials Science Forum, 235-238 (Proceedings of the International Symposium on Metastable, Mechanically Alloyed, and Nanocrystalline Materials, Rome, Italy, May 20-24 1996). Fischer, J. J., deBarbadillo, J. J. and Shaw, M.J. (1991) Heat Treating, 23(5), 15. Froes, E H. and deBarbadillo, J. J. (1990) Structural Applications of Mechanical Alloying (ASM International, Materials Park, OH). Gaffet, E., Malhouroux, N. and Abdellaoui, M. (1993) J. Alloys & Compounds, 194, 339. Gente, C., Oehring, M. and Bormann, R. (1993) Phys. Rev., B48, 13244. Gerasimov, K. B., Gusev, A. A., Ivanov, E. Y. and Boldyrev, V. V. (1991) J. Mater Sci., 26, 2495. Gialanella, S., Newcomb, S. B. and Cahn, R. W. (1992) in Ordering and Disordering in Alloys, ed. Yavari, A. R. (Elsevier, London), p. 67. Gilman, E S. and Benjamin, J. S. (1983) Ann. Rev. Mater Sci., 13, 279.

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Gleiter, H. (1989) Prog. Mater Sci., 33, 223. Goodwin, R S. and Ward-Close, C. M. (1995) Mater Sci. Forum, 17%181, 411. Guo, W., Iasonna, A., Magini, M., Martelli, S. and Padella, E (1994) J. Mater Sci., 29, 2436. Hack, G. A. J. (1987) Metals and Materials, 3, 457. He, L. and Ma, E. (1996) J. Mater Res., 11, 72. Heinicke, G. (1984) Tribochemistry (Akademie Verlag, Berlin). Hellstern, E., Fecht, H. J., Garland, C. and Johnson, W. L. (1989) in Multicomponent Ultrafine Micrtostructures, Vol. 132, eds. McCandlish, L. E., Polk, D. E., Siegel, R. W. and Kear, B. H. (Materials Research Society, Pittsburgh, PA), p. 137. Herr, U. (1995) Key Eng. Mater., 103, 113. Huang, J. Y., Wu, Y. K., He, A. Q. and Ye, H. Q. (1994) NanoStructured Mater., 4, 293. Huang, J. Y., Yu, Y. D., Wu, Y. K., Li, D. X. and Ye, H. Q. (1997) J. Mater. Res., 12, 936. Ivanov, E. (1992) Mater Sci. Forum, 88-90, 475. Ivanov, E., Konstanchuk, I. G., Bokhonov, B. D. and Boldyrev, V. V. (1989a) Doklady Phys. Chem., 304(1-3), 82. Jang, J. S. C. and Koch, C. C. (1990) J. Mater Res., 5, 498. Jangg, G. (1989) in New Materials by Mechanical Alloying Techniques, eds. Arzt, E. and Schultz, L. (Deutsche Gesellschaft ftir Metallkunde, Oberursel, Germany), p. 39. Jangg, G., Kuttner, R and Korb, G. (1975) Aluminium, 51, 641. Klassen, T., Herr, U. and Averback, R. S. (1997) Acta Mater., 45, 2921. Koch, C. C. (1988) Inter Mater Rev., 33, 201. Koch, C. C. (1989) Ann. Rev. Mater Sci., 19, 121. Koch, C. C. (1991) in Processing of Metals and Alloys, Vol. 15 of Materials Science and Technology--A Comprehensive Treatment, ed. Cahn, R. W. (VCH Verlagsgesellschaft mbH, Weinheim, Germany), p. 193. Koch, C. C. (1993) NanoStructured Mater., 2, 109. Koch, C. C. (1994) Inter J. Mechanochem. Mech. Alloy., 1, 56. Koch, C. C. (1996) Scripta Mater., 34, 21. Koch, C. C. (1997) NanoStructured Mater., 9, 13. Koch, C. C., Cavin, O. B., McKamey, C. G. and Scarbrough, J. O. (1983) Appl. Phys. Lett., 43, 1017. Li, W., Suryanarayana, C. and Froes, F. H. (1995a) in Synthesis and Processing of Lightweight Metallic Materials, eds. Froes, E H., Suryanarayana, C. and Ward-Close, C. M. (TMS, Warrendale, PA), p. 203. Li, W., Suryanarayana, C. and Froes, E H. (1995b) in Advances in Powder Metallurgy and Particulate Materials, Part 1, compiled by Phillips, M. A. and Porter, J. (Metal Powder Industries Federation, Princeton, NJ), p. 145. Liebermann, H. H. (1993) Rapidly Solidified Alloys: Processes, Structures, Properties, Applications (Marcel Dekker, New York). Lu, K. (1996) Mater. Sci. Eng. Rept., R16, 161. Lu, L. and Lai, M. O. (1998) Mechanical Alloying (Kluwer, Norwell, MA). Luton, M. J., Jayanth, C. S., Disko, M. M., Matras, S. and Vallone, J. (1989) in Multicomponent Ultrafine Micrtostructures, Vol. 132, eds. McCandlish, L. E., Polk, D. E., Siegel, R. W. and Kear, B. H. (Materials Research Society, Pittsburgh, PA), p. 79. Ma, E., He, J.-H. and Schilling, R J. (1997) Phys. Rev., B55, 5542. Magini, M. and Iasonna, A. (1995) Mater Trans. JIM, 36, 123. Magini, M., Iasonna, A. and Padella, E (1996) Scripta Mater., 34, 13. Martelli, S., Mazzone, G., Scaglione, S. and Vittori, M. (1988) J. Less-Common Metals, 145, 261. Massalski, T. B., Okamoto, H., Subramanian, R R. and Kacprzak, L., eds. (1990) Binary Alloy Phase Diagrams, second edition, (ASM International, Materials Park, OH). Maurice, D. R. and Courtney, T. H. (1990) Metall. Trans., 21A, 289. McCormick, R G. (1995) Mater Trans. JIM, 36, 161. McDermott, B. T. and Koch, C. C. (1986) Scripta Metall., 20, 669. Mizutani, U. and Lee, C. H. (1990) J. Mater Sci., 25, 399. Moore, J. J. and Feng, H. J. (1995) Prog. Mater Sci., 39, 243. Mukhopadhyay, D. K., Suryanarayana, C. and Froes, E H. (1994) Scripta Metall. Mater., 31, 333. Mukhopadhyay, D. K., Prisbrey, K. A., Suryanarayana, C. and Froes, E H. (1996) in Proceedings of the International Conferene on Tungsten and Refractory Metals 3, eds. Bose, A. and Dowding, R. J. (Metal Powder Industries Federation, Princeton, NJ), p. 239.

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C. S u r y a n a r a y a n a

Murty, B. S. and Ranganathan, S. (1998) Inter Mater Rev., 43 (in press). Murty, B. S., Mohan Rao, M. and Ranganathan, S. (1993) NanoStructured Mater., 3, 459. Nasu, S., Miglierini, M., Ishihara, K. N. and Shingu, P. H. (1992) J. Phys. Soc. Japan, 61, 3766. Oleszak, D., Burzynska-Szyszko, M. and Matyja, H. (1993) J. Mater Sci. Lett., 12, 3. Pathak, D., Yamada, K. and Koch, C. C. (1993) in Mechanical Alloying for Structural Applications eds. deBarbadillo, J. J. et al. (ASM International, Materials Park, OH), p. 205. Politis, C., Krauss, W., Leitz, H. and Schommers, W. (1989) Modern Phys. Lett., 3, 615. Radlinski, A. P. and Calka, A. (1991) Mater Sci. Eng., A134, 1376. Rawers, J., Govier, D. and Cook, D. (1995) Scripta Metall. Mater., 32, 1319. Russell, K. C. (1985) Prog. Mater Sci., 28, 229. Saji, S., Abe, S. and Matsumoto, K. (1992) Mater. Sci. Forum, 88-90, 367. Schaffer, G. B. and McCormick, P. G. (1989) Appl. Phys. Lett., 55, 45. Schaffer, G. B. and McCormick, P. G. (1992) Mater Forum, 16, 91. Schropf, H., Kuhrt, C., Arzt, E. and Schultz, L. (1994) Scripta Metall. Mater., 30, 1569. Schulz, R. (1996) Materials Science Forum, 225-227 (Proceedings of the International Symposium on Metastable, Mechanically Alloyed, and Nanocrystalline Materials, Quebec, Canada, 1995). Schwarz, R. B. and Johnson, W. L. (1983) Phys. Rev. Lett., 51, 415. Schwarz, R. B., Petrich, R. R. and Saw, C. K. (1985) J. Non-Cryst. Solids, 76, 281. Seki, Y. and Johnson, W. L. (1990) in Solid State Powder Processing, eds. Clauer, A. H. and de Barbadillo, J. J. (TMS, Warrendale, PA), p. 287. Shingu, P. H. (1992) Materials Science Forum, 88--90 (Proceedings of the International Symposium on Mechanical Alloying, Kyoto, Japan, May 1991). Siegel, R. W. (1991) in Processing of Metals and Alloys, Vol. 15 of Materials Science and Technology--A Comprehensive Treatment, ed. Cahn, R.W. (VCH Verlagsgesellschaft mbH, Weinheim, Germany), p. 583. Singer, R. E, Oliver, W. C. and Nix, W. D. (1980) Metall. Trans., llA, 1895. Stoloff, N. S. and Davies, R. G. (1966) Prog. Mater Sci., 13, 77. Suryanarayana, C. (1995a) Bibliography on Mechanical Alloying and Milling (Cambridge International Science Publishing, Cambridge, UK). Suryanarayana, C. (1995b)Intermetallics, 3, 153. Suryanarayana, C. (1995c) Inter Mater Rev., 40, 41. Suryanarayana, C. (1996) Metals & Materials, 2, 195. Suryanarayana, C. (1998) in Powder Metal Technologies and Applications, ASM Handbook, Vol. 7 (ASM International, Materials Park, OH) (in press). Suryanarayana, C. (1999) Prog. Mater Sci. (in press). Suryanarayana, C. and Froes, E H. (1990) J. Mater Res., 5, 1880. Suryanarayana, C. and Jones, H. (1988) Inter J. Rapid Solidification, 3, 253. Suryanarayana, C., Zhou, E., Peng, Z. and Froes, E H. (1994a) Scripta Metall. Mater., 30, 781. Suryanarayana, C., Li, W. and Froes, E H. (1994b) Scripta Metall. Mater., 31, 1465. Suryanarayana, C., Korth, G. E. and Froes, E H. (1997) Metall. Mater Trans., 28A, 293. Suryanarayana, C., Ivanov, E., Noufi, R., Contreras, M. A. and Moore, J.J. (1999) J. Mater Res., 14 (in press). Takeuchi, T., Yamada, Y., Fukunaga, T. and Mizutani, U. (1993) J. Japan Soc. Powder Powder Metall., 40, 951. Tessier, P., Zaluski, L., Zaluska, A., Str6m-Olsen, J. O. and Schulz, R. (1996) Mater Sci. Forum, 225-227, 869. Trudeau, M. L. (1994) Appl. Phys. Lett., 64, 3661. Trudeau, M. L., Schulz, R., Dussault, D. and Van Neste, A. (1990) Phys. Rev. Lett., 64, 99. Wang, G. M., Campbell, S. J., Calka, A. and Kaczmarek, W. A. (1997) J. Mater Sci., 32,1461. Weeber, A. W. and Bakker, H. (1988) Physica B, 153, 93. Yamada, K. and Koch, C. C. (1993) J. Mater Res., 8, 1317. Yang, E., Wagner, C. N. J. and Boldrick, M. S. (1993) Key Eng. Mater., 81-83, 663. Yavari, A. R. (1995) Materials Science Forum, 179-181 (Proceedings of the International Symposium on Metastable, Mechanically Alloyed, and Nanocrystalline Materials, Grenoble, France, June 27-July 1, 1994). Yavari, A. R. and Desr6, R (1990) Journal de Physique 51, Colloque C4, Supplement No. 14 (Proceedings of the International Symposium on Multilayer Amorphisation by Solid-State Reaction and Mechanical Alloying, Grenoble, France, February 21-23 1990). Yermakov, A. E., Yurchikov, E. E. and Barinov, V. A. (1981) Phys. Met. Metallogr., 52(6), 50. Yermakov, A. E., Barinov, V. A. and Yurchikov, E. E. (1982) Phys. Met. Metallogr., 54(5), 90.

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Yvon, E J. and Schwarz, R. B. (1993) in Structural Applications of Mechanical Alloying, eds. deBarbadillo, J.J. et al. (ASM International, Materials Park, OH), p. 393. Zhang, H. (1993) J. Phys.: Condens. Matter, 5, L337. Zhang, E X. and Wang, W. K. (1996) J. Alloys & Compounds, 240, 256. Zhou, E., Suryanarayana, C. and Froes, E H. (1995) Mater. Lett., 23, 27. Zhou, G. E and Bakker, H. (1996) Scripta Mater., 34, 29.

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Chapter 5 Laser Processing 5.1. Principles of Lasers 5.2. Classifications of Laser Processing 5.3. Analysis of the Laser Melting and Quenching Process 5.3.1 Heat Transfer 5.3.1.1 Absorption Phenomena of Laser 5.3.1.2 Pulsed Laser 5.3.1.3 Continuous Wave (CW) Laser 5.3.2 Kinetic Conditions for Solidification 5.3.2.1 Kinetic Conditions for Amorphous Phase Formation 5.3.2.2 Solidification Modes in the Laser Processing of Materials 5.3.3 Microstructure Selection Maps (MSM) for Alloys 5.4. Laser-Quenching 5.4.1 Amorphous Phase Formation 5.4.2 Formation of Crystalline Phases 5.5. Laser Surface-Alloying and Cladding 5.5.1 Ferrous Alloys 5.5.2 Non-Ferrous Alloys 5.6. Laser-Annealing 5.7. Laser-Beam Joining 5.7.1 Welding of Structural Materials 5.7.1.1 Ferrous-Based Alloys 5.7.1.2 Non-Ferrous Alloys 5.7.2 Microjoining 5.8. Conclusions References

89 89 92 92 92 93 94 96 96 97 97 98 98 99 100 100 101 102 106 106 106 108 112 115 115

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Chapter 5 Laser Processing KOJIRO E KOBAYASHI

5.1. P R I N C I P L E S O F L A S E R S

The laser (light amplification by stimulated emission of radiation) is a unique source of electromagnetic radiation. Lasers are distinguished from other light sources by their property to emit a coherent radiation of high spectral purity (Bimberg 1986). The total power output of a laser can be focused down to a spot of an extremely small diameter and so large intensities can be obtained. The use of focused laser energy allows one to accomplish the heat treatment of materials in a controlled spatial and temporal domain. There are at present a large number of commercially available laser sources that could be used to induce the melting needed for surface alloying. Of these a relatively small group, including the carbon dioxide (CO2), neodymium-doped yttrium-aluminum garnet (Nd:YAG), and EXCIMER (excited dimer), account for nearly all the practically used sources (Draper and Poate 1985, von Allmen 1983). The laser characteristics such as wavelength, output power, beam diameter, output mode, pulse length, and repetition rate are important processing factors. Table 5.1 lists the powers and wavelengths of different lasers. The laser processing system for a pulsed or Q-switched source offers cooling rates exceeding those of rapid solidification from the melt by several orders of magnitude. Qswitched mode operation can provide extremely higher peak power pulses and higher pulse rates than normal pulsed mode. Each individual focused laser event produces a melt spot and those melt spots must be raster-scanned on the surface if area coverage larger than the spot itself is desired. In the case of the continuous wave laser the area per unit time scanned is dependent on the effective spot size, the degree of overlap, and the relative transverse velocity.

5.2. C L A S S I F I C A T I O N S O F L A S E R P R O C E S S I N G

Figure 5.1 shows schematically the major laser processes that relate to the formation of non-equilibrium phases. The definition and characteristics of each of these processes is outlined below (Roessler 1986). Laser surface melting and quenching (Fig. 5.1(a)): A laser-beam is scanned over the surface of a substrate to be melted. Behind the laser-beam the molten pool cools very rapidly because of high heat conduction from the molten pool into the still cold substrate. 89

K. F. Kobayashi

90

Laser beam Molten / pool ~ / .............

~ ~

Laser beam /\Molten Alloying / ~pool powder / ~/

Rapidly solidified layer

(a) Laser surface melting and quenching

(b) Laser surface alloying Laser beam

Laser beam Cladding / powder [ \ [ ~,~,~~,~; 9

Alloyed layer

/ Molten ~ pool ~,/

Clad layer

Region heated above annealing / temperature ~

(c) Laser surface cladding Focused IA~mrhmam

Annealed and

/ rapidlycooled layer

(d) Laser annealing

Weld bead

(f) Laser beam soldering

(e) Laser beam welding l-- ......substrate .....

Laser beam ~ A ~ ~

p

I

~Clusters Plasma plume

(g) Laser ablation

Figure5.1. Schematicrepresentationsofvariouslaserprocessesformaterials.

Laser Processing

91

Table 5.1. Powersand wavelengths of different lasers. Source c02 laser (gas) YAG laser (solid: Nd) EXCIMER (gas: ArF) (gas: KrF)

Wavelength (microns)

Power

10.6 1.06

~50 kW CW: ~6 kW Pulse: ~20 kW ~50 W ~150 W

0.193 0.248

This produces a rapidly solidified surface layer and thereby results in the formation of amorphous or metastable phases on the surface. This process is known as laser-glazing. Laser surface-alloying (Fig. 5.1(b)): In surface alloying with a laser, an alloying material, which is delivered to the surface of a substrate by either preplacing a powder on the substrate or blowing the powder into the laser-generated melt pool, is fused to and mixed with the substrate so that the surface results in an alloyed layer after solidification. The alloyed layer shows a fine and homogeneous microstructure with minimal segregation because of the rapid solidification. The thickness of the alloyed layer can be from a few microns to about 1 mm. Laser surface-cladding (Fig. 5.1(c)): Laser surface cladding is similar to laser surface alloying except that the surface clad layer is fused to the substrate with the minimum of dilution with the substrate. Thus, the objective in this process is to provide the characteristics of a cladding material itself to the surface of a substrate. Clad layers having a thickness of up to several millimeters can be formed. Laser-annealing (Fig. 5.1(d)): A laser-beam is scanned over the surface of a substrate and heats it above an annealing temperature before rapid cooling. The objective of laser annealing is to achieve homogenization, solutionizing, recrystallization and/or removing defects locally in the surface layer. Laser annealing can allow very selective annealing of the substrate and extremely high cooling rates in comparison with conventional furnace annealing. Laser-beam welding (Fig. 5.1(e)): Fusion welding with a laser is a deep-penetration welding process. This process is characterized by the formation of a "keyhole". The laser energy causes local vaporization of material and generates a keyhole filled with vaporized metal through the thickness of the workpiece (Fig. 5.1(e)). As the focused laser-beam is scanned over the weld line, the keyhole surrounded by a molten pool moves. Continuous melting and vaporization occur at the leading edge of the keyhole. The molten metal flows back around the keyhole and resolidifies at the rear edge of the molten pool. Thus, a fusion weld joint with deep-penetration can be obtained. The major advantage of laser welding is selective heat input with narrow fusion and heat-affected zones. On the other hand, rapid heating and cooling cycles may result in the formation of unusual metastable phases in the weld fusion and heat affected zones. So, the effect of microstructural changes on the properties of the weld joint should be considered. Laser-beam soldering (Fig. 5.1(f)): Laser-beam joining can be applied to microjoining in electronics industries. Figure 5.1 (f) shows schematically laser-beam soldering between

K. E Kobayashi

92

the lead of a quad-flat pack surface-mounted component and a pad of a printed wiring board. A laser is irradiated onto the lead and the absorbed laser energy melts a solder. The solder solidifies and the solder joint is made after the laser irradiation is stopped. Laser power, beam diameter, and duration time are precisely controlled. Laser-ablation (Fig. 5.1(g)): Laser ablation is a process for the deposition of thin films. When a high energy laser-beam is irradiated onto a target surface, atoms and ions are emitted from the target to form a plasma plume. The plasma expands, cools and is condensed into clusters. The atoms, ions, clusters and/or molecules are collected on a substrate, resulting in a thin film deposition on the substrate surface. The deposited film usually has a non-equilibrium amorphous structure. The applications and characteristics of the above laser processes will be described in detail below in Sections 5.4 to 5.7.

5.3. A N A L Y S I S OF T H E L A S E R M E L T I N G AND Q U E N C H I N G P R O C E S S

5.3.1 Heat transfer The recent availability of lasers with high-power directed energy sources has helped in the development of rapid melting and solidification techniques in which a semi-infinite bulk substrate in intimate contact with a molten layer acts as the quenching medium. All mathematical treatments of classical and laser heat treatments are based on the fundamental equation for heat conduction:

OT = OT { 8t

)~ OT

)~ OT ~

O ~ C p p Ox ! Cpp 8y

)~ OT~

Cpp Oz)

qi

Cpp

(5.1)

Here T stands for the absolute temperature, t for the time, x, y and z are the coordinates, ~. the thermal conductivity, C p the specific heat, p the density and q i is the rate at which heat is supplied per unit time and volume. If all the parameters involved are known, Eq. (5.1) can be solved numerically by either finite difference or finite element methods (Mehrabian et al. 1979, Bergmann et al. 1981, Fritsch 1986).

5.3.1.1 Absorption phenomena of laser Absorption of laser power by the substrate is related to the laser-solid interaction. When an electromagnetic wave interacts with the atoms in the surface of a substrate, the field causes an electrical current. The Joule effect increases the surface temperature of the substrate due to the finite electrical conductivity of the material. Generally, the power of the laser beam is subdivided into the refected part, the absorbed part and the transmitted part. For metals that are not transparent, the last part can be neglected. Assuming that the Wiedemann-Franz law, i.e., the ratio between the electrical and thermal conductivity is constant, is valid for metals, metals with good conductivity exhibit poor absorption of laser. Similarly, the absorption ratio (fraction of the incident laser energy absorbed by the metal surface) increases with increasing temperature because the conductivity of metals decreases. When the surface is heated continuously by a laser, the surface layers start melting which makes the electrical conductivity lower and makes the absorption ratio higher.

Laser Processing

o 0.8

93

Ag

,,m

(u L.

s 0

0.6

~-

0.4

Cu

i n

k,.

0 .Q <

Fe

0.2

0.4

1

2

4

8

W a v e length (l~nn) Figure 5.2. Changein absorption ratio as a function of wavelengthfor metal surfaces.

In electromagnetic fields, the frequency affects the electrical conductivity and hence the laser light with different wavelengths shows different absorption values for a given material. Figure 5.2 shows the change in absorption ratio as a function of laser wavelength for metal surfaces. It is recognized that the CO2 laser with a 10.6 # m wavelength has an inferior absorption ratio to Nd: YAG laser with shorter wavelength (1.06/zm). Additionallyl reactions with the molten metal and the surrounding atmosphere may also affect the absorption ratio.

5.3.1.2 Pulsed laser The nature of interaction between a laser-beam and a material surface is controlled by two variables. One is the absorbed power density from the laser and the other is the interaction time. In the pulsed mode, interaction time means the time of laser irradiation. In the rapid solidification surface treatment process, glazing requires a high power density and substantially shorter interaction times. These requirements are met by the application of an Nd: YAG laser with Q switch. As discussed in detail by Bloembergen (1979) and others (Spaepen and Turnbull 1982, Linet al. (1984), picosecond pulse lasers can be used to induce very high quench rates, up to 1012 K/s, in thin molten overlays on a metallic substrate. During laser-glazing, the modified surface is considered to go through the following stages: (i) energy deposition in a layer on the order of the absorption depth (~50 nm), (ii) very fast melting ( ~ 1000 m/s) of a ~ 100 nm thick layer, and (iii) cooling and solidification.

94

K. F. Kobayashi

The quenching rate, u, is determined as a function of temperature gradient at the interface from the thermal balance.

u --

(

~.V 6 Hf

0T) Ox interface

9

(5.2)

o

Here V and A H f represent the molar volume and molar heat of fusion, respectively. On the other hand from the kinetics of interface motion u = u

Tm-T/ ~ Ti

0

(5.3)

where Tm stands for melting temperature, T/ for interface temperature and u0 is the velocity prefactor. For collision-limited motion u0 almost equals to the speed of sound, while for diffusion-limited motion uo = D~ I, where D is the diffusivity in the liquid and 1 is the interatomic distance. Generally u0 has a value between these two extremes because of a mixed mechanism. Since the temperature gradients in this process are very steep (1011 K/m during melting and 101~ K/m during cooling), u from Eq. (5.2) becomes very large and at the same time, Ti becomes considerably different from Tm from Eq. (5.3). The lifetime of the melt, determined experimentally (Lin and Spaepen 1981, Lin and Spaepen 1983) or estimated from heat flow calculations, is reported to be about 1 ns. Since the temperature of the material rises several thousand degrees in that period, the cooling rate can be estimated to be at least d T / d t = 103 K/10 -9 s = 1012 K/s (Lin and Spaepen 1984). This rapid solidification can lead to the formation of new metastable crystalline phases, or, if the crystal growth kinetics are slow enough, glass formation results. 5.3.1.3 Continuous wave (CW) laser In the pulsed mode, very high cooling rates are achieved due to the short interaction time. At the same time, however, this makes the modified region extremely small. Normally for industrial purposes, very high cooling rates are not required and the continuous mode is applied to modify the surface effectively 9 The absorbed heat-flux distribution, which is a function of the position and time, is treated moving in a CW laser while it is treated as stationary in a pulsed laser. In the basic investigation of laser-beam processing, an important aspect is the use of appropriate heat flow models. If all the parameters in Eq. (5.1) are known, it can be solved numerically by either finite difference or finite element methods 9 However, the material parameters are not always constant and their temperature dependence is not always known. Internal heat sources or sinks during transformation, melting or evaporation are not also always negligible. Cline and Anthony (1977) showed that the moving point source model, which involves an analytical solution of the thermal diffusion equation using Green's functions, could be used to estimate the temperature distribution during CW laser irradiation. During exposure and removal of a laser with a uniform energy flux onto the surface, the heating and cooling of a semi-infinite plate during the glazing process can be modeled conveniently and simply (Greenwald et al. 1979).

95

Laser Processing

During heating, the temperature at a point x below the surface at time t is given by (5.4)

T ( x , t) -- To + qJ(x, t)

and 2qi[~t 7*(x,t)----~

(

-exp

where To is the ambient temperature and tc During cooling

-

-

-~

x2 )

-

x err ( 2 ~ x t ) ]

(5.5)

)~/(CplO) is the thermal diffusivity.

T ( x , t) -- To + qJ(x, t) - ~P(x, t - r)

(5.6)

where r is the interaction time, defined in continuous mode as (beam diameter)/(beam traveling speed). The times for initial melting, r* and final solidification r** are given by Eqs. (5.7) and (5.8). r* =

Tm-T0

ab

(5.7)

(r + r*) 2 r** =

4r*

(5.8)

Here a and b are defined as follows: a -- v/rr/4K

(5.9)

b - qi/)~.

(5.1 O)

From these relations, the trailing liquid zone length, L, may be determined, since L is simply equal to v(r** - r*), in which v is the beam traveling speed. The compatibility of the model is checked by comparing this calculated length with the experimental value. Under conditions in which the energy inputs range from 1 to 100 J/cm, the heating effect is concentrated in a very thin region at the material surface. This gives rise to extremely high thermal gradients, which promote rapid solidification of the melt. At the same time, the melt and the substrate are in intimate contact, which makes the heattransfer coefficient between them infinity. These effects enable the solidification to occur faster in the case of atomization and splat-cooling types of processes. As a consequence of the extremely rapid solidification, a variety of interesting metallurgical microstructures can be developed, some of which are unique to the laser glazing process.

96

K. E Kobayashi

5.3.2 Kinetic conditions f o r solidification Crystallization kinetics, cooling rate, and laser parameters are closely connected for the intelligent processing of materials. 5.3.2.1 Kinetic conditions for amorphous phase formation The conditions for amorphous phase formation are to be derived from the nucleation and growth theory (Uhlman 1972, Davies et al. 1974, Kobayashi and Shingu 1988). Classical nucleation theory gives the following expression for the nucleation frequency J (T) (Turnbull 1950, Pound and Lamer 1974, Miyazawa and Pound 1974, Spaepen 1975, Spaepen and Turnbull 1976):

100 [

J ( T ) -- ~ e x p rl

-

3

1 - Tr

(5.11)

where r/= viscosity of liquid, Tr = reduced temperature (=T/Tm) and Tm = equilibrium melting point; c~ is a factor which depends on the atomic arrangement at the interface with a value close to unity, and/3 is the entropy change due to solidification in units of the gas constant. The growth rate of crystal nuclei is given by molecular kinetic considerations assuming that the atomic transport across the interface region, which usually is assumed as several atoms diameters in thickness, takes place by a diffusive process according to the equation: Ri(T) :

K r/(1 -- Tr)

(5.12)

where K = d 3L f/32rq6 and d = atomic diameter, L f = heat of fusion, and q~ = interface thickness. The temperature dependence of viscosity in the undercooled liquid region is represented assuming a Tamrnann and Hesse type of relation, as log 7? = A +

B

T - T~

(5.13)

where A, B and T~ are constants which are determined assuming r/ -- 1013 poise (10 le Pa sec) at T -- Tg (glass temperature); r/ is the value for the Arrhenius relation at T = Tm and d(log r/)/d(1/T) is the slope of the Arrhenius relation at T = Tm. The general form of the fraction of phase transformed X may be given as - In(1 -- X) -- -4n. ~ fot J (T) [fo t Ri (co) do)]3 dO

(5.14)

where J ( T ) is the nucleation frequency, Ri(w) is the linear growth velocity, X the crystallized fraction and t the time after passing the melting point during cooling. The temperature-time-transformation (TTT) diagram and the continuous-coolingtransformation (CCT) diagram can be obtained from Eq. (5.14) (Cahn 1956, Grange and Kiefer 1941). By the use of Eq. (5.14) we can estimate the condition for amorphous

Laser Processing

97

phase formation by requiring that the crystallized volume fraction X in the case of laser quenching be very low. 5.3.2.2 Solidification modes in the laser processing o f materials

The surface treatment of materials by laser is of great interest in many technical applications. By moving at high speed, a high power beam, such as that produced by a CO2 laser, over a surface heat treatment or rapid melting-solidification of the base material can be obtained at depths varying from a few microns to several millimeters. For the case of CW laser sources and metals, heat flow has been modeled by Mehrabian and co-workers (Mehrabian 1982, Hsu et al. 1978, Hsu et al. 1980, Kou et al. 1981, Sekhar and Mehrabian 1981) and others (Mazumder and Steen 1980, Chande and Mazumder 1982, Chande and Mazumder 1983, Lin and Spaepen 1984). During laser remelting, the rate of solidification governs the possible formation of metastable phases, their composition, and the scale of the microstructure (Boettinger et al. 1980, Kurz et al. 1986, Trivedi et al. 1987). Laser resolidification starts epitaxially on a crystalline substrate and nucleation plays a less important role. At the high rates of absolute stability, one cannot assume equilibrium partitioning any more. Such departure from equilibrium partitioning was discussed by Boettinger et al. (1984) who used two response functions to describe the temperature and the solid composition at the interface, where the variation of the partition coefficient with growth velocity was given by Aziz (1982) and Jackson et al. (1980). An amorphous phase should form in two different cases. In the first case, the cooling rate is too fast for partitioning of atoms to occur. In the second, partitioning and/or rearrangement of atoms for crystallization is too slow for a given cooling rate (Kobayashi and Shingu 1982, Kobayashi et al. 1984). 5.3.3 Microstructure selection maps ( M S M ) f o r alloys

Recently a series of solidification models have been developed which are valid at low and high growth velocities for different microstructures (Trivedi et al. 1987, Kurz et al. 1986, Boettinger et al. 1988, Lipton et al. 1987, Carrad et al. 1992). These models predict phase selection and the solidification microstructure of an alloy as a function of solidification parameters such as the growth rate and temperature gradient. These are discussed for dendrites (Kurz et al. 1986, Gill and Kurz 1995, Gilgien et al. 1995, Kurz 1995) or eutectic (including peritectic) structures (Rappaz et al. 1987, Kurz and Trivedi 1994, Gill and Kurz 1995, Gilgien et al. 1995) in rapid solidification processing such as in welding and laser treatment. MSMs contain information on the phase selection, solidification morphology and sometimes size of the microstructure as a function of composition and growth rate (or growth temperature, often combined with the equilibrium phase diagram). The modification of laser treatment conditions by controlling the atmosphere and the initial temperature of the sample allowed the reduction of instabilities in the liquid bath at low beam velocities. This made it possible to observe the metastable eutectic A1-A16Fe and the transition from the metastable eutectic to Al-dendrites for A1-4at%Fe alloy during laser resolidification (Gilgien et al. 1995). In the case of stainless steel, metastable phases such as y may form in a composition range where under normal welding conditions

98

K. F. Kobayashi

still forms. This becomes clear when one considers the fact that a phase with less solute rejection is preferred at high growth velocities even if its liquidus temperature is lower than that of the equilibrium phase (David and Vitek 1989). In high power laser treatment (>5 kW) or in welding, the melt pool might become so large that buoyancy driven convection becomes important (Kurz and Trivedi 1994). Preheating lowers the surface temperature gradient that decreases the driving force for Marangoni convection. Preheating is therefore useful for low velocity laser resolidification.

5.4. LASER-QUENCHING

5.4.1 Amorphousphaseformation Laser quenching yields an amorphous phase over most of the phase diagram in Pd-Si and Pt-Si systems (von Allmen et al. 1980). For the formation of an amorphous phase by laser-alloying methods, it must not only be produced as a result of an individual laser event, but the amorphous region thus produced must remain amorphous throughout subsequent heat cycling needed for overlap in multi track treatment. Fe-B alloys between 5 and 24 at%B were found to be amorphous following laser-alloying (Lin and Spaepen 1982). In this work, Nd : YAG laser melted the alloy for 30 picoseconds (ps) and it was quenched at the high cooling rate of about 101~ K/s. The formation of amorphous phases from pure metallic melts by short laser pulses has also been reported. Frequency-quadrupled Nd-laser pulses (wavelength: 0.265 #m) of 10 ns duration incident on single-crystal Si formed amorphous zones a few tens of nm thick (Tsu et al. 1979). Liu et al. (1979) formed amorphous Si by using 30 ps pulses of both 0.53 and 0.265 #m radiation and the thickness of the zone was estimated to be about 30 nm. Amorphization kinetics of Si prepared by short laser pulses were discussed by Cullis et al. (1982) and by Rozgonyi et al. (1982). In the case of pure A1, the reported amorphous region contained small grains of crystalline A1N (Mazzoldi et al. 1980). Recently, amorphization of Co-Ti thin films was achieved after quenching following melting by a 5 ns Nd: YAG laser pulse (Vitta et al. 1994). The Co-Ti system is suitable for the study of amorphous phase formation at compound-forming compositions since the compounds exhibit a variety of homogeneity ranges and structures. The competition between the formation of an amorphous phase and the terminal solid solution was investigated in this work. The solidification was partitionless and complete amorphous phase formation was found for Ti contents greater than about 23 at% in laser quenched sampies. This compositional limit of the amorphous phase formation agrees well with the prediction based on the influence of atomic size difference on the topographical stability of terminal solid solution that was proposed by Egami and Waseda (1984). Amorphous alloy films were deposited by EXCIMER laser ablation using sintered Ta-Ni targets (Yano et al. 1996). By irradiating a solid target with a pulsed laser-beam of high energy density, atoms and ions forming plumes are emitted from the target, and deposited on a substrate as an amorphous phase. These Ta-Ni amorphous alloy films are known to have extremely high corrosion resistance in both acid and alkaline solutions (Mitsuhashi et al. 1987, Lee et al. 1995). The pulsed laser deposition was performed with

Laser Processing

99

a KrF EXCIMER laser (Lamda Physics LPX305i) having a wavelength and beam dimensions 248 nm and 15 mm x 35 mm, respectively. In another case, the effect of nitrogen on the pulsed laser deposition of amorphous silicon carbide films was investigated to evaluate the properties and structure (Yee et al. 1996). The mechanical, optical, and electrical properties of amorphous carbon-based films (Kurokawa et al. 1987, Fraunheim et al. 1989, Elhossary et al. 1990, Liu and Cohen 1990, Lu et al. 1992, Schwan et al. 1994) can be controlled by focusing laser pulse from ArF EXCIMER laser (Questek 2740SC) operating at 193 nm onto a target. CO2 laser surface treatment consisting of laser surface melting and subsequent selfquenching is also useful to form amorphous phases (Bergmann and Mordike 1981, Kumagai et al. 1986, Hirose et al. 1992, Hirose et al. 1994). An amorphous surface has a variety of useful characteristics such as extremely high corrosion resistance and unique electrocatalytic activity. Hirose et al. (1994) fabricated an amorphous magnetic alloy layer by CO2 laser quenching. The laser-quenched Fe78B 13Si9 alloy layer consisted of an amorphous phase and was utilized to construct a non-contact-type static torque sensor with good sensitivity and linearity.

5. 4.2 Formation o f crystalline p h a s e s

Laser melting and quenching is used not only for the formation of amorphous phases but also for understanding the solidification process. During CO2 laser surface resolidification of Ni3 (A1, Ti) alloys, a sharp transition was observed with increasing growth velocity from the growth of an ordered intermetallic to metastable structures (Hunziker and Kurz 1997). This transition is attributed to disorder trapped in the ordered phase and the corresponding undercooling. High-intensity laser irradiation was used for melting silicon carbide in a high-pressure inert gas atmosphere (Sadler and Shamsuzzoha 1997). By this technique, an entirely new class of materials such as cast silicon carbide shapes and silicon carbide fibers spun or blown directly from the melt might be possible. Badan et al. (1985) formed cellular or dendritic layers, with carbon and chromium supersaturation and metastable austenite as the main constituent on the steel surface by CW-CO2 laser. In the case of YAG laser processing, non-crystalline Nb40Ni60 was formed (Lin and Spaepen 1986) and the morphology of nanostructures on the scale of 1-3 nm could be altered by the laser quenching process (Scandella et al. 1991). Using KrF (248 nm) EXCIMER laser of relatively long (45 ns full width half maximum (FWHM)) pulse duration, melting and subsequent solidification of thick (190410 nm) amorphous-silicon (a-Si) layers produced by ion implantation was investigated (Lowndes et al. 1987). "Fine grained polycrystalline silicon" region produced by explosive crystallization of a-Si contained large numbers of disk-shaped silicon flakes, and the onset of epitaxial regrowth was marked by a slight decrease in surface melt duration. KrF EXCIMER pulsed laser deposition (PLD) has been employed successfully to grow high-quality films that are fully dense, long lasting, and have low friction surfaces. Film properties such as crystallinity, stoichiometry, morphology and dopant concentration can easily be adjusted by varying PLD parameters, since it is exactly these properties that are important for determining the tribological behavior. The feasibility of using PLD to grow

100

K. F. Kobayashi

WS2 films has been evaluated (Prasad et al. 1992). PLD was used to grow adherent films of WS2 on 440C stainless steel substrates (Zabinski et al. 1994). EXCIMER laser is quite a useful tool for the synthesis of electronic materials. A pulsed laser deposition technique formed strained dielectric superlattices of BaTiO3 (BTO) and SrTiO3 (STO) (Tabata et al. 1994). A large stress of 400-500 MPa is introduced at the interface between the BTO and STO layers. A large dielectric constant of 900 was observed with a stacking periodicity of 2 unit cells/2 unit cells. The superlattices show drastically different electrical behavior from that of the solid solution (Sr, Ba)TiO3 films. Dielectric superlattices of (Sr, Ca)TiO3/(Ba, Sr)TiO3 also formed by a pulsed EXCIMER laser ablation (Tabata and Kawai 1997a). The crystal structure is controlled with atomic order accuracy and a large lattice stress of 0.5-1 GPa can be introduced periodically at the interfaces owing to the lattice mismatch of the constituent layers. For the multi-layered films, a laser ablation technique produced a new concept device with a combination of ferroelectric and ferromagnetic materials (Tabata et al. 1997b). An ideal hetero-epitaxy can be obtained owing to the similar crystal structure of perovskite type ferroelectric Pb(Zr, Ti)O3; (so called PZT) and ferromagnetic (La, Sr)MnO3, which is useful for application in DRAM (dynamic random access memory) and SRAM (static random access memory) devices. For the superconducting materials, a laser ablation technique using ArF EXCIMER laser formed Bi-based layered perovskite oxide films (SrBi2Ta207) on the Nb-SrTiO3 substrate or Bi2Sr2CuO6 superconducting bottom electrode with c-axis perpendicular to the substrate surface (Tabata et al. 1995). A multi-target, laser-ablation technique inserted layers of undoped, semiconducting La2CuO4 or SmzCuO4 in a controlled manner in the Lal.ssSr0.15CuO(LSCO) compound superlattices (Tabata et al. 1993). 5.5. L A S E R S U R F A C E - A L L O Y I N G

AND CLADDING

Laser surface treatment is a material-processing method that utilizes the high power density (CO2 laser) available from focused laser sources to melt metal coatings and a portion of the underlying substrate. Laser cladding uses a laser-beam to melt and alloy powders of different composition placed on top of the substrate as well as a thin layer of the substrate. The main purpose of laser cladding is to form a thin interfacial layer of an alloy on a given substrate with minimum dilution of the clad layer and to impart good surface wear and corrosion resistance to the material. The advantages of this technique are production of noble alloys, minimized clad dilution, reduced alloy material loss, reduced machining costs, and reduced distortion. 5.5.1 Ferrous-alloys

A typical use of laser surface treatment is to modify the surface structure of gray cast iron, with the aim of producing a white ledeburitic microstructure while retaining the gray core (Bergmann 1985). The advantage of such a treatment is the combination of good core ductility, a result of stable Fe-C solidification producing an iron-graphite microstructure, and surface wear resistance resulting from the Fe-Fe3C layer obtained in a metastable state.

Laser Processing

101

The phase analysis of the metastable structure was investigated in a laser melted Fe-C alloy (Zarubova et al. 1996). As the laser transformation hardening is characterized by extremely fast heating and self-quenching rates, the presence of alloying elements such as Cr and Mn in steels, that promote hardenability and tend to partition to the carbides in the initial microstructures, is likely to prevent complete austenitization and development of a nearly fully martensitic hardened layer (Bradley and Kim 1988). To improve wear resistance of the AISI 304L austenitic stainless steel, without degrading the inherent corrosion resistance, laser surface treatment was utilized. For this purpose, molybdenum and tantalum were used as alloying elements and they formed a substitutional solid solution in the austenite phase (Akgun and Inal 1995). A 10 kW CO2 laser clad Fe-Cr-Mn-C alloy on an AISI 1016 steel produced a very fine grained microstructure of ferrite and complex carbide precipitation (Singh and Mazumder 1987a). The clad exhibited far superior wear properties compared to Stellite 6 during block-on-cylinder tests. The improved wear resistance was attributed to the fine distribution of metastable M6C carbides in the ferrite matrix. Laser clad Ni-Fe-Cr-A1Hf alloys showed a high degree of grain refinement, considerable increase in solubility of Hf in matrix and the occurrence of Hf-rich precipitates and new metastable phases which contributed towards an increase in the oxidation resistance of the alloy (Singh and Mazumder 1987b). A 10 kW CO2 laser-clad Ni-Cr-A1-Y alloy with either niobium carbide (NbC) or chromium carbides (Cr3C2, Cr7C3, Cr23C6) had higher wear resistance than that of the Ni-Cr-A1-Y alloy cladding without the carbides (Hirose and Kobayashi 1995). Laser cladding of Ni-Cr and Ni-P alloys on steel enhanced the corrosion resistance (Renaud et al. 1991, Garcia-Alonso et al. 1996).

5.5.2 Non-ferrous alloys

Aluminum, titanium, and copper alloys are the most popular materials, but their low wear resistance prevents their widespread applications. A titanium aluminide layer was formed on an aluminum surface by laser processing. By applying a laser-beam to a powder mixture of aluminum and titanium previously placed on the aluminum base metal, the mixture was superficially melted and titanium aluminide alloy layers with a thickness of 400 # m were formed at a relatively high speed. By forming this titanium aluminide alloy layer on an aluminum surface, the wear properties of aluminum have been considerably improved without losing the advantage of lightness of this material (Uenishi et al. 1992a, Uenishi et al. 1992b). A13Ti-based intermetallic compound matrix composite (IMC) layers were formed by using a mixture of aluminum, titanium and ceramic powders as preplaced powders (Uenishi and Kobayashi 1993, Uenishi and Kobayashi 1994, Uenishi and Kobayashi 1996a, Uenishi and Kobayashi 1996b). These IMC layers exhibited good wear resistance. The same technique can be applied to modify the surface of titanium (Hirose et al. 1993). The intermetallic compound clad layers were successfully formed on titanium substrate. A typical example of the clad layers is shown in Fig. 5.3. The clad layers showed good wear and oxidation resistance.

K. E Kobayashi

102

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Figure 5.3. Ti3AI laser alloyed layer formed on pure titanium surface using 2.5 kW CO2 laser.

Surface modification of copper using a laser is suggested to improve the strength and wear resistance of pure copper at elevated temperatures without sacrificing its electrical properties. In particular, Cu-Cr alloyed layers were produced on a pure copper substrate by irradiating a laser-beam onto chromium powder beds or plasma spray-deposited on the substrate surface (Hirose and Kobayashi 1994).

5.6. L A S E R - A N N E A L I N G

Laser-annealing is characterized by the following advantages in comparison with conventional processes such as furnace annealing" 9 Chemical cleanliness,

Laser Processing

103

9 Extremely short duration of heating cycle and, therefore, high heating and cooling rates, 9 Less thermal damage and distortion, and 9 Controlled shape and location of the treated region. Because of its selective heating of the workpiece, laser-annealing is promising for heat treatment of semiconductor materials as an alternative to furnace-annealing in the microelectronics industry. In conventional IC (integrated circuit) fabrication, the lattice damage caused by ion implantation is removed and the implanted dopants are activated by furnace annealing for at least 30 min at temperatures of 820 K to 1470 K. Laser annealing can more rapidly and effectively remove the lattice damage and activate the dopants. Although significant damage usually remains in the form of dislocation loops after furnace annealing, pulsed-laser annealed specimens have no damage down to a resolution of approximately 1 nm (Young and Wood 1982, Young et al. 1982). Figure 5.4 shows the carrier concentration as a function of implanted dose for pulsed laser and thermally annealed boron-implanted silicon (Young and Wood 1982). In laser-annealed silicon the carrier concentration increases linearly beyond the dose of 1.5 • 1016/cm2 that corresponds to the equilibrium solubility limit of boron in silicon in this case, whereas it saturates for thermal annealing. It is clearly shown in this figure that the non-equilibrium process of laser annealing can allow active carrier concentration in excess of the equilibrium solubility limit. This is achieved due to the extremely high heating and cooling rates in the laser processing. In laser-annealing, the very short duration time and the ease of controlling the shape and the location of the treating region are also beneficial to form shallow junctions, which are needed for highly integrated devices such as megabit SRAMs (Tsukamoto et al. 1992). In conventional furnace annealing or rapid thermal annealing systems, shallow p + n junction formation using boron ion implantation is difficult, because the implanted boron diffuses very rapidly into silicon on subsequent annealing (Tsukamoto et al. 1992). An ultrashallow p + n junction, less than 50 nm deep, has been formed by low power BF2 ion implantation followed by EXCIMER laser-annealing with single pulse irradiation (Tsukamoto et al. 1992, Tsukamoto et al. 1996). Laser-annealing has also been applied to recrystallize amorphous layers of materials to improve their electrical properties. Polycrystalline silicon thin-film transistors (polySi TFTs) are used for large-area electronics applications such as liquid crystal displays and image sensors, because poly-Si TFTs have higher carrier mobility than amorphous silicon (a-Si) TFTs. EXCIMER laser annealing has been applied to the fabrication of highperformance poly-Si TFTs on non-heat-tolerant glass substrates, because its extremely short duration and shallow melt-regrowth avoids thermal damage to the glass substrate (Sameshima et al. 1986, Sera et al. 1989, Kuriyama et al. 1991). Recently large-grain poly-Si TFTs have been obtained by controlling the solidification velocity during dualEXCIMER laser irradiation both from the front side and from the back side (Shimizu et al. 1993) and by low temperature (<673 K) substrate heating during EXCIMER laser annealing (Kuriyama et al. 1992, Kuriyama et al. 1993). Besides the microelectronics applications, annealing of amorphous materials by laserbeams has been applied to improve the magnetic properties of materials (Matteazzi et al.

K. F. Kobayashi

104 A 04 I

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1985, Sato et al. 1985, Lanotte and Lannotti 1995). Due to the very short duration time and rapid temperature change, laser-heating can produce uniform crystallization at a temperature that is much higher than that reached by other methods (Lanotte and Lannotti 1995). This is very useful for the synthesis of nanocrystalline materials through crystallization of amorphous materials. It has been known that nanocrystalline ferromagnetic materials have improved soft magnetic properties in comparison with crystalline or amorphous materials (Herzer 1992). Pulsed YAG laser irradiation of thick Fe80.5(BSiC)19.5 amorphous ribbons (>50 mm), which is a transformer core material, has decreased the core losses because of refinement of the domain widths (Matteazzi et al. 1985). CW CO2 laser annealing of Fe73.sCulNb3Si13.sB9 amorphous ribbons produced nanocrystalline structure of 50 nm grain size (Lanotte and Lannotti 1995). The laser annealed materials had good soft magnetic properties (Lanotte and Lannotti 1995). It has also been reported that continuous laser-annealing improved the soft magnetic properties of Fe40Ni40P14B6 amorphous ribbons (Sato et al. 1985). Selective and local annealing by laser-beam can be applied to structural materials. Laser irradiation for solution-annealing (laser solution-annealing) has been performed on the surface of a Ni-base superalloy, Inconel 718, to reduce its sensitivity to hydrogen embrittlement (Hirose et al. 1996). Inconel 718, which is one of the most used Ni-base superalloys, is age-hardened by the homogeneous precipitation of yt and y~f phases, which allows the alloy to have a high strength at moderate temperatures, and good creep and fa-

105

Laser Processing A

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tigue resistance. However, the high sensitivity of Inconel 718 to hydrogen embrittlement is a serious problem in using the alloy, for example, in a liquid hydrogen fueled rocket engine. Hydrogen-induced cracking tends to occur at the surface of stress concentrated regions. Reducing the hardness of such surface regions locally is known to be effective in preventing such cracking. To this end, laser surface solutionizing has been applied to Inconel 718 using a 2.5 kW CO2 laser. The surface-solutionized zones can be obtained without melting the treated zones by optimizing the laser conditions. While the base metal has a hardness of approximately 450 Hv, the hardness of the laser solutionized zones is reduced to below 250 Hv as shown in Fig. 5.5. No precipitates of yl and y1~ phases are identified in the laser solutionized zones by TEM observations. The surface regions are heated beyond the solvus temperatures of y~ and y~ phases and these phases are dissolved in the matrix. The subsequent rapid cooling can prevent the precipitation of the age hardenable phases. Surface-solutionizing can be achieved by laser-scanning. The desired depth of the solutionized zones can be obtained by controlling the laser parameters, traverse speed and defocus distance, as shown in Fig. 5.6. The ductility of the surface solutionized specimen is almost twice that of the base metal in a tensile test under a 29.4 MPa hydrogen atmosphere at room temperature with hydrogen pre-charging. Thus, the laser surface softening is effective in reducing the sensitivity of Inconel 718 to hydrogen embrittlement with little sacrifice to the strength.

106

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10 15 20 25 30 Traverse speed (mm/s)

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Figure 5.6. Effectof traverse speed on depth of the laser surface solutionized zone for different defocus distances (Hirose et al. 1996).

5.7. LASER-BEAM JOINING Joining is one of the most important applications of lasers including both solid-state and gas lasers, ranging from joining and bonding devices in microelectronics to welding structural materials. Laser-beam joining is characterized by its high energy density and flexibility, and has obvious advantages, such as high speed, narrow fusion and heat affected zones (minimizing distortion and damage of adjacent areas), ability of remote noncontact processing, and ease of automation. 5. 7.1 Welding o f structural materials 5. 7.1.1 Ferrous-based alloys Laser-beam welding is an attractive process for joining steels due to its high productivity and possibility of automation. In particular, automotive and shipbuilding industries have an interest in the applications of laser-beam welding. Laser-beam weld fusion and heat affected zones are subjected to significantly more rapid heating and cooling cycles compared to conventional arc welds. In the case of laser-beam welding of C-Mn steels, HY (high yield strength) steels and HSLA (high strength low alloy) steels, metastable phases such as martensite and bainite having high hardness are produced in the weld fusion and heat affected zones (Metzbower and Denney 1986, Moon et al. 1987, Metzbower 1990). Thus, these laser-beam welded zones have higher hardness in comparison with arc welded zones. Therefore, predictions of hardness of laser-beam welds are desired. Ion et al. have developed simplified analytical models of heat flow and phase transformations, and applied them to laser-beam welding of C-Mn steels to predict the maximum heat affected zone hardness (Ion et al. 1996). In this model the heat cycles in the heat affected

107

Laser Processing

Cooling time from 1073K to 773K o~" 0.30 W W r

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Figure 5.7. Diagram of carbon equivalent against absorbed energy showing theoretical contours for maximum heat affected zone hardness in carbon manganese steels (Ion et al. 1996).

zones have been simulated with a formula based on the thin plate analysis of Rosenthal (1946). Microstructural constituents in the heat affected zones have been estimated from the simple carbon equivalent, Ceq = C + Mn/12 + Si/24 (where the concentrations of the different elements are given in wt.%), and cooling time between 1073 K and 773 K using empirical formulas obtained by analyzing experimental data. The hardness of individual phases, namely martensite, bainite and ferrite/pearlite, has been empirically obtained from the carbon equivalent. Finally the maximum heat affected zone hardness has been determined by summing the contributions from the individual phases using a simple rule of mixtures. Figure 5.7 shows the diagram for predicting the maximum heat affected zone hardness from the carbon equivalent and absorbed energy using this model. In laser-beam welding of stainless steels the microstructural changes in weld fusion and heat affected zones due to rapid cooling, which affect both the mechanical and corrosion properties, should be considered. In conventional arc welding, the room-temperature microstructures of weld-bead are estimated using Schaeffler diagram (Schaeffler 1949). However, the solidification microstructures of laser welds subjected to more rapid cooling and higher solidification rate than arc welds are significantly different from those estimated by Schaeffler diagram (Vitek et al. 1982, David and Vitek 1985, David et al. 1987, Khan et al. 1988, Zambon and Bonollo 1994). It has been shown that a significant narrowing of the duplex ferrite-austenite region in the diagram occurs with cooling rates typical of laser-beam welding as schematically shown in Fig. 5.8 (David et al. 1987). Thus, while

K. F. Kobayashi

108 A

28

"

._+ o" 4)

\

4~,~ 0

,usten,te(,)

/

~M+F~....,..~"'~

4 8 12 16 20 24 28 Creq=%Cr+%Mo+l.5x%Si+O.5x%Nb

Ferrite (F)

"

32 36 40 (mass%)

Figure 5.8. Schaeffler diagram showing schematically the convergence of the 0 and 100% ferrite lines in laser-beam welding (Davidet al. 1987).

austenitic stainless steels such as type 304, 308 and 316 have two-phase ferrite (3) plus austenite (y) microstructures that are found after primary ferrite solidification under normal welding conditions, these steels often have a fully austenitic microstructure under rapid solidification conditions such as laser-beam welding. This is caused by a change in the solidification mode from the equilibrium primary ferrite solidification mode to a non-equilibrium primary austenite solidification mode due to rapid solidification (Vitek and David 1994). 5. 7.1.2 Non-ferrous alloys The low overall heat input and, thereby, the small fusion and heat affected zones makes laser-beam joining suitable for welding age-hardenable alloys, which have a problem of softening in welded zones, in comparison with conventional arc welding processes. The characteristics of laser-beam welds of age-hardened A1 and Ni base alloys will be hereinafter presented in comparison with their tungsten inert gas (TIG) welds. Age-hardened A1-Mg-Si alloy, 606 l-T6, thin plates (1 mm in thickness) have been welded with full penetration using a 2.5 kW CO2 laser (Hirose et al. 1997b). In laser-beam welding a full penetration weld can be achieved even at a welding speed of 167 mm/s, whereas in TIG welding a welding speed of 5 mm/s or less is required for the full penetration weld. The cross section of a typical laser-beam weld is shown in Fig. 5.9. The bead width of the laser-beam weld is approximately 1/3 that of the TIG weld. Figure 5.10 shows the hardness profiles obtained from the laser-beam and TIG welds in the as-welded condition and after an artificial aging treatment at 448 K for 28.8 ks. In the as-welded condition, although a softened region is found in both the laser-beam weld and the TIG weld, the width of the softened region in the laser-beam weld is less than 1/4 that of the TIG weld. The softened region is caused by reversion (dissolution) of strengthening/~

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(Mg2Si) precipitates due to the heat input of welding. The hardness of the softened region in the laser-beam weld is almost fully recovered to the level of the base alloy after an aging treatment at 448 K for 28.8 ks. In contrast to this, the TIG weld subjected to the aging treatment after welding has an appreciable hardness drop in a region between 4 and 7 mm from the bead center. This region is thought to be the region where the solute (Mg and Si) content in the matrix is insufficient for precipitation of the strengthening fl (MgzSi) phase because of precipitation of nonhardening fif and/or fl1~ phases during weld thermal cycle. Therefore, the TIG weld will require a solution annealing treatment before artificial aging to fully recover the heat-affected zone hardness unlike the laser-beam welds. In laser-beam welding, since the cooling rate in the heat affected zone is rapid enough to prevent the precipitation of fl~ and/or fl~ phases, supersaturated solid solutions can be obtained over the whole region of the heat affected zone, and thereby sufficient age-hardening due to precipitation of fi1~ phase can be achieved after the artificial aging. As expected from the hardness profile in Fig. 5.10, the tensile strength of the specimen subjected to the aging treatment after laser-beam welding is almost equivalent to that of the base plate. Laser-beam welding has been applied to joining an age-hardened Ni base superalloy, Inconel 718, plate (3.2 mm in thickness) using a 2.5 kW CO2 laser (Hirose et al. 1997a, Hirose et al. 1998) and its properties have been compared with those of TIG welds. Inconel 718 is strengthened by intermetallic ~,~ and V ~ phases, which are precipitated after solution annealing treatment followed by a two step aging treatment. Microstructures of the weld fusion zones are shown in Fig. 5.11. The TIG weld fusion zone has a cellular dendritic structure having austenite (Vf)/Nb-rich Laves phase eutectic and Nb-rich MCtype carbides in the interdendritic regions (Fig. 5.11 (a)). The laser-beam weld fusion zone

110

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has significantly finer solidification structure than that of the TIG weld (Fig. 5.1 l(b)). In the laser-beam welding, the cooling rate during solidification that is estimated from the relationship between dendrite arm spacing and cooling rate is 100 times greater than that in the TIG welding. Thus, the high energy concentration of the laser-beam welding results in rapid solidification and thereby fine solidification structure. Figure 5.12 shows the hardness distributions obtained from the laser-beam and TIG welds in the as-welded condition and after the two-step aging treatment. The fusion zone width of the laser-beam weld is approximately 1/6 of that of the TIG weld. Apparent softening occurs in the weld fusion and heat-affected zones because of a reversion (dissolution) of strengthening y' and y" phases. However, the laser-beam weld has a significantly narrower softened region that is approximately 1/7 of that of the TIG weld. In the laser-beam welds, the hardness of the fusion and heat-affected zones is fully recovered to that of the aged base alloy after the aging treatment. In the TIG welds, the hardness of the weld fusion zones after the aging treatment is not completely recovered to the level of the base metal. Since in the TIG weld fusion zones coarse Nb-rich phases are formed in interdendritic regions, the solute concentrations of the y' matrix may be insufficient for full precipitation of the y' and y" phases. These hardness distributions apparently reveal an advantage of laser-beam welding that its low overall heat input results in fine solidification structure and minimize the weld fusion and heat affected zones. The tensile strength of the laser-beam weld joint subjected to the aging treatment after welding fully recovers to that of the aged base alloy, whereas that of the TIG weld joint does not because of the coarse interdendritic phases and lower hardness in the weld fusion zone.

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111

Figure 5.11. Microstructures of weld fusion zones of Ni-base superalloy, Incone1718. (a) TIG welding and (B) CO2 laser-beam welding.

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112

K. E Kobayashi

Narrow weld-fusion and heat-affected zones, and minimizing distortion in laser-beam welding, are also beneficial to fusion welding, of brittle materials such as intermetallic compounds. Laser-beam welding of a molybdenum-containing TiA1 intermetallic compound (Ti-46at%A1-2at%Mo) has been performed using a 2.5 kW CO2 laser (Hirose et al. 1995a, Hirose et al. 1995b). The cooling rates in the welded zones have been controlled by changing the welding speed and pre-heating temperature. The base metal has the lamellar structure containing y(TiA1) + otz(Ti3A1). However, the microstructure and hardness of the laser fusion zones change depending upon the cooling rates as shown in the schematic CCT diagram (Fig. 5.13). When the cooling rate at 1390 K is above 3000 K/s, the ot ~ y transformation is suppressed and the fusion zone structure results in a single phase of ot2. If the cooling rate is less than 3000 K/s, massive y phase is seen together with the Ot2 phase. In laser-beam welding, where the cooling rates are approximately 80 K/s to 1800 K/s, the fusion-zone microstructure mainly consists of massive ot2, massive y and lamellar (c~2 + y) phases. Thus metastable phases of massive Ot2 and massive y form owing to rapid cooling during laser-beam welding. A fully lamellar structure is formed if the cooling rate is less than 20 K/s. The hardness of the fusion zones increases from 315 Hv of the base metal hardness to above 500 Hv with increasing cooling rate. Cracking in the weld fusion zones is sensitive to the cooling rates and hence to the microstructures of the fusion zones. When the hardness value of the fusion zones is reduced below 410 Hv, cracking dramatically decreases. As a result, crack-free welds can be obtained at welding speeds below 50 mm/s and pre-heating temperatures above 673 K (Fig. 5.14). The cross-section of a crack-free weld bead is shown in Fig. 5.15. During tensile testing, the laser welded specimens without cracking fractured in the base metal. Thus, laser-beam welding with preheating can provide narrow weld fusion and heat-affected zones, and optimal microstructures of the weld fusion zones that are beneficial to welding of brittle materials.

5.7.2 Microjoining

Microjoining in electronics industries also is an important application of laser-beam processing. Laser-beam soldering has been applied to surface-mount technology, namely interconnection between electronic components and printed circuits since 1980's (Chang 1986, Semerad et al. 1993, Messier and Millard 1994). In laser-beam soldering, ease of controlling the shape and location of the heating region brings about reliable solder joints with minimal component heating and is suitable for fine-pitch soldering and high-density packaging. While CO2 laser and Nd:YAG laser are used for laser soldering, Nd: YAG laser is more suitable because of preferential absorption by solders as opposed to board materials (Messier and Millard 1994). Since only the solder joint area is heated during laser-beam soldering process without damaging the boards and components, this process is more beneficial to interconnections using solders with a wide range of melting temperatures in comparison with conventional infrared reflow process (Messier and Millard 1994). Thus, this can overcome the problem of soldering with Pb-free solders with higher melting temperatures than those of eutectic or near eutectic Sn-Pb solders (Yang et al. 1995). Rapid heating and cooling thermal

Laser Processing

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2at%Mo using 2.5 kW CO2 laser.

cycles in laser-beam soldering result in the microstructures and mechanical properties of the solder joints different from those of the conventional infrared reflowed joints. Laser-beam soldering of eutectic Sn-Ag solder has been performed and the effect of the soldering process on the microstructures and mechanical properties have been evaluated (Yang et al. 1995). In this study eutectic Sn-Ag solder joints between quad-flat pack surface-mounted components and printed wiring boards have been produced using a Nd : YAG laser soldering system. Much lower overall heat input and shorter duration time in laser-beam soldering results in very high cooling rates (~104 K/s) from peak temperature compared with the infrared reflow soldering (~10 K/s). Therefore, in laserbeam solder joints, much finer eutectic Ag3Sn intermetallic in pure tin and thinner Cu-Sn intermetallic layers at copper pad and wire interfaces form compared with the infrared

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115

reflowed joints. The high peak temperature (673 K to 873 K) caused by laser irradiation compared to the infrared reflow (~528 K) dissolves a certain amount of copper from leads and pads. This results in the formation of a considerable number of equiaxed Cu6Sn5 dendrites in the bulk solder. The difference of microstructure between laser-beam soldered and infrared reflowed joints causes higher microhardness of the laser-beam soldered joints (~27 Hv) than that of infrared reflowed joints (~16 Hv), suggesting better mechanical properties under static, creep and fatigue loading. Laser-beam processes have also been applied to bond Tape Automated Bonding (TAB) leads to integrated circuit pads (Spletter and Crowley 1990). Since laser inner lead bonding can be performed using lower temperature and lower pressure compared to conventional gang bonding process, it can minimize the thermal and mechanical stressinduced damages of circuit. TAB laser bonding is performed by positioning a focused laser-beam over a lead and then pulsing the laser once, creating a reliable metallurgical bonding between the plated lead and the bump (Spletter and Crowley 1990). This bonding process takes advantage of a thin tin plating to absorb the laser energy and accelerates melting of the copper lead and gold bump. Since this process achieves rapid alloying of copper and gold, bonds without the brittle intermetallic compounds often found in gold/tin alloys formed with conventional methods can be obtained (Spletter and Crowley 1990).

5.8. C O N C L U S I O N S Lasers are lights that obey straightforward physical laws and make it easy to predict their behavior. Knowing the relevant optical equations and how to apply them will explain most of the features of laser materials processing. For example, the high energy density available from focused sources is utilized to melt only the surface of a material in a very short time. This characteristic can apply to form non-equilibrium materials. In other cases like EXCIMER laser, it is suitable for applications in the field of micromachining. Although laser processing is prominent, the user needs to consider the following problems. What is the best quality that one can achieve and what is the fastest way to get there? What will be new and different are the orders of magnitude differences between the feature sizes and quality. These differences will demand better optics, stable environments, and system design, part handling and inspection systems.

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Sameshima, T., Usui, S. and Sekiya, M. (1986) IEEE Electron Device Lett., 7, 276. Sato, T., Yamada, T. and Ozawa, T. (1985) in Rapidly Quenched Metals V, eds. Steeb, S. and Warlimont, H. (Elsevier Science Pub., Amsterdam), p. 1643. Scandella, L., Staufer, U., Brodbeck, D., Reimann, P., Gtintherodt, H. J., Zehringer, R., Hauert, R. and Moser, E. M. (1991) Mater Sci. Eng., A133, 601. Schaeffler, A. L. (1949) Metal Prog., 56, 680. Schwan, J., Dworschak, W., Jung, K. and Ehrhardt, H. (1994) Diamond Relat. Mater., 3, 1034. Sekhar, J. A. and Mehrabian, R. (1981) Metall. Trans., 12B, 411. Semerad, E., Musiejovsky, L. and Nicolas, J. (1993) J. Mater. Sci., 28, 5065. Sera, K., Okumura, E, Uchida, H., Itoh, S., Kaneko, S. and Hotta, K. (1989) IEEE Trans. Electron Devices, 36, 2868. Shimizu, K., Sugiura, O. and Matsumura, M. (1993) IEEE Trans. Electron Devices, 40, 112. Singh, J. and Mazumder, J. (1987a) Metall. Trans., 18A, 1987. Singh, J. and Mazumder, J. (1987b) Acta Metall., 35, 1995. Spaepen, E (1975) Acta Metall., 23, 729. Spaepen, E and Turnbull, D. (1976) in Rapidly Quenched Metals, eds. Grant, N.J. and Giessen, B. C. (MIT Press, Cambridge), p. 205. Spaepen, E and Turnbull, D. (1982) in Laser Annealing of Semiconductors, eds. Poate, J. M. and Mayer, J. W. (Academic Press, New York), p. 15. Spletter, P. J. and Crowley, R. T. (1990) in Proc. Electronics Components & Technology Conf., p. 757. Tabata, H. and Kawai, T. (1997a) Appl. Phys. Lett., 70, 321. Tabata, H. and Kawai, T. (1997b) IEICE Trans. Electron, E-80-C, 918. Tabata, H., Kawai, T. and Kawai, S. (1993) Phys. Rev. Lett., 70, 2633. Tabata, H., Tanaka, H. and Kawai, T. (1994) Appl. Phys. Lett., 65, 1970. Tabata, H., Tanaka, H. and Kawai, T. (1995) Jpn. J. Appl. Phys., 34, 5146. Trivedi, R., Magnin, P. and Kurz, W. (1987) Acta Metall., 35, 971. Tsu, R., Hodgson, R. T., Tan, T. Y. and Baglin, J. E. (1979) Phys. Rev. Lett., 42, 1356. Tsukamoto, H., Yamamoto, H., Noguchi, T. and Suzuki, T. (1992) Jpn. J. Appl. Phys., 31, L659. Tsukamoto, H., Yamamoto, H., Noguchi, T., Masuya, H. and Suzuki, T. (1996) Jpn. J. Appl. Phys., 35, 3810. Turnbull, D. (1950) J. Appl. Phys., 21, 1022. Uenishi, K. and Kobayashi, K. E (1993) in Proc. 3rd Japan Internat. SAMPE Symposium, p. 1038. Uenishi, K. and Kobayashi, K. E (1994) in Proc. of the 4th lnternat. Conf. on Aluminum Alloys, (Atlanta, GA), p. 714. Uenishi, K. and Kobayashi, K. E (1996a) Solidification Science and Processing, eds. Ohnaka, I. and Stefanescu, D. M. (TMS, Warrendale), p. 231. Uenishi, K. and Kobayashi, K. E (1996b) Intermetallics, 4, $95. Uenishi, K., Sugimoto, A. and Kobayashi, K. E (1992a) Z. Metallkde., 83, 241. Uenishi, K., Sugimoto, A. and Murakami, T. (1992b) in Proc. Int. Conf. on Laser Advanced Materials Processing (Nagaoka, Japan), p. 807. Uhlman, D. R. (1972) J. Non-Cryst. Solids, 7, 337. Vitek, J. M. and David, S. A. (1994) Laser Materials Processing-IV, eds. Mazumder, J., Mukherjee, K. and Mordike, B. L. (TMS, Warrendale), p. 153. Vitek, J. M., Das Gupta, A. and David, S. A. (1982) Metall. Trans. 13A, 1141. Vitta, S., Greer, A. L. and Somekh, R. E. (1994) Mater. Sci. Eng., A179/A180, 243. von Allmen, M. (1983) in Glassy Metals 1I, eds. Beck, H. and Gtintherodt, H.-J. (Springer-Verlag, Berlin), p. 261. von Allmen, M., Lau, S. S., Maenpaa, M. and Tsaur, B. Y. (1980) Appl. Phys. Lett., 37, 84. Yang, W., Messier, Jr., R. W. and Felton, L. E. (1995) Weld. J. 74, 224-s. Yano, T., Ooie, T., Yoneda, M. and Katsumura, M. (1996) J. Mater. Sci. Lett. 15, 1994. Yee, A. L., Ong, H. C., Xiong, F. and Chang, R. P. H. (1996) J. Mater. Res. 11, 1979. Young, R. T. and Wood, R. E (1982) Ann. Rev. Mater. Sci., 12, 323. Young, R. T., Leeden, G. A., Narayan, J., Christie, W. H., Wood, R. E, Rothe, D. E. and Levatter, J. I. (1982) IEEE Electron Device Lett., 3, 280. Zabinski, J. S., Donley, M. S., Prasad, S. V. and McDevitt, N. T. (1994) J. Mater. Sci. 29, 4834. Zambon, A. and Bonollo, E (1994) Mater. Sci. Eng., A178, 203. Zarubova, N., Wolf, P., Cermak, J. and Cernansky, M. (1996) J. Mater. Sci. 31, 137.

Chapter 6 Thermal Plasma Processing 6.1.

Introduction 6.1.1 Advantages of Plasma Processing 6.2. Thermal Plasmas 6.2.1 Principles of Plasma Generation 6.2.1.1 DC Plasma Torches 6.2.1.2 AC Plasma Torches 6.2.1.3 RF Plasma Torches 6.2.2 Plasmagen Gases 6.2.3 Plasma-Particle Interaction 6.2.4 Plasma Processing Systems 6.3. Processing of Materials 6.3.1 Plasma-Spraying 6.3.1.1 Structure of Sprayed Deposits 6.3.1.2 Plasma Spray-Deposited Materials 6.3.1.3 Low-Pressure Plasma-Spraying 6.3.1.4 Reactive Plasma-Spray Forming 6.3.1.5 Plasma Spheroidization 6.3.2 Plasma Reactors and Furnaces 6.3.2.1 Plasma Decomposition 6.3.2.2 Plasma Metallurgy 6.3.2.3 Processing of Ceramics 6.3.2.4 Treatment of Hazardous Wastes 6.3.3 Processing of Metastable Phases 6.3.3.1 Plasma Deposition of Diamonds 6.3.3.2 Thermal Plasma Synthesis of Ultrafine Alumina Powder 6.4. Summary and Conclusions Acknowledgments References

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Chapter 6 Thermal Plasma Processing R V. ANANTHAPADMANABHAN and N. VENKATRAMANI

6.1. INTRODUCTION The twentieth century, variously termed at different times the semiconductor age, the silicon age, and the oxide age, has seen some breathtaking achievements in the development of new materials. These achievements have been made possible by matching developments in the material processing technologies. Today's advanced industries need precisely tailored materials to meet stringent operational requirements. However, a close look at the conventional manufacturing technology tells us that it is beset with barriers, both theoretical and man-made. Table 6.1 summarizes the limitations of the current material processing technologies and how they can be overcome in the emerging technologies of the future. It may be seen that many of the shortcomings of conventional technologies can be overcome by the use of electron, plasma and laser beams. In particular, thermal plasma processing has the capability to transcend thermodynamic and kinetic barriers by opening up new reaction paths inconceivable in the conventional sense. This technology, which until recently was confined to the realms of laboratory, is fast emerging as a major technique not only to synthesize and purify materials, but also to modify and shape surfaces (Boulos 1996, Venkatramani 1995, Kong and Lau 1990, Pauloski 1995). Plasmas of technological interest can be classified into two categories: nonequilibrium plasmas and thermal plasmas (Fauchais et al. 1983, Fineman 1987). Nonequilibrium or cold plasmas, more popularly known as glow-discharge plasmas, are low-pressure plasmas characterized by high electron temperatures and low ion and neutral temperatures. They are widely used in lighting, surface cleaning, etching, film deposition and polymerization. Thermal plasmas are characterized by the electron temperature being equal to the gas temperature. Normally, plasmas in the temperature range of 2,00030,000 K and with charged particle density of 1019-1021 m -3 are termed thermal plasmas. Thermal plasma processing has emerged as a major processing route to synthesize many alloys and ceramics of industrial importance. The different categories of plasmas and their applications are listed in Table 6.2.

6.1.1 Advantages of plasma processing The high temperature, high enthalpy, and fast quench rate which are characteristics of the process, offer unlimited scope to the materials scientist to synthesize newer ma121

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P V. Ananthapadmanabhan and N. Venkatramani

Table 6.1. Limitations of current material processing technologies Barrier

Approach to overcome the barrier

Emerging technology

Thermodynamic barrier

Processing in non-equilibrium and thermal plasmas

Spatial barrier

Access to non-equilibrium regimes, use of exotic intermediates and precursors Transfer heat to volume rather than across surface Micro- and nanoscale processing

Temporal barrier

Nano- and picosecond processing

Purity barrier

Absolute chemical control and limit impurity content

Heat transfer barrier

and

specificity

Ecological barrier

One-step eco-friendly technology, conserving materials and recycling wastes

Electromagnetic processing Beam processing; electron and laser beams and low-pressure plasma processing Laser processing, thermal plasma processing with rapid quench Processing by electron, plasma and laser beams and low-pressure plasmas Thermal plasma processing

Table 6.2. Plasmas of technological interest Category of plasma

Applications

Low pressure plasmas (10-4-10 -2 torr)

Sputtering and surface modification processes, plasma source for ion implantation Etching, microelectronic processing Plasma chemistry, plasma polymerization

Medium pressure plasmas (10-2-1 torr) Subatmospheric pressure plasmas (1-100 torr) Atmospheric plasmas (100+ torr)

Plasma-spraying, plasma-melting, material synthesis

terials with improved properties. The temperature encountered in the plasma medium is 10,000-15,000 K, whereas the maximum temperature available in conventional heat sources is 3,000 K. This enables the adoption of novel approaches to accomplish chemical reactions and metallurgical processes, inconceivable in the conventional sense. The high-temperature gas phase chemistry in the plasma medium promotes various chemical reactions at much faster rates than those encountered in conventional processing techniques. The high quench rate, typically of the order of 105-107 K/s, prevents recombination of the products and favors homogeneous nucleation resulting in nano-sized particles. Furthermore, it offers flexibility in the choice of reactants, which can be solids, liquids or gases. Plasma processes are, usually, one-step processes, with the capability of handling large throughputs in small reactor volumes in relatively short processing times. They also do not substantially contaminate the process environment, thereby ensuring high purity of the product.

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123

Plasma processing has been applied to synthesize a variety of alloys, intermetallics, and refractory ceramics. Plasma-spray technology is a well-established surface treatment technique to prepare protective coatings for a variety of industrial applications. The focus of this chapter will be on thermal plasma processing of materials. The organization of the chapter is as follows: Introduction to the subject of thermal plasma processing will be followed by a description of the different techniques of plasma generation. Applications of thermal plasmas in material processing are discussed in detail in Section 6.3. Plasma synthesis of ultrafine alumina is discussed in detail towards the end of this section to illustrate the formation of non-equilibrium phases in reactions occurring in plasma medium.

6.2. T H E R M A L

PLASMAS

The classical plasma is a fully ionized gas consisting of electrons and ions that interact with each other only through Coulomb forces (Eliezer and Eliezer 1989). This model is valid only for astrophysical or fusion plasmas; it does not hold good for industrial plasmas, which are partially ionized gases and are dominated by atomic and molecular processes. Thermal plasma is a viscous, electrically and thermally conducting fluid. Plasma retains many of the properties of gases, and behaves in conformance with the physical laws valid for gases. The specific properties of plasma that distinguish it from a gas become apparent in the presence of a magnetic field, when the plasma acquires non-isotropic properties. Thus, the fundamental difference between plasma and a non-ionized gas lies in their responses to electromagnetic forces. The electrically charged particles present in plasma are affected by externally applied electric and magnetic fields, and also interact with one another. The electric field that they set up is so extensive that every particle is affected by a multitude of other particles. Consequently, the collective behavior of the particles is much stronger in plasma than in a non-ionized gas. The unique feature of the thermal plasma jet that distinguishes it from other heat sources is its high power density. The energy density of thermal plasma devices is of the order of several GW/m 2, which is 10-100 times the power density of conventional oxy-fuel flames. Although higher power densities are obtainable with electron and laser beams, the source strengths available with these devices, especially laser devices, are not high enough for large-scale material processing applications.

6.2.1 Principles of plasma generation Thermal plasmas for material processing applications are usually of electric origin. Laboratory thermal plasmas are produced in devices known as plasma torches or plasmatrons (Heberlein 1992). Plasma torches are electrothermal devices, which convert the electrical input energy to thermal energy of the plasma. Depending on the primary source, which can be direct current (DC), alternating current (AC) at mains frequency, or at radio frequency (RF), they are known as DC, AC or RF torches. There are also combinations of these torches known as hybrid devices. Figure 6.1 gives a schematic of the different types of plasma torches.

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Figure 6.1. Schematicof different plasma sources: (a) DC plasma torch--rod cathode type (non-transferred arc mode); (b) DC plasma torchmrod cathode type (transferred arc mode); (c) DC plasma torch--hollow cathode type; and (d) RF plasma torch.

6.2.1.1 DC plasma torches The most commonly used plasma device for material processing applications is the DC arc plasma torch. A direct current arc is struck between a cathode rod and a nozzle anode and the energy from the arc is abstracted by the plasma gas injected in the inter-electrode region. Forced gas flow extends the arc in the anode nozzle which is strongly water-cooled (Fig. 6.1(a)). A thermal arc pinch effect is produced by the joint action of the cold wall arc channel and the cold gas sheath around a very high-temperature conducting core (the arc column). The DC arc in a plasma torch needs to be stabilized, i.e., it should remain stationary against fluctuations. This is often done by constricting the arc to a well-adjusted narrow high-temperature, highly conducting arc column. Various torch configurations are possible depending upon the arc stabilization mode: tangential vortex gas input in the arc channel, axial gas input along the cathode, segmented anode arc and magnetic stabilization (Fauchais et al. 1983). The magnetic field can be self-induced (by an arc current greater than 8000 A), or externally generated. The electric arc is stabilized by using a constricted anode nozzle and by the resultant aerodynamic effects in the streaming plasma gas. Stabilizing action of the vortex gas flow provides a cold boundary layer near the anode wall so that heat loss to the wall is reduced. This results in the thermal energy being highly concentrated with improved torch stability and efficiency. The localized heat flux at the nozzle can be as high as 160 W/mm 2. The obvious choice of material for the anode is copper, although molybdenum and graphite are also used. Depending on the nature of the gas and the working parameters, the anode losses range

Thermal Plasma Processing 3

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Figure 6.2. Cross-section of a typical DC plasma torch: (1) cathode; (2) anode; (3) anode segment; (4) insulator; (5) cathode segment; (6) water inlet; (7) plasma gas inlet; and (8) water outlet.

between 80 and 10% of the energy input in the arc. The anode losses are proportional to the current density and are a function of the arc voltage. The cathodes can be of the thermionic type such as tungsten, carbon or molybdenum, which, obviously, must be used in non-oxidizing atmosphere. Under certain conditions of oxidizing atmosphere, one can use zirconium or hafnium cathodes. Heat losses at the cathode are generally quite low (less than 10% of the input power). For a tungsten cathode tip, the erosion rate has been estimated to be of the order of 10 -9 to 10-10 g per Coulomb, so that its lifetime is about 150 h. Non-thermionic cathodes are normally made of copper. The cross-section of a typical DC plasma torch is shown in Fig. 6.2. This design is based on a rod type cathode and nozzle-shaped anode with tangential gas entry through the insulator module. When a gas is injected into the electrode gap and a high intensity current is passed, a DC arc is established between the electrodes. The plasma gas extracts energy from the arc and emerges out of the nozzle as a high-temperature, high velocity jet. The temperature at the core of the plasma jet ranges between 15,000-20,000 K (Gerdeman and Hecht 1972). The typical operating parameters of a low power, laboratory DC plasma torch (non-transferred arc mode) are given in Table 6.3 The plasma torch can be operated in a transferred or non-transferred mode, depending on whether the arc is electrically transferred to the work piece or not, as illustrated in Fig. 6.1 (a) and 6.1 (b). The transferred arc mode is mainly used for cutting metal sheets, refining, and remelting operations. The non-transferred arc mode is used in plasma-spraying and other material processing operations.

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Table 6.3 Typical operating parameters of a low-power laboratory DC plasma torch (non-transferred arc mode) Arc voltage (V) Arc current (A) Arc gas Torch input power (kW) Water Inlet temperature ((C) Water outlet temperature ((C) Cooling water flow rate (1/min) Power loss at the electrodes (kW) Flow rate of Argon (g/s) Cold gas flow velocity (m/s) Thermal energy of emerging plasma (kW) Enthalpy of plasma (kJ/g) Plasma temperature at the nozzle exit (K) Plasma exit velocity (m/s)

40 250 Argon 10 33.2 36.3 15 3.2 1.19 13.27 6.8 5.71 10,600" 468.9*

* Calculated.

6.2.1.2 AC plasma torches When alternating current at mains frequency is used in a two-electrode device, each electrode alternately acts as the cathode and anode. In each half cycle a DC discharge is established once the breakdown potential is crossed and the discharge is maintained so long as the applied voltage remains above the maintenance potential. When the voltage falls below the maintenance potential, the discharge gets extinguished, and the space charge gets swept away. The continuous alternation of the polarity of the electrodes limits strongly the stability of the discharge and the choice of the electrode materials.

6.2.1.3 RF plasma torches Radio-frequency (RF) inductive or capacitive coupling (between 1-20 MHz) can be used to generate a plasma column, which can be stabilized in a quartz tube. The main advantage of this technique is its electrodeless feature, allowing the use of aggressive gases such as oxygen, chlorine (Mackinnon and Reuben 1975) or UF6, etc. which cannot be used as plasmagen gases in DC plasma torches. The minimum power necessary to sustain the discharge is determined by the nature of the gas, its pressure and the frequency of the electromagnetic field (Rykalin 1978). The gas velocity in an RF torch is several tens of m/s, which is much lower than that encountered in DC torches (hundreds of m/s) (Dembovski 1985). Furthermore, the diameter of the reactor is important (20-200 mm) and the energy density is much lower than in arc plasma generators. RF plasma torches capable of operating up to power levels of 1 MW (Ionarc Tafa torches) and a lifetime of a few thousand hours are presently employed for material processing applications. However, the electrothermal efficiency of these torches is low (40-50%) and are mostly limited to specialized applications.

Thermal Plasma Processing

,•,4

11,S00K

10,500K

9,600 K

8,600 K

7,600 K

127

6,800 K

E2

0

4

8

12

16

20

24

Axial Distance(mm) Figure 6.3. Axial and radial temperature distribution in a nitrogen plasma jet (Fauchais et al. 1983).

The arc column can be extended by using hybrid torches, which consist of a combination of torches connected in tandem. The extended hot zone increases the residence time and ensures that the reaction goes to completion.

6.2.2 Plasmagen gases The most commonly used gases for plasma generation are argon, nitrogen, helium, hydrogen, and air. Plasma gas flow rate and the electric power to the plasma torch must be properly balanced in order to get a stable arc. The choice of plasma gas depends on many factors, such as the design features of the torch, in particular the electrode materials (Venkatramani 1995). In the case of plasma torches employing tungsten cathode, the choice of the plasma gas is limited to inert gases and non-oxidizing gases. Gas enthalpy is another important factor deciding the choice of the gas (Gerdeman and Hecht 1972). The energy content of nitrogen and hydrogen, which are diatomic, is considerably higher than that of argon or helium. This is due to the dissociation reaction in the case of nitrogen and hydrogen prior to ionization. The heat transfer coefficient is also higher, enabling efficient plasma-particle heat transfer. The low cost and high internal energy of nitrogen makes it the most commonly used gas. If a completely inert atmosphere is required, argon is usually preferred. Reactive gases like hydrogen, oxygen, chlorine, methane and ammonia/nitrogen can be used to impart reducing, oxidizing, chloriding, carburizing, or nitriding effects, respectively, to the plasma. The axial and radial temperature fields in the plasma jet issuing out of a nontransferred arc torch are shown in Fig. 6.3. It is seen from the figure that the temperature falls off rapidly downstream the axis and radially away from the centre-line. The maximum temperature and the temperature profile depend on the input power, nature and flow rate of the plasma gas and the torch design. But the significant point to be noted is the steep gradient in the axial and radial temperatures. Particle injection into the plasma stream is not a trivial problem. In fact, it can be very difficult due to the high viscosity of the plasma. Precursor powders are usually entrained

128

P V. Ananthapadmanabhan and N. Venkatramani

in a carrier gas for injection into the plasma. The carrier gas can be inert or reactive. There is an increasing tendency to use liquid reactants, which is known as liquid injection plasma synthesis (LIPS) (Kong and Lau 1990).

6.2.3 Plasma-particle interaction The heat transfer process from the plasma to the particles is of fundamental importance in all plasma processing applications. Plasma-particle heat transfer is very complicated and depends on the thermo-physical properties of the material as well as the enthalpy and composition of the plasma. Many theoretical models (Fauchais et al. 1983, Bonet 1976) have been developed for the plasma-particle heat transfer phenomenon. The heattransfer process is decided by the thermal boundary at the particle-plasma interface and the properties of the gases in this boundary layer, which is characterized by steep temperature gradient and strong deviations from chemical and kinetic equilibrium (Pfender 1994). The heat transfer coefficient to the particle may be calculated from a correlation based on the well-known Ranz-Marshall equation (Ranz and Marshall 1952): Nu = 2.0 + 0.6Re~

~

(6.1)

where Nu, Re and Pr are the Nusselt, Reynolds and Prandtl numbers, respectively. Although this equation must be corrected by additional factors to account for the strong variation of gas properties in the boundary layer, it does give a fair estimate of the heat transfer coefficient. The coefficient of heat transfer between the plasma gas and the surface of the particle can now be calculated from the definition of the Nusselt number: Nu - hdp/k

(6.2)

where h is the heat transfer coefficient, k is the thermal conductivity of the plasma, and

dp is the diameter of the particle. The goal pursued while optimizing the process parameters is to ensure that all particles entering the plasma jet are heated to the desired temperature, which may be the melting point in the case of plasma-spraying or the reaction temperature in chemical synthesis or dissociation reactions. If the particles are small or if they conduct heat well, this objective is easier to achieve. On the other hand, in the case of large or poorly conducting particles, it might be necessary to choose processing conditions that lead to intense evaporation from the surface of the particles. The large thermal gradients within the particle may also lead to partially molten particles, which lead to porosity in the coatings. Numerous onedimensional and elaborate two-dimensional models have been worked out for plasmaparticle heat transfer problem. In the case of good heat-conducting particles, the equation describing particle melting in the plasma is given by (Pauloski 1995):

telqd2h(Tg - Tp)

-

3

1/61-IdpppCp{(Tm

- 300) -k- A H f

}.

(6.3)

The left-hand side of Eq. (6.3) describes heat transferred from the flame to the particle,

Tg is the gas temperature and Tp is the particle temperature, which is assumed to be uniform in the entire volume of the particle. The right-hand side describes the heat absorbed

Thermal Plasma Processing

129

by the particle, pp and Cp being the density and specific heat of the particle respectively. A Hu is the latent heat of fusion of the material per unit volume and te the time for heating the particle from room temperature to its melting point. It is assumed that the energy transferred by radiation from the plasma is small.

6.2.4 Plasma processing systems Plasma-processing encompasses a wide range of operations including plasma-cutting and welding, plasma-melting and machining, plasma-spraying, chemical synthesis and mineral dissociation. A wide variety of plasma reactors have been developed based on DC non-transferred and transferred arc plasma torches. The various types of plasma reactors used for material processing applications have been described by Fauchais et al. (1983). The exact design of the reactor set-up depends on the specific application, but the general features remain more or less the same. The schematic of a typical plasma reactor developed in our laboratory for material synthesis is shown in Fig. 6.4. The major subsystems are plasma torch, powder-feeder, reaction chamber, torch power supply and control console. The reaction chamber consists of double walled water-cooled stainless steel segments provided with sampling and diagnostic ports. The plasma torch is mounted on one end of the reaction chamber as shown in the figure. There is also provision to collect the product and feed any reactive gas. The power sources generally employed are DC motor generator sets, transductor controlled rectifier units or silicon controlled rectifier (SCR) system. The SCR systems are currently the most extensively used power sources. A high frequency (HF) igniter to initiate the plasma is also included in the power supplies. The HF unit is connected in parallel with the DC power supply, which must be protected from it. The control console has instruments for the measurement and control of the torch inputs (viz. electrical, cooling water, gas and powder feed rates). Powder is injected into the plasma by means of the powder feeder. The turn-table-feed type powder feeder is commonly used. This ensures smooth volumetric flow of powder and has very good control over the feed rate and requires low carrier gas. 6.3. P R O C E S S I N G OF M A T E R I A L S

Plasma processing is an emerging technology that utilizes the exotic properties of the plasma medium to effect physical, chemical or metallurgical reactions to produce new materials or impart new functional properties to conventional materials. Plasma processing has been successfully applied to develop advanced ceramic coatings, synthesis of nanocrystalline materials, processing of minerals and ores, and treatment of hazardous wastes.

6.3.1 Plasma-spraying Plasma-spraying is a surface modification technique that combines particle melting, rapid solidification, and consolidation in a single process. This technique has been widely used to develop protective coatings of ceramics, alloys, and composites to enhance the surface properties of critical components operating in severe environment (Pauloski 1995,

130

P. V. Ananthapadmanabhan and N. Venkatramani 5

6

7

8

=

=

QIO

OlO

['~

Figure 6.4. Schematic of thermal plasma reactor for material synthesis: (1) cooling water; (2) control console; (3) power supply; (4) powder feeder; (5) plasma torch; (6) reaction chamber; (7) sampling duct; (8) powder collector; and (9) auxiliary instruments.

Gerdeman and Hecht 1972). In conventional plasma-spraying, the material to be coated is introduced in the plasma jet in powder form by means of a carrier gas. The powder particles, as they enter the plasma jet, are heated and melted and the molten droplets are accelerated to very high velocities, as high as 100-400 m/s. As these molten droplets strike the substrate surface, they flatten and get anchored to the surface irregularities to form an adherent coating. Plasma-spraying has several advantages over other coating processes. Because the high temperature generated by the plasma arc is far beyond the melting point of any known material, any material that melts and remains in the liquid state over a distinct range can be spray-coated. Although the plasma temperature may be in the range of 7,00010,000 K, the substrate temperature can be kept low (below 400 K), avoiding the risk of substrate distortion. The spray technique does not impose any restriction on the workpiece dimensions and large samples can be coated. The adherence of the coating to the substrate surface is predominantly mechanical in nature (Pauloski 1995). When the molten particles impact a roughened surface, they get interlocked to the surface irregularities resulting in a mechanical adherence, which is believed to be the predominant mechanism in plasma-sprayed coatings. However, depending on the nature of the substrate material, its temperature and the impinging particles, a metallurgical bond can result as in the case of nickel aluminide or molybdenum on steel substrates.

Thermal Plasma Processing

9

Figure 6.5.

,

~

,

131

-

SEM photograph of a typical sample of plasma-sprayed coating of alumina.

The coating quality is affected by a number of interdependent experimental parameters, which need to be carefully controlled. Thermal interaction of the particles with the plasma jet, the particle trajectory and the dwell time are important factors influencing coating adhesion and other properties. Particle size decides the particle trajectory and degree of melting, which controls the porosity and adhesion of the coatings. As it is practically impossible to have monosized powders, particle size in a narrow range is selected. Although the exact particle size depends on the nature of the material to be spray-coated, the recommended size for ceramic powder is 3 0 - 4 0 / z m and for metal powders 40-60 /~m.

6.3.1.1 Structure of sprayed deposits The sprayed layer has a typical lamellar structure, built up by successive layers on the material (Gerdeman and Hecht 1972). Heat transfer calculations show that freezing of the particles occurs in a few microseconds and that complete solidification occurs in perhaps 100/zs. The zone of thermal effect in the underlying material is, therefore, quite small and temperature gradients reach 105 K/cm. The wetting and flow properties of the liquid droplets are important since they influence the porosity within the coating and at the substrate interface. The surface morphology of a typical sample of plasma-sprayed alumina is shown in Fig. 6.5. The deposit shows grains consisting of melt-quenched particles with interganular and intragranular porosity, which are characteristic of plasma-sprayed coatings. The porosity of the coating is about 15-20%, which can be considerably reduced by optimizing the experimental parameters and carrying out the coating operation under reduced pressure.

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The flow and solidification of molten particles on impact with a cold surface is a difficult problem to treat theoretically because of the interaction of heat transfer and crystal growth. The very high velocity at which the various processes occur also makes direct experimental observation extremely difficult. Madejski (1976) has carried out a theoretical treatment of the impact of a molten droplet on a cold substrate, taking into account surface tension, viscosity and crystallization effects. The results show that the degree of flattening (D/d), defined as the ratio of the diameter of the flattened disc (D) to that of the original drop (d), is related to the density p, viscosity/z and impact velocity of the drop v by the following equation: D / d = 1.29(pvd/#) ~

(6.4)

Substituting appropriate values for A1203, in the velocity range of 100-400 m/s gives a value of 3-6 for the degree of flattening, which seems to be in reasonable agreement with various experimental observations (Zolotowski 1968). This suggests that at least for alumina, flattening is completed before crystallization becomes significant and the two processes can largely be treated separately. Plasma-spraying conditions are very similar to those occurring during splat-quenching of metals which has been widely studied. The cooling rate has been of particular interest and various techniques have been used to estimate the thermal conditions as a molten particle spreads and solidifies (Jones 1973). The main results of this work show that the cooling rate is largely independent of substrate material and is controlled by the interface between substrate and particle if the flattened particle has a thickness of less than a few #m. The estimated heat transfer coefficient under these conditions is of the order of 105 W m -2 K -1 . Applying this result to the case of plasma sprayed A1203 gives a cooling rate of the order of 107 K/s and a freezing time of about 10 #s. This also suggests that there is a large temperature gradient across the interface between the particle and the substrate (or particle and previously deposited layer) and the temperature of the underlying material remains relatively low. This is particularly so if the substrate is cooled during spraying by an air blast to prevent heating by plasma gases and to remove the sensible heat and heat of fusion of the sprayed particle. 6.3.1.2 Plasma spray-deposited materials Applications for plasma-sprayed materials range from microelectronics to large engineering structures. There are excellent reviews and publications (Pauloski 1995, Gerdeman and Hecht 1972, Zaat 1983, Fauchais and Vardelle 1994, Notomi and Hisatome 1996, Boulos 1992, Matejka and Benko 1989) exclusively devoted to the coating process and applications, which can be referred to for further details. Plasma-sprayed coatings of metals, alloys and ceramic materials have been used for various applications. Plasma-spraying has been applied to produce wear-resistant coatings of ceramics and metals on various components. Successful wear-resistant coatings have been developed using cobalt- and nickel-based alloys containing carbides of chromium, tungsten and titanium as abrasion-resistant particles (Rangaswamy and Herman 1986, Mohanty and Smith 1995, Knight et al. 1995). Self-lubrication can be provided by incorporating low-friction materials or dry lubricants such as graphite, molybdenum disul-

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phide, or organic lubricants. Plasma-spraying has been effectively used to protect metallic and non-metallic components from high temperature and oxidation. Plasma-sprayed thermal barrier coatings consisting of multi-layered and graded ceramic and metal coatings are being extensively used for protecting the surfaces of metal components, gas turbines and other critical engine parts exposed to corrosive environment. Another important application of plasma-spraying is reclamation, where it is used to repair worn out or badly machined components. Worn out surfaces of brake drums, grinders, spindles, etc. can be reclaimed by plasma coating. Oxide coatings are used for many high-tech as well as conventional applications. Plasma-sprayed coatings of stabilized zirconia, alumina, and alumina-titania have been extensively studied (Bratton and Lau 1981, Stecura 1977, Ananthapadmanabhan et al. 1991, Ramachandran et al. 1997). Aluminium oxide (often mixed with titania) and chromium oxide are used on many textile machinery components such as thread guides (Matejka and Benko 1989). They are also widely used on pumps and mechanical seals to provide resistance to both abrasion and corrosion. Plasma-sprayed zirconia coatings, as thermal barriers, are effective in improving the efficiency of gas turbine engines. The preferred material is 6-8 wt.% Y203-ZrO2 or 24 wt.% MgO-ZrO2 deposited over Ni-CrA1-Y bond coat applied by vacuum plasma-spraying (Meetham 1988). The latter provides oxidation resistance to the superalloy substrate and is also strain tolerant, which helps to accommodate stresses arising out of thermal expansion mismatch between the base metal and the ceramic coating. Thermal shock resistance of zirconia coatings has been greatly enhanced by using composite coatings of alumina and calcia stabilized zirconia. The improved thermal shock behavior is attributed to destabilization of calcia stabilized zirconia leading to the generation of microcracks which act as thermal stress relief centers (Ananthapadmanabhan et al. 1991). When ceramic coatings are applied on metal substrates, the large mismatch in the thermal expansion coefficients of the metal base and the coating material can result in interface stresses. This can lead to poor adhesion and even coating failure. In order to overcome this, an intermediate coating, referred to as bond coat, is applied on the metal substrate before applying the ceramic coating. The primary function of the bond coat is to reduce the interface stress and improve coating adhesion (Matejka and Benko 1989). Nickel aluminide and its derivatives, Ni-Cr-A1-Y, are frequently used as the bond coat material. The reaction of nickel and aluminum to form the intermetallic is highly exothermic, resulting in high adhesion strength. Nickel aluminide coatings are usually prepared by using special composite powder of nickel and aluminum (Matejka and Benko 1989). This composite powder forms the compound during the spray process. A one-step process to form plasma-sprayed Ni3A1 coatings has been developed (Ananthapadmanabhan et al. 1995) by using a blend of nickel and aluminum powders as the feedstock material. The process parameters have been optimized to effect compound formation as the powder particles deposit on the substrate surface. Coating adhesion was measured by tensile adhesion strength as described in ASTM C 633 (Lin and Berndt 1994). The adhesion strength of the coatings was found to be much greater than 60 MPa.

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The cooling rate for plasma-sprayed particles approaches 106 K]s and this produces a fine-grained microstructure within the splat (Scott and Kingswell 1991). This rapid solidification can produce amorphous deposits from some ferrous alloys (Herman and Bhat 1980), while metastable or non-stoichiometric phases can be retained in some ceramics. Phase composition and stability of yttria stabilized zirconia coatings have been extensively studied. The phase structure of yttria stabilized zirconia coatings have been shown to consist of non-equilibrium phases. The work of Miller et al. (1981) and Bratton and Lau (1981) on plasma-sprayed zirconia thermal barrier coatings showed that the coatings consisted of non-equilibrium phases. Irrespective of the starting material, the as-sprayed material was found to contain a non-transformable tetragonal (high yttria) phase. It was observed that partially stabilized zirconia (ZRO2-8% Y203) as well as fully stabilized zirconia (ZRO2-20% Y203) contain primarily the non-equilibrium tetragonal phase, which transformed to the equilibrium cubic and tetragonal (low yttria) during thermal annealing. Plasma-sprayed coatings of alumina and alumina-titania have been shown to consist of metastable alumina (Heintze and Uematsu 1992, Ramachandran et al. 1997). The work of Heintze and Uematsu (1992) has shown that plasma-sprayed coatings of alumina predominantly consist of y-alumina. They also found that u-alumina can be formed by careful optimization of the process parameters.

6.3.1.3 Low-pressure plasma-spraying Atmospheric plasma-spraying (APS) is seldom satisfactory when used to prepare coatings of metals, alloys, and non-oxide ceramics. When coatings of refractory carbides, nitrides and borides are prepared by APS, the high velocity plasma jet acts as an aspirator and sucks in air from the surrounding atmosphere resulting in oxide inclusions, which can seriously affect the coating performance. Studies on the phase composition of TiB2 coatings prepared by APS showed appreciable amount of boric oxide, formed by partial in-flight oxidation of TiB2 particles (Ananthapadmanabhan et al. 1993). The low melting point of B203 (about 800 K) affects the high temperature performance of TiB2 coatings. This problem can be overcome by low-pressure plasma-spraying (LPPS) or vacuum plasma-spraying (VPS), where the coating operation is performed in a vacuum chamber. Alternately, the chamber can be filled with inert gas to minimize coating contamination. The quality of the coatings produced by LPPS is superior to that of those prepared by APS, in terms of coating properties as well as chemical purity. LPPS reduces porosity and improves coating adhesion. Reactive metals like titanium, zirconium and tungsten can be spray-coated with minimum oxide contamination. LPPS has been extensively used to prepare coatings of refractory carbides, nitrides and other non-oxide materials for various applications. Recent developments include the successful application of LPPS to fabricate components for solid oxide fuel cell (SOFC) components (Notomi and Hisatome 1996). Dense electrolyte coatings of yttria stabilized zirconia have been prepared by low pressure plasma-spraying. A 1 kW module has been developed and successfully operated for 3,000 h.

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6.3.1.4 Reactive plasma-spray f o r m i n g Conventional plasma-spraying consists in feeding metal, alloy, or ceramic powders in the plasma jet, where the particles melt and the molten droplets are deposited on to a substrate surface, forming an adherent coating. It is essentially a physical process involving melting and resolidification of the powder. In reactive plasma-spraying, the high temperature plasma environment is utilized to realize inflight chemical reactions by using suitable reactive gases along with the feed powder (Mutasim et al. 1992). The metal or ceramic powders are introduced in the plasma jet along with a suitable reactive gas. The reactant powder particles melt and the molten droplets react with the reactive gas, resulting in the in situ formation of the desired compound, which is deposited on the substrate. This technique has been used to develop refractory carbide coatings. Reactive plasma-spraying of Ti using methane, propylene, or acetylene as the reactive gas has been carried out (Mutasim et al. 1992, Ananthapadmanabhan et al. 1993) to produce Ti-TiC composite coatings. Dissociation of the hydrocarbon to form finely dispersed aerosol of carbon is the first step in the formation of the carbide. These fine particles of carbon subsequently react with the molten droplets of the metal to form the metal carbide, which is deposited on the substrate. By proper selection of the process parameters, composite coatings containing 60 vol.% of TiC dispersions have been obtained. It was also observed that higher substrate temperatures increased the carbide content. A recent development in this field is a one-step process to form Fe-TiC cermet coatings by reactive plasma-spraying of ilmenite (Ananthapadmanabhan and Taylor 1998). The principle involves the reduction of ilmenite by methane in a thermal plasma environment according to the following chemical reaction:

FeTiO3 + 4CH4 = Fe + TiC + 3CO + 8H2.

(6.5)

Ilmenite powder is introduced into the plasma jet using a mixture of methane and argon as the carrier gas and the substrate is kept at about 1200 K. Ilmenite is reduced to iron and titanium dioxide, which is further reduced to its lower oxides and finally to TiC. Fig. 6.6 shows the surface morphology of a typical Fe-TiC coating, revealing excellent inter-particle cohesion and very little porosity. Reactive plasma-spraying can be extended to develop other refractory carbide and nitride coatings by suitable choice of precursor materials and reactive gases. Reactive plasma-spraying is also used for depositing hydroxy apatite (HA) coatings for dental implants. Conventional atmospheric plasma-spraying of hydroxy apatite, Cal0(PO4)6(OH)2 results in the formation of biodegradable phases such as, fi-tricalcium phospate and tetracalcium phosphate by the thermal dissociation of HA, rendering it unsuitable for biomedical applications (Hench 1991). The decomposition of HA to biodegradable compounds has been eliminated by a reactive plasma-spray technique developed by Patil et al. (1996).

6.3.1.5 Plasma spheroidization Powders of metals, alloys and ceramics can be melted in a plasma jet. Melting of a particle results in the formation of a spherical drop under the action of the surface tension forces and this shape is usually retained after solidification, thus providing the name to

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Figure 6.6. SEM photograph of Fe-TiC coating formedby reactive plasma-spraying.

the process. This treatment may be used simply to give a spherical particle shape for particular application, such as Fe304 for photocopying, plasma spray-quality powders for thermal spray applications and UO2 for dispensed nuclear fuels. Spheroidization is particularly useful to prepare spray-quality powders of special materials for thermal spray applications. Spheroidization of a powder blend of Ni-15% A1 has been carried out in a thermal plasma reactor (Ananthapadmanabhan et al. 1997). The as-collected powder showed poor crystallinity as indicated by the broad X-ray diffraction pattern (Ananthapadmanabhan et al. 1995). Annealing at 800 K in an atmosphere of flowing argon resulted in the formation of well-crystalline Ni3A1. The process parameters have been optimized to yield a precursor powder of Ni-A1, which forms Ni3A1 on the substrate surface during plasma-spraying. A typical photomicrograph of plasma spheroidized Ni-A1 powder is shown in Fig. 6.7. Besides possessing spherical morphology, the powder has excellent flow characteristics, making it ideally suited for thermal-spray applications. Spheroidized particles may also be prepared by atomization of rods or wires fed into a plasma. The diameter of the particles produced depends on the diameter of the torch nozzle, the gas flow rate, plasma density and temperature and surface tension of the material in the liquid state (Rykalin 1976). Alumina particles of 40 # m diameter have been produced by feeding alumina rod in a hydrogen plasma. The heat transfer can be improved if the arc is transferred to the wire but this is not possible with non-conducting materials. An example of large scale industrial application is the spheroidization of magnetite for photocopying applications, where magnetite particles, 125/zm diameter, are produced in 600 kW AC air plasma heater with a power consumption of 2 kWh&g (Fey et al. 1975).

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Figure 6.7. SEM photograph of plasma spheroidized Ni-A1 powder.

6.3.2 Plasma reactors and furnaces

Plasma technology can be gainfully utilized to treat minerals and ores to yield value-added materials. Decomposition processes which are already proved both thermodynamically and technically feasible and favorable include dissociation of carbonates, oxides, sulfides, halides and complex polymetallic ores (Dembovski 1985). It has also been used for metallurgical applications and in the synthesis of advanced ceramics. Table 6.4 shows the high power plasma systems used for diverse applications.

6.3.2.1 Plasma decomposition Decomposition of zircon sand in thermal plasma reactors to produce zirconia is well established. The principle involved is the thermal dissociation of zircon into its component oxides, followed by quenching the reaction products to yield a mixture of ZrO2 and SIO2. Zircon is thermodynamically stable up to 1950 K, above which it spontaneously dissociates into zirconium dioxide and silicon dioxide. The reverse reaction is prevented by freezing the high temperature equilibrium by rapid quenching of the products. Zircon sand is fed into the plasma by means of a powder feeder. As the zircon particles pass through the plasma, they are heated to temperatures above 2,000 K, when dissociation takes place and the product is rapidly quenched. Zirconia is separated from the oxide mixture by leaching away silica with 30-50% solution of sodium hydroxide. Extensive studies on the influence of operating parameters on the dissociation reaction have been carried out by Ananthapadmanabhan et al. (1994a). Results of the study showed that the extent of dissociation of zircon was affected by the power input to the

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Table 6.4 High-powerplasma torches for material processing applications Torch make Huls Linde-Retech Tioxide Aero-Spatiale SKF Ionarc Metco Daido Tetronics Voest-Alpine Chemische Werke Huls AG Voest-Alpine

Power ( M W ) 8.5 0.7 5.0 5.0 7.0 0.35 0.1 1.0 3.0 7.5 8.0 4.0

Applications Chemical synthesis Plasma melting Titanium dioxide Ferro-manganese Recovery from dust Zirconia production Plasma-spraying Scrap melting Tundish heating Steel remelting Synthetic gas for direct reduction of iron ore Ferromanganese

plasma, plasma gas flow rate, carrier gas flow rate, feed rate of zircon, etc. Ionarc Tafa has developed a 300 kW multi-torch furnace for processing zircon sands at 300 kg/h feed rate. The process has been shown to be economically superior to the conventional process. The product purity is superior with an electricity consumption of 1.32 kWh/kg of the product compared to 9.9 kWh/kg for the conventional process (Dembovski 1985). A plasma reactor has also been used to convert ilmenite to titanium dioxide. The process involves reductive dissociation of ilmenite to TiO2 and iron in a thermal plasma medium using controlled amount of methane as the reductant. The process has been extended to prepare Fe-TiC powder (Taylor et al. 1995).

6.3.2.2 Plasma metallurgy The use of RF and DC plasmas for metallurgical applications is getting established in many countries throughout the world. The introduction of plasma furnaces has revolutionized the metallurgical industry by improving the operational efficiency and quality of the products. The conventional iron and steel industry is highly capital-intensive. The blast furnace operation utilizes dwindling resources of expensive high-grade coking coal, which is used for heating and reducing operations. Consumption of this expensive coke can be considerably reduced by augmenting the conventional blast furnace operation with the plasma process (Dembovski 1985). The process consists in partially replacing coking coal by superhot reducing gas synthesized by plasma-heating of a mixture of natural gas and air or carbon dioxide. Superhot reducing gas can reduce the use of coke by as much as 75%, resulting in substantial cost savings. This process is referred to as the direct reduction process and the direct-reduced iron (DRI) is used for producing low-carbon steels. Mintek, South Africa has developed pilot-plant-scale furnaces for producing low-carbon steel by smelting DRI and iron ore (McRae et al. 1985).

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The application of thermal plasma technology to steel-plant dust is a recent development. Dust and metal-bearing wastes from steel plants consist of oxides of various metals and their disposal in landfill sites poses environmental hazards, since toxic metals may be leached into drinking water supplies. Thermal plasma processing of steel-plant dust not only offers an effective way to destroy the hazardous waste, but also to recover valuable metals such as zinc, nickel and chromium. SKF of Sweden has developed the Plasmared processes for recovering metals from metal-bearing wastes in a reducing atmosphere generated by the decomposition of fossil fuels in a plasma stream. The Enviroplas process is a DC plasma arc process, developed by Mintek, South Africa, which can treat arc-furnace dust, alloy-steel dust and other metal-bearing wastes. Production of ferrochrome, ferromanganese and nitrogen alloying of steel are other areas where plasma technology has been extensively used. DC transferred arc furnaces have replaced the conventional submerged arc furnaces for producing ferrochromium from chromite (FeCrO3) ore. The DC arc process ensures higher chromium recovery with reduced levels of phosphorous and sulfur (Jones et al. 1997). Nitrogen alloying of steel by plasma treatment is illustrative of the unique advantages of the process over conventional metallurgical alloying processes. Interstitially dissolved nitrogen improves the mechanical properties and corrosion behavior of steel significantly. The strength and corrosion resistance of some steel grades are commonly enhanced by addition of nitrogen-enriched ferrochrome, ferromanganese or calcium nitride. In conventional induction furnaces with nitrogen partial pressure of 105 Pa, the amount of nitrogen dissolved in steel at 2,500 K is hardly 0.06%, which can be increased to about 0.2% by weight under very high partial pressures of nitrogen. In plasma furnaces, however, a substantial fraction of nitrogen is present in atomic, excited and ionic state. Under identical conditions of temperature and pressure, plasma processing can increase the dissolved nitrogen content in iron or steel by as much as an order of magnitude over the levels attained in induction furnaces. The tensile strength, yield point and creep strength of 18Cr-8Ni steel are found to increase with the amount of dissolved nitrogen. A nitrogen content of 0.8% raises the yield strength from about 200-700 MPa (Dembovski 1985). Substantial improvement in the long-term creep strength and corrosion resistance in various corrosive media are also observed. Reduction of metal halides and oxides to metal in hydrogen plasma offers many technological advantages. The atomic hydrogen and other active species present in the plasma medium considerably enhance the kinetics of the reduction reaction. The high residual hydrogen contents of metals gained by this technique can be brought down to acceptable levels by subsequent vacuum annealing. Zirconium tetrachloride vapor has been reduced with hydrogen to produce metallic zirconium in a hybrid DC transferred arc plasma reactor (Jones et al. 1997). This plasma reduction technique has been successfully applied to reduction of WO3 and MOO3. Arc plasma devices operating in the range of 150-250 kW have been used to produce tungsten and molybdenum. Similar oxide reduction processes have been extended to produce metals of Ni, Cr and other elements using an argon plasma medium containing additions of hydrogen and hydrocarbons.

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6.3.2.3 Processing o f ceramics High-temperature vapor phase chemistry in the thermal plasma medium offers an attractive route for powder synthesis. The process offers flexibility in the choice of the reactants, which can be solids, liquids or gases. The starting material is introduced in the form of a gas, liquid or solid. The reaction proceeds in several steps: vaporization, vapor phase reaction, nucleation and condensation. Carbides, nitrides, borides and oxides of refractory metals have been successfully produced in RF and DC torch plasma-based reactors. 6.3.2.3.1 Oxide powders. Plasma processing has been widely used to produce ultrafine powders of alumina, magnesia, titania and other oxide powders. A plasma-based technology has been developed for large-scale manufacture of pigment-grade titanium dioxide by the Tioxide Group in UK under the name ICON process. The process essentially makes use of a DC plasma torch to oxidize TIC14. The principal advantages of the plasma process over the conventional process are simplification of the process steps, narrow particle size range and precise control over product quality. Ultrafine silica is used as a filler material in composites and as a thixotropic additive material in the paint industry. The active silica product is prepared by vaporizing SiO2 in an oxidizing atmosphere followed by rapid quenching. The resulting powder is highly active and amorphous. Ultrafine aluminum oxide, iron oxide and zirconium oxide have been synthesized in an inductively coupled plasma (ICP) reactor (Kagawa et al. 1983). Nitrate solutions of the respective metals were introduced into the plasma jet, where decomposition of the fine droplets of the nitrate into the oxide occurred. The fast quench rate favored the formation of metastable ),-phase of A1203 from aluminum nitrate. The dissociation of zirconium nitrate in the ICP reactor resulted in the formation of ZrO2, which was predominantly in the metastable tetragonal phase. 6.3.2.3.2 Non-oxide ceramics. The carbides of silicon, titanium, boron, and tungsten are characterized by high hardness and excellent wear and corrosion resistant properties. These materials find application in the cutting tool industry and also as abrasives. The high energy density of the transferred arc has been used to produce metallic vapors that on condensation produce submicron aerosols. These vapors are also reacted in a suitable environment to produce a variety of products. Submicron size silicon carbide powders have also been produced with SIC14 and methane in a DC plasma torch (Allaire et al. 1991). Ultrafine powders of silicon nitride and aluminum nitride have also been produced using arc plasmas using ammonia as the nitriding gas (Ageorges et al. 1994). Processing of ultrafine powders of refractory carbides and nitrides is one of the widely studied areas of plasma processing. Most of the research efforts have been directed to the study of gas-phase reactions in plasma medium involving metal halides and suitable reactive gas to give ultrafine powders of refractory carbides and nitrides (Yoshida et al. 1979). Titanium carbide and tungsten carbide have been synthesized in DC and RF plasma reactors from titanium tetrachloride in a methane atmosphere. Ti and W metal powders have also been used to synthesize TiC and WC powders (Ananthapadmanabhan et al. 1994b). Plasma processing in RF plasma torches has also been used to prepare titanium

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141

nitride. The process involves the use of titanium halide or titanium metal powder with ammonia or nitrogen as the reactive gas. The liquid injection plasma synthesis (LIPS), where a liquid precursor is injected into the plasma jet, is widely used to process nanosized ceramic powders (Kong and Lau 1990). The liquid injection can be either parallel to the flow or counter current to the flow. Ultrafine oxide and non-oxide ceramic powders and mixed oxide compounds have been synthesized using the LIPS and counter current plasma synthesis (CCPS). Refractory carbides, nitrides, high temperature super-conducting oxide and mixed oxide powders have been produced in RF and DC based plasma reactor systems. 6.3.2.4 Treatment of hazardous wastes Conventional incinerators are, often, unsatisfactory for effective treatment and disposal of hazardous wastes, which are often mixtures of metals, organic and inorganic compounds. Production of toxic gases such as dioxin, insufficient thermal destruction of inorganic constituents, production of polluted flue gases, etc., are some of the drawbacks of conventional waste management systems. Thermal plasma technology has been shown to be potentially useful in the management of industrial waste (Boulos 1996), hospital waste, etc. Multi-torch plasma furnaces have been developed for destruction of highly toxic waste materials such as polychlorinated biphenyls (PCBs), pesticides, herbicides and organic solvents. These systems are also suitable for nuclear waste treatment (Ageorges et al. 1994).

6.3.3 Processing of metastable phases The unique feature of plasma processing is its ability to overcome thermodynamic and kinetic barriers and gain access to non-equilibrium regimes. The high temperature and the high quench rate (105-107 K/s) help in processing metastable and amorphous phases. As the reactants enter the plasma jet, they are heated to their melting temperatures and very often, an appreciable fraction of the particles vaporize and react with each other. The atomic and molecular clusters of the product are carried to the cooler regions of the reactor, where they are quenched. The entire process is completed in a few milliseconds. This time frame is too small for atomic rearrangement to form crystalline structures. It has been observed that thermal plasma processing of silica and other silicate minerals, such as zircon invariably produces highly reactive silica in an amorphous state as confirmed by X-ray diffraction of the product. Plasma process is capable of producing metallic glasses, amorphous silicon, diamond and other metastable phases of ceramic materials.

6.3.3.1 Plasma deposition of diamonds Conventionally, diamond is synthesized by recrystallization of metal-solvated carbon under high temperature (2,000~ and extreme pressures (50-100 kbar). Vapor-phase synthesis of diamond by the hot-filament chemical vapor deposition (CVD) technique has been known for many years. However, the technique is characterized by extremely low deposition rates. The use of a thermal plasma jet for diamond deposition is a recent development. Thermal dissociation of methane and other hydrocarbons in DC and RF plasma jets have been successfully applied to develop diamond coatings (Yoshida et al.

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Figure 6.8 Surfacemorphologyof plasma deposited diamond

1979, Meyer et al. 1989, Metaxas 1996). In contrast to typical growth rates of 0.05/xm/h in hot filament CVD methods, the plasma process is characterized by high growth rates of the order of 900 #m/h (Ohtake et al. 1991). Ramachandran et al. (1996) have injected a mixture of methane and hydrogen gas into an argon plasma jet issuing out of a DC arc plasma torch operated in the non-transferred arc mode. Figure 6.8 shows the surface morphology of diamond deposited on silicon substrate. The cauliflower-like structure is characteristic of high growth rate. The Raman spectrum, indicating the characteristic Raman Shift of the diamond deposit, is shown in Fig. 6.9. The precursor gases employed consisted of a mixture of methane and hydrogen in the ratio 1:99. The experimental conditions strongly influence the diamond growth rate, both qualitatively and quantitatively (Baldwin et al. 1994). In particular, the substrate temperature and hydrogen/methane ratio are significant. Although the exact mechanism of diamond formation is still not understood, high temperature gas phase reactions under plasma environments involving methyl radicals is proposed to be the main operating mechanism.

6.3.3.2 Thermal plasma synthesis of ultrafine alumina powder A classic example illustrating the non-equilibrium nature of reactions occurring in the plasma medium is exemplified by plasma oxidation of aluminum powder (Ananthapadmanbhan et al. 1996). The experimental set-up used is shown in Fig. 6.4. Aluminum metal powder, stored in a powder feeder, was fed into the plasma using argon as carrier gas. Compressed air was injected into the reactor. The powder particles melt and vaporize and the molten droplets and vapors of aluminum react with oxygen, resulting in the

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143

800

4330

cr~ t

600

400 ,,

A

15IXl

Wovenumbers ( cn~ I )

|

1000

Figure 6.9 Raman spectrum of diamond deposit.

Table 6.5 Typicalexperimental parameters Power input (kW) Arc voltage (V) Plasma gas Ar (LPM) Secondary gas N2 (LPM) Carrier gas Ar (LPM) Reactive gas air (LPM) Powder feed rate (g/min)

16 40 20 02 10 40 05

formation of aluminum oxide that is deposited on the reactor walls. Typical operating parameters are given in Table 6.5. X-ray diffraction results indicate that the powder consists of the metastable ),aluminum oxide. Formation of the ),-phase is consistent with the highly non-equilibrium nature of the process. Vapors and molten droplets of aluminum react with oxygen to form aluminum oxide, which crystallizes in the metastable ),-phase by virtue of the rapid quench rate associated with the process. Figure 6.10 is a scanning electron micrograph (SEM) of a typical sample of alumina powder collected from the main section of the reactor. Individual particles are not discernible, but agglomerates consisting of nanosized particles can be seen. A transmission electron micrograph (TEM) of a typical sample is shown in Fig. 6.11. The clusters are resolved into individual particles with a size ranging

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Figure 6.10. SEM photographof plasma synthesized alumina

from a few nanometers to about 30 nm. The spherical morphology of the particles is also evident from the figure.

6.3.3.2.1 Formation of },-alumina. The free energy minimization plot for the system A1O-N has been developed using the CSIRO thermopackage (Turnbull and Wadsley 1988). Figure 6.12 shows the thermodynamic equilibrium diagram for the system aluminumoxygen-nitrogen in the molar ratio 1:2:8 (close to the experimental conditions) under 1 atmosphere pressure. The Y-axis shows the concentration of the species in number of mols and the X-axis shows the temperature. Argon as the plasma gas is also included in the thermodynamic calculations. However, only the oxides of aluminum, atomic species of oxygen and nitrogen and aluminum vapor are shown in the figure. It can be seen from Fig. 6.12 that at temperatures greater than 5,000 K, which are commonly encountered in plasma reactors, the stable species present are aluminum vapour, atomic oxygen and atomic nitrogen. In the temperature region 3,500-4,500 K, A10(g) becomes increasingly stable. The other oxide species of aluminum such as A120, A102 and A1202 are also formed in this temperature zone, although at lower concentration levels, possibly by the following reactions: Al(g) + O(g) = A10(g)

(6.6)

145

Thermal Plasma Processing w

~

" ,'~t~" ",

. . ~.. .

.:,:

*

'

.:

~

.

.

..

'*~" , ~

~--

.

.

100

nm

Figure 6.11. Transmission electron micrograph of plasma synthesized alumina.

2Al(g) + O(g) = A120(g)

(6.7)

2Al(g) + 20(g) = A1202(g)

(6.8)

Al(g) + 20(g) = A102(g)

(6.9)

2A10(g) = A120(g) + O(g).

(6.10)

As the temperature falls below 4,500 K, oxygen molecules and nitrogen molecules (not shown in the figure due to their high concentration) start forming by the recombination of the respective atomic species. At temperatures below 3,500 K and above 2,500 K, liquid aluminum oxide is formed. Solid ot-A1203, which is the only thermodynamically stable form of alumina, results as the temperature drops to below 2,300 K, which is the

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P. V. Ananthapadmanabhan and N. Venkatramani

1 0.9 0.8 0.7 0.6 o.s

~s)

N(g)

0.4 0.3 0.2 0.1 0 0 0 0 '~--

0 0 it) ~--

0 0 0 tN

0 0 rid r

0 0 0 03

0 0 ur) ~

0 0 0 ~I*

0 0 it') ~

0 0 0 IZ)

0 0 it') it')

0 0 0 t,,ID

T ~ ~ Figure 6.12. Free energy minimization plot for the system AI-O-N (1:2:8).

melting point of alumina. Below 2,300 K, the stable species are solid ot-A1203, nitrogen and oxygen molecules (not shown due to their high concentration). Thermodynamic considerations would lead us to conclude that ct-Al203 is formed. However, the conditions existing in the plasma reactor, augmented by thermochemical effects associated with the oxidation of aluminum, favor the formation of non-equilibrium phases. The first step leading to the formation of alumina is the melting and vaporization of aluminum powder injected into the plasma. Molten droplets and aluminum vapor react with oxygen downstream of the reactor to produce aluminum oxide according to the following equation: 2A1 + 3 0 2 -- A1203, AH ~ -- - 1675.7 kJ/mol.

(6.11)

Oxidation of aluminum to alumina is highly exothermic, the standard enthalpy of formation of aluminum oxide from the elements being -1675.7 4- 1.2 kJ/mol (Chase et al. 1985). The enthalpy of formation increases with the reaction temperature and is -2243.73 kJ/mol at 3,000 K. The heat released during oxidation coupled with the high temperature in the plasma medium vaporizes aluminum oxide, leading to the formation of suboxides of aluminum and if the temperature is higher than 4,500 K aluminum vapor and oxygen atoms are formed. As aluminum vapor and oxygen atoms move down to regions where the temperature is between 3,500-4,500 K, they react to form A10 gas and other oxides of aluminum. This is followed by a rapid quench of the vapors as they move to the cooler zones of the reactor tube. The quench rate in plasma processing has been estimated to be

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of the order of 106-108 K/s and the extent of supercooling achievable is much greater than 0.2Tm, favoring homogeneous nucleation resulting in nanosized particles of ),-alumina. In principle, the selection of a metastable phase during solidification is associated with kinetic factors and requires suppression of the nucleation of the stable phase. (See also Chapter 2 on "Thermodynamics and Kinetics of Metastable Phase Formation" by K. N. Ishihara in this book). The nucleation rate of a solid condensing from liquid or gaseous phase is decided by its critical nucleation free energy, A G*, which is the energy barrier for nucleation. The relative nucleation rates of a and y phases of alumina depend on the ratio of their respective critical nucleation free energies, A G * / A G•*. McPherson (1973) has estimated the critical free energies of nucleation for a- and v-alumina as a function of temperature. The estimated ratio of the relative critical free energy for nucleation of a- and y-alumina indicates that v-alumina would nucleate in preference to a-alumina at temperatures lower than 1,740~ The estimated solidification temperature for alumina particles falls in the range of 1,500-1,600~ under plasma conditions, suggesting the formation of v-alumina in preference to a-alumina. Formation of v-alumina in preference to the thermodynamically stable a-phase is typical of rapid quench techniques such as plasma-spraying (Heintze and Uematsu 1992), arc discharge (Hirayama 1987) and electro-hydrodynamic atomization (Jayaram and Levi 1989). Spherical particles of alumina prepared by electro-hydrodynamic atomization have been shown to consist of metastable phases like y- and 6-alumina. Aluminum oxide prepared by vapor condensation from arc, flame oxidation of aluminum salts also consists of metastable phases rather than the thermodynamically stable form. However, the final phase obtained depends on the thermal excursion the particles experience during recalescence and subsequent cooling of the solid powder. This depends on the particle size and the cooling rate. The transformation of y-A1203 to a-A1203 for plasma-sprayed alumina droplets has been analyzed by McPherson (1981), who estimated the achievable supercooling and the thermal history of droplets of different sizes. The analysis shows that particles less than 1 # m in diamater consist entirely of y-A1203, whereas particles >50 # m are completely transformed to the a-phase. Levi et al. (1988) have developed a more rigorous model to examine microstructure evolution continuously through the liquid cooling stage, recalescence and subsequent cooling of the solid powder. Their analysis shows that decreasing the droplet size results in higher supercoolings and more efficient cooling leading to the formation of the metastable phase. The particle size range normally observed in plasma processed alumina is of the order of a few tens of nanometers suggesting that high supercoolings favor the formation of the v-phase. The above analysis helps to illustrate the physicochemical principles involved in the formation of non-equilibrium phases under plasma environment. The generalities of the model can be extended to any other specific case. 6.4. S U M M A R Y AND C O N C L U S I O N S

An overview of the recent developments in the field of thermal plasma processing has been presented. Thermal plasma processes are characterized by high temperature, high enthalpy and power density, 10-100 times higher than those encountered in conventional

148

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a n d N. V e n k a t r a m a n i

furnaces. Different plasma devices ranging from a few hundreds of k W to a few M W p o w e r levels have been designed and developed, over the past decade, to meet a variety of requirements ranging from p o w d e r synthesis to primary and secondary metallurgical operations and treatment of hazardous wastes. The use of plasma-sprayed ceramic and alloy coatings for industrial applications is a well established technique. But, there have been newer developments such as the reactive plasma-spray technique, which combine plasma processes and chemical reactions in one single operation. A further development in this area is the direct conversion method using cheap raw materials. The gas-to-solid conversion in plasma process to fabricate diamond coatings is a variant of this method. Contrary to conventional processing techniques, which are limited by thermodynamic and kinetic barriers, plasma processing, by virtue of the high temperature, high power density and rapid quench, offers unlimited scope for preparation of metastable phases and amorphous structures. The presence of reactive species provides new reaction paths to synthesize a variety of novel materials.

ACKNOWLEDGMENTS

The authors thank Mr. U. K. Chatterjee, Head, Laser & Plasma Technology Division for his constant support and encouragement. The cooperation and help received from Mr. K. P. Sreekumar and Dr. D. S. Patil are also gratefully acknowledged.

REFERENCES

Ageorges, H., Megy, S., Chang, K., Baronnet, J. M., Williams, J. K. and Chapman, C. (1994) Plasma Chem. Plasma Process., 13, 613. Allaire, E, Parent, L. and Dalleuie, S. (1991) J. Mater. Sci., 26, 4160. Ananthapadmanabhan, P. V. and Taylor, P. R. (1998)J. Alloys and Compounds (in press). Ananthapadmanabhan, P. V., Sreekumar, K. P., Muraleedharan, K. V. and Venkatramani, N. (1991) Surf. Coat. Technol., 49, 62. Ananthapadmanabhan, P. V., Sreekumar, K. P., Ravindran, P. V. and Venkatramani, N. (1993) Thin Solid Films, 224, 148. Ananthapadmanabhan, P. V., Sreekumar, K. P., Iyer, K. V. and Venkatramani, N. (1994a) Mater Chem. Phys., 38, 15.

Ananthapadmanabhan, P. V., Sreekumar, K. P. and Venkatramani, N. (1994b) Plasma Chemical Processing of Refractory Carbides, Plasma Science and Technology, ed. Das, I. M. L. (Allied Publ., New Delhi, India), p. 627. Ananthapadmanabhan, P. V., Sreekumar, K. P., Ravindran, P. V. and Venkatramani, N. (1995) in Thermal Spraying--Current Status and Future Trends, Vol. 2, ed. Ohmori, A. (High Temperature Soc. of Japan, Tokyo), p. 1127. Ananthapadmanabhan, P. V., Sreekumar, K. P., Venkatramani, N. and Taylor, P. R. (1996) J. Alloys and Compounds, 244, 70. Ananthapadmanabhan, P. V., Sreekumar, K. P., Venkatramani, N. and Taylor, P. R. (1997) in EPD Congress 1997, ed. Mishra, B. (TMS, Warrandale, PA), p. 209. Baldwin, Jr., S. K., Owano, T. G. and Kruger, C. H. (1994) Plasma Chem. Plasma Process., 14, 383. Bonet, C. (1976) Chem. Eng. Prog., 72, 63. Boulos, M. (1992) J. Thermal Spray Technol., 1, 33. Boulos, M. I. (1996) Pure Appl. Chem., 68, 1007. Bratton, R. J. and Lau, S. K. (1981) inAdvances in Ceramics, Vol. III, Science and Technology ofZirconia, eds. Heuer, A. H. and Hobbs, L. W. (American Ceramic Society, Columbus, OH), p. 226.

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Chase, Jr., M. W., Davies, C. A., Downey, Jr., J. R., Frurip, D. J., McDonald, R. A. and Syverud, A. N. (eds.), (1985) JANAF Thermochemical Tables, third edition (The American Chemical Society and The American Institute of Physics for the NIST, New York). Dembovski, V. (1985) Plasma Metallurgy: The Principles, Materials Science Monographs, No. 23, (Elsevier, Amsterdam). Eliezer, Y. and Eliezer, S. (1989) The Fourth State of Matter: An Introduction to the Physics of Plasma (Adam Hilger, Bristol, UK). Fauchais, E and Vardelle, M. (1994) Pure Appl. Chem., 66, 1247. Fauchais, P., Bourdin, E., Coudert, J. E and McPherson, R. (1983) in Topics in Current Chemistry, Vol. 107, eds. Veprek, S. and Venugopalan, M. (Springer-Verlag, Berlin/Heidelberg), p. 59. Fey, M. G., Wolf, C. B. and Harvey, J. E (1975) in Proceedings of the International Round Table: Transport Phenomena in Thermal Plasmas (Odeillo, France). Fineman, J. (1987) Plasma Technology in Metallurgical Industry (Iron and Steel Society, Warrendale, PA). Gerdeman, D. A. and Hecht, N. L. (1972) Arc Plasma Technology in Materials Science (Springer-Verlag, Wien). Heberlein, J. V. R. (1992) Pure Appl. Chem., 64, 629. Heintze, G. N. and Uematsu, S. (1992) Surf. Coat. Technol., 50, 213. Hench, L. L. (1991) J. Amer. Ceram. Soc., 74, 1487. Herman, H. and Bhat, H. (1980) in Synthesis of Metastable Phases, eds. Machlin, E. S. and Rowland, T. J. (TMS- AIME, Warrendale, PA), p. 115. Hirayama, T. (1987) J. Amer. Ceram. Soc., 70, C-122. Jayaram, V. and Levi, C. G. (1989) Acta. Metall., 37, 569. Jones, H. (1973) Rep. Prog. Phys., 36, 1425. Jones, R. T., Barcza, N. A. and Curr, T. R. (1997) Plasma Developments in Africa (Internal Report, Mintek, Randburg, South Africa). Kagawa, M., Honda, E, Onodera, H. and Nagae, T. (1983) Mater Res. Bull., 18, 1081. Knight, R., Smith, R. W. and Lawley, A. (1995) Inter J. Powder Metall., 31, 205. Kong, E C. and Lau, Y. C. (1990) Pure Appl. Chem., 62, 1809. Levi, C. G., Jayaram, V., Valencia, J. J. and Mehrabian, R. (1988) J. Mater Res., 3, 969. Lin, C. K. and Berndt, C. C. (1994) J. Thermal Spray Technol., 3, 75. Mackinnon, I. M. and Reuben, B. G. (1975) J. Electrochem. Soc., 122, 806. Madejski, J. (1976) Inter J. Heat and Mass Trans., 19, 1009. Matejka, D. and Benko, B. (1989) Plasma Spraying of Metallic and Ceramic Materials (John Wiley, New York). McPherson, R. (1973) J. Mater Sci., 8, 851. McPherson, R. (1981) J. Mater Sci., 16, 3141. McRae, L. B., Barcza, N. A. and Curr, T. R. (1985) in Proceedings of the International Conference on Mineral Science and Technology, Vol. 2, ed. Haughton (Council for Mineral Technology, Randburg, South Africa), p. 827. Meetham, G. W. (1988) Mater Design, 9, 308. Metaxas, A. C. (1996) Foundations ofElectroheat: A Unified Approach (John Wiley, Chichester, UK). Meyer, D. E., Dillon, R. O. and Woollam, J. A. (1989) J. Vac. Sci. TechnoI., AT, 2325. Miller, R. A., Smialek, J. L. and Garlick, R. G. (1981) inAdvances in Ceramics, Vol. III, Science and Technology ofZirconia, eds. Huer, A. H. and Hobbs, L. W. (American Ceramic Society, Columbus, OH), p. 241. Mohanty, M. and Smith, R. W. (1995) J. Therm. Spray Technol., 4, 384. Mutasim, Z. Z., Smith, R. W. and Mohanty, M. (1992) in Thermal Spray--International Advances in Coating Technology, ed. Berndt, C. C. (ASM International, Materials Park, OH), p. 1019. Notomi, A. and Hisatome, N. (1996) Pure Appl. Chem., 68, 1101. Ohtake, N., Mashimo, Y. and Yoshikawa, M. (1991) New Diamond Sci. Technol., 173, 178. Patil, D. S., Sreekumar, K. E, Venkatramani, N., Iyer, R. K., Ram Prasad, Koppikar, R. S. and Munim, K. R. (1996) Bull. Mater Sci. (India), 19, 115. Pauloski, L. (1995) The Science and Engineering of Thermal Spray Coatings (John Wiley, New York). Pfender, E. (1994) in Proceedings of the International Symposium on Heat and Mass Transfer, ed. Fauchais, E (Bell House, Turkey), p. 223. Ramachandran, K., Patil, D. S., Venkatramani, N., Biswas, A. R., Venkateswaran, S. and D' cunha, R. (1996) Indian Vac. Soc. Bull., 27, 23. Ramachandran., K., Selvarajan, V., Ananthapadmanabhan, P. V. and Sreekumar, K. P. (1997) Plasma Devices and Operations, 5, 1.

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Rangaswamy, S. and Herman, H. (1986) Advances in Thermal Spraying. (Pergamon Press, Oxford, UK), p. 101. Ranz, W. E. and Marshall, W. R. (1952) Chem. Eng. Prog., 48, 141. Rykalin, N. N. (1976) Pure Appl. Chem., 48, 179. Rykalin, N. N. and Kudinov, V. V. (1976) Pure Appl. Chem., 48, 229. Scott, K. T. and Kingswell, R. (1991) in Advanced Surface Coatings, eds. Rickerby, D. S. and Matthews, A. (Blackie, London). Stecura, S. (1977) Amer. Ceram. Soc. Bull., 56, 1082. Taylor, E R., Manrique, M., Pirzada, S. A. and Abdel-Latif, M. (1995) Plasma Chem. Plasma Proc., 15, 545. Turnbull, A. G. and Wadsley, M. W. (1988) The CSIRO Thermochemistry System, version 5 (IMEC, Australia). Venkatramani, N. (1995) Bull. Mater. Sci. (India), 18, 741. Yoshida, T., Kawasaki, A., Nakagawa, K. and Akashi, K. (1979) J. Mater. Sci. 14, 1624. Zaat, J. H. (1983) Ann. Rev. Mater. Sci., 13, 9. Zoltowski, P. (1968) Rev. Int. Htes. Temp. et. Refract., 6, 65.

Chapter 7 Spray-Forming 7.1. 7.2. 7.3.

Introduction Principles Variations and Distinctions 7.3.1 Variations 7.3.1.1 Crucible: Induction Skull Melting/Spray-Forming 7.3.1.2 Atomizer: Circular vs. Linear Spray-Forming 7.3.1.3 Atomizer: Close-Coupled vs. Free-Fall Spray-Forming 7.3.1.4 Atomization Gas: Reactive Spray-Forming 7.3.1.5 Substrate: Near Net Shape Spray-Forming 7.3.1.6 Spray-Forming and Co-Injection 7.3.2 Nomenclature 7.3.2.1 Spray Atomization 7.3.2.2 Related Processing Techniques 7.4. Applicability 7.5. Non-Equilibrium Phenomena in Spray-Forming 7.5.1 Non-Equilibrium Nature 7.5.1.1 Rapid Solidification in Atomized Droplets 7.5.1.2 Transient Semi-Solid Layer in Deposition Stage 7.5.2 Non-Equilibrium Related Features in the Deposit 7.5.2.1 Metastable Phases 7.5.2.2 Extended Solid Solubility 7.5.2.3 Absence of Macrosegregation/Minimized Microsegregation 7.5.2.4 Refinement 7.5.3 Effects of Non-Equilibrium Features on Mechanical/Physical Properties 7.5.3.1 Direct Effects 7.5.3.2 Indirect Effects 7.6. Concluding Remarks Acknowledgments References

153 154 157 157 158 158 159 160 163 163 166 166 167 168 172 173 173 176 178 178 178 179 180 181 183 185 189 189 189

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Chapter 7 Spray-Forming B ING LI and ENRIQUE J. LAVERNIA

7.1. I N T R O D U C T I O N

Materials science and engineering is a very broad, interdisciplinary field which encompasses several sciences, primarily chemistry, physics, and metallurgy. The materials under investigation include metals, semiconductors, ceramics, polymers, glasses, and composites. The areas of study range from the nature of atomic arrangement to the phenomenology of the macroscopic material properties; and from elegant fundamental theories to exciting practical applications. Among the manifold aspects, an everlasting important philosophy in materials science and engineering is to produce a material having the required physical and mechanical properties at the lowest cost. To fulfill this goal, it is essential to explore advanced materials with superior properties while simultaneously maintaining a competitive cost. This is made possible by development/utilization of novel processing techniques. Each of the above mentioned categories of materials has its own unique processing requirements, either conventional or novel. For metallic materials, there are mainly two most frequently utilized conventional processing techniques, i.e., casting (ingot casting, die casting, investment casting, etc.) and powder metallurgy. Meanwhile, many novel processing techniques are being currently explored/investigated, such as mechanical alloying (Koch 1989), laser surface modification (Steen 1985), ion implantation (Picraux 1984), self-sustained high-temperature processing (Mabuchi et al. 1995, Alman et al. 1995, Rawers and Wrzesinski 1992), and spray-forming, to name a few. Each of these novel processing techniques has its own advantages and disadvantages. Among these many techniques, spray-forming is especially attractive in the sense that it is potentially capable of fabricating advanced materials with improved properties at costs that are industrially competitive. The improved properties are directly associated with the non-equilibrium nature of spray-forming, which is the focus of this chapter and will be discussed in detail in the following sections. The cost effectiveness is derived from the inherent simplicity of this technique, which, in turn, is partially attributable to the non-equilibrium phenomena. It is instrumental to illustrate its inherent simplicity by comparing spray-forming with other processing techniques, especially the two well-established processing techniques: powder metallurgy and ingot casting. Compared with powder metallurgy, spray-forming combines powder production, sieving, canning, cold pressing, degassing and hot press153

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Bing Li and E. J. Lavernia

ing into a single step. This is inherent in the concept of spray-forming which incorporates consolidation into the process of powder production. Compared with the conventional ingot casting, spray-forming eliminates macrosegregation and minimizes microsegregation, hence avoiding many intermediate forming steps, such as long term homogenization treatment and extensive thermomechanical processing, which are typically required in conventional ingot casting techniques to eliminate/minimize segregation. The concept of spray-forming was first introduced by Singer in the early 1970s (Singer 1970, 1972). His initial intent was to develop a processing technique which was capable of reducing the overall cost in manufacturing structural materials. It was soon realized that spray-forming is also unique in obtaining superior microstructures and hence improved mechanical properties. Because of the many potential advantages of this technology over conventional processing techniques, spray-forming aroused more and more interest in the academic field, as well as in the industrial community worldwide over the last three decades. There are now more than 20 research organizations and 20 industrial companies located in the United States, Europe, and Eastern Asia (China, Japan, Korea, etc.) actively involved in the development/utilization of this technology (Leatham 1993). These developmental activities are coupled with a large volume of publications in the field of spray-forming, including journal articles, conference papers, book sections (Yule and Dunkley 1994), and books (Lavernia and Wu 1996). These efforts may be categorized as: (a) understanding of the fundamental phenomena associated with this technology, such as heat transfer, undercooling, and solidification; (b) application of this technology to new alloy systems rather than the widely studied aluminum alloys, such as Cu, Fe, Ni, Ti, Zr, and Mo based alloys; (c) optimization of the processing parameters to achieve better microstructural characteristics and mechanical properties; (d) modification of this technology to extend its application domain from monolithic alloys to metal matrix composites; (e) novel design of this technology to enhance its capability in the control of product geometry (near net shape, such as flat sheet, circular tube, etc.); and (f) commercialization of this technology.

7.2. P R I N C I P L E S

The principle of spray-forming is schematically shown in Fig. 7.1. It consists of two stages: spray atomization and spray deposition. Accordingly, spray-forming is also termed as spray atomization and deposition in the academic field as well as the industrial community. The spray atomization stage further contains the following detailed physical changes: (a) A stream of molten metallic material is impinged upon by highly energetic gas jets in different, normally symmetrically opposite directions, as shown in Fig. 7.1. The result is disintegration of the molten metallic stream into small, irregular ligaments. (b) The small, irregular ligaments transform into spherical droplets shortly after their formation (Lavernia et al. 1992). This change is normally termed spheroidization in

Spray-Forming

155

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Figure 7.1. Two stages of spray-forming: sprayatomization and spray deposition.

spray-forming. The driving force for this change comes from the decrease of the surface energy associated with the ligament (minimum energy principle). The surface energy is the product of surface tension and surface area. The surface tension is a constant for the same material, while the surface area depends on the shape. For a given volume, a sphere has the minimum surface area. Accordingly, there is always a tendency for any shape to change into spherical to minimize the total surface energy. In other words, formation of spherical droplets from irregular ligaments is thermodynamically favorable. In a practical situation, whether spheroidization really occurs depends on the associated kinetics. It has been demonstrated, both experimentally and theoretically, that spheroidization can be completed almost immediately after the formation of the irregular ligaments during spray-forming, mainly because of the high surface tension of most molten metallic materials and the size of the ligaments. (c) The sizes of the spherical droplets are normally on the order of microns, but typically fall into a wide range. Ample experimental evidence suggest that, in any spray-forming experiment, the size distribution of these droplets obeys a specific distribution called log-normal distribution (Lubanska 1970, Liang et al. 1992a, Marinkovich et al. 1989, Irani and Callis 1963, Allen 1981, Grant et al. 1993a, Delplanque et al. 1997). Most of the droplets have sizes close to a certain finite value (median size). When the size deviates further from this value, the probability of finding a droplet having this size decreases. (d) The atomization gas flow field forms a specific geometrical shape, normally a spray cone. The micron-sized droplets are dispersed in the spray cone. (e) The spherical droplets experience momentum transfer with the atomization gas flow field. This momentum transfer governs the movement of the droplets in the spray

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Bing Li and E. J. Lavernia

cone, including trajectory, flight time, and subsequent deformation behavior during the deposition stage. (f) The spherical droplets experience thermal exchange with the atomization gas flow field. This thermal exchange causes the spherical droplets to cool down and solidify during their flight towards the substrate. The thermal exchange is affected, to a certain extent, by the momentum transfer between the droplet and the gas field. The spray deposition stage involves collection of the droplets in the spray cone onto a water-cooled substrate, forming a dense bulk material that is normally referred to as a deposit: (a) By the time when these droplets reach the substrate, they are in the solid, liquid or semi-solid states. (b) At the beginning of the deposition, the droplets impinge directly on the top surface of the substrate. Once these droplets are collected on the substrate, they act as the substrate themselves for the subsequent impinging droplets. (c) Impingement of the droplets with the substrate leads to extensive modification of the droplets, including their morphology and the associated microstructural features, such as dendrites (see Fig. 7.2), which are developed during solidification in flight. This deformation plays a critical role in the evolution of the microstructure of the deposit. (d) The deposited bulk material experiences thermal exchange with the water-cooled substrate and the surrounding environment, leading to the dissipation of its thermal energy. Simultaneously, during spray deposition, the arrival of droplets from the spray cone onto the substrate continuously incorporates thermal energy into the deposited bulk material. These two effects compete with each other, determining the overall thermal behavior of the deposit. This further affects some aspects of the microstructure, such as grain size. The microstructure in turn determines the mechanical properties. Spray-forming is normally carried out in an environmental chamber. The chamber is evacuated and back-filled with a protective gas such as N2, He, and Ar. The purpose of this arrangement is to enhance the capability to either suppress or exploit the possible reactions, such as oxidation, between the atomization/environmental gas and the metallic materials (the droplets and the deposit). The products of spray-forming include the bulk deposit and powders that are usually termed as oversprayed powders. During the deposition stage, some powders (solidified droplets) cannot be collected into the substrate to form the bulk deposit. Instead, they are carried by the gas flow field into a collecting apparatus (cyclone). These powders are termed as oversprayed powders in spray-forming technology. In commercial applications, the bulk deposit is the main product of practical value, while the oversprayed powders must be minimized in quantity. In scientific research, both the bulk deposit and the oversprayed powders are useful in providing valuable information on the understanding of the fundamentals in spray-forming.

Spray-Forming

Figure 7.2. SEM

157

micrograph showingdendrites formed in spray atomized A1-Mn-LiZr powders (Lavernia et al. 1992).

7.3. VARIATIONS AND D I S T I N C T I O N S

Advances in our understanding of the science and technology of spray-forming have led to the development of novel designs, as well as modifications to the original technology, although the fundamental principle remains the same. A comprehensive discussion should encompass a description of these advanced designs. Since spray-forming is currently a very active field, an exhaustive description of all these variations is difficult. Accordingly, in the following sections, several representative variations will be described. It is also worth noting that there is some confusion in the literature regarding the term sprayforming, as applied to related manufacturing technologies. This issue is addressed in the sections that follow. 7.3.1 Variations

Most of the variations to spray-forming involve modification of one or more of the key components of a spray-forming facility, which include the following: crucible, atomizer, atomization gas, and substrate, as will be described in sequence below.

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Bing Li and E. J. Lavernia

7.3.1.1 Crucible: induction skull melting~spray-forming In spray-forming, there is generally a crucible used to melt the master alloy and then to hold the molten metal. During spray-forming of non-reactive alloys with low to intermediate melting temperatures, such as aluminum and magnesium based alloys, ferritic steels, and nickel based intermetallics, the crucible is generally made of graphite or ceramic materials such as zirconia and alumina. For reactive/refractory alloys, these conventional crucibles are no longer usable. In spray-forming, the molten metal is generally heated to a temperature about 50 ~ 100 K higher than its melting temperature and then maintained there for a significant time duration (> 10 min). Such a processing procedure is essential to ensure smooth flow of the molten metal through a delivery tube into the atomizer or, in the worst scenario, to avoid freeze-up of the molten metal at the delivery tube. For reactive alloys, such as Ti (Morris et al. 1992, Li et al. 1996a, Li and Lavernia 1997) and Zr (Li et al. 1997) based alloys, their highly reactive nature in the molten state will lead to significant reaction between the molten alloys and the crucible material during such a long period, making selection/utilization of conventional crucibles extremely difficult if not completely impossible. For refractory alloys, such as Zr (Li et al. 1997), Mo, Nb, V, Ta (Delplanque et al. 1997), and W based alloys, their melting temperatures (>1800~ are normally close to or above those of the commonly used crucible materials (zirconia or graphite). Accordingly, the crucible will be destroyed before the alloys can be melted. To circumvent these problems, the facility has to be significantly modified, and an alternate melting approach, namely skull melting, has to be incorporated. The novelty in the concept of skull melting/spray-forming may be best illustrated by a comparison between the usage of the conventional crucible and the skull crucible (refer to Fig. 7.3). In spray-forming experiments that employ a conventional crucible, the crucible is charged with a master alloy and then heated up using an induction furnace. In the process of heating, the crucible is thermally insulated from its environment, preventing unnecessary loss of thermal energy and enhancing melting efficiency. Consequently, both the crucible and the charge remain at an elevated temperature as they are heated up. In skull melting/spray-forming, the crucible is water-cooled during the entire process of heating of the charged material. When the center of the charge is melted, its exterior regions remain solid because of their intimate contact with the cold crucible. These solid regions of the charged material form a crucible of the shape akin to the water-cooled crucible. This crucible is called a "skull crucible", which leads to the name of this technique. The most intriguing feature of skull melting technique is that the molten metal will be contained in a crucible (the skull crucible) made of the metal itself; hence contamination resulting from the reaction between the molten metal and the crucible is minimized, if not totally eliminated. 7.3.1.2 Atomizer: circular vs. linear spray-forming

In terms of the geometrical shape of the atomizer, there are two typical types of sprayforming techniques: circular and linear. In a circular atomizer, the gas jet nozzles are arranged concentrically (Fig. 7.4a). The number, cross-section area, and shape of the nozzles are critical parameters that control the efficiency of the atomizer. During sprayforming, the atomization gas streams exit these nozzles, forming a circular cone (Fig. 7.5b). The deposit produced this way normally exhibits a bell shape (Fig. 7.5b). In a

Spray-Forming

159

Thermocoupl, Moltenmetal

Stopper Moltenmetal Ceramic cruci Thermalinsul Inductioncoil

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(b)

Figure 7.3. Comparison on the usage of: (a) conventional ceramic crucible and (b) skull crucible.

linear atomizer, the gas jets are located along two parallel lines (Fig. 7.4b). During sprayforming the atomization gas forms a prism-like geometry, giving rise to a deposit with a high aspect ratio (Fig. 7.5a). Hence, a linear nozzle, in combination with a moving substrate, may be readily used to attain a sheet-like geometry (Fig. 7.5a).

7.3.1.3 Atomizer: close-coupled vs. free-fall spray-forming Schematic diagrams illustrating these two types of spray-forming are shown in Fig. 7.6. In close-coupled spray-forming, the molten metal is released from the crucible through a delivery tube. When it exits the tube, the high-speed atomization gas jets impinge on the melt stream at the tip of the tube (Fig. 7.6a). In free-fall spray-forming, the molten metal is released from the bottom of the crucible, forming a stream. The molten metal stream travels downwards, through the atomizer unit, until it is atomized at a certain point below the atomizer. From the point it exits the bottom of the crucible to the point it is atomized, the molten stream is not constrained by any tubes as in close-coupled atomizer (Fig. 7.6b). Instead, it falls freely, which leads to its name: free-fall sprayforming. Close-coupled spray-forming and free-fall spray-forming are complementary to each other. Close-coupled spray-forming is more susceptible to freeze-up problems. Since the delivery tube is in close contact with the cold atomization gas, there is a higher probability for the melt stream to freeze at the tip of the tube. When this happens, further delivery of the molten metal through the tube is impossible, and the entire experiment fails and has to be restarted. Free-fall spray-forming can avert freeze-up problem because of the absence of the delivery tube. A concomitant problem is the instability of the molten stream in the absence of the delivery tube. The molten metal stream may be easily deflected by the turbulence in the surrounding gas envelope. This instability may lead to oscillation and asymmetry of the atomization cone (Lavernia and Wu 1996).

160

Bing Li and E. J. Lavernia

Gas jet nozzles Orl ac(

delivery tube

(a)

Space to accommodate delivery tube

Linear atomizer twins

~as jet slits

(b) Figure 7.4. Arrangementof atomization gas jet nozzles (slits) in circular and linear atomizers: (a) circular, and (b) linear.

7.3.1.4 Atomization gas: reactive spray-forming There are two types of atomization gases: protective and reactive. Helium and argon are common protective gases, while 02, CO, CO2 and H2 are normally categorized as reactive gases. Nitrogen (N2) is either protective or reactive, depending on the nature of the material being synthesized. In many situations, the atomization gas employed is protective. In a recently developed variation, i.e., reactive spray-forming (Zeng et al. 1992, Zeng et al. 1995), the atomization gas is deliberately selected to contain certain amount of reactive gases. Reactive spray-forming is mainly utilized to synthesize a special category of high strength materials: dispersion-strengthened alloys. The most representative example in this category is the oxide dispersion-strengthened (ODS) material. Its earliest application may be dated back to 1910 when Coolidge incorporated thorium oxide into tungsten wire to improve its elevated temperature mechanical properties (Coolidge 1910). The first

Spray-Forming Molten metal

161

...,.<,-~..,-~" " .--<-,.%':"~~ ,.--.2---%'-2---

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~9 ~ ~ ' ; ' " " . . . . . . . . . . . . . . . . . . . .

I

Linear aton~izer ~ ~ - - - ~ . _

\

Substrate motion

(a) Molten metal Delivery tube Circular atomizer

Atomized droplets . . . . .

Deposit Water-cooled copper substrate -Rotation and/or translation movement

......

~,._~

Figure 7.5. Linear vs. circular spray-forming: (a) linear (Lavernia and Wu 1996), and (b) circular.

162

Bing Li and E. J. Lavernia

('ruciMc

Liquid metal

Crucit"

~netal

v tube Ammiz assemM

Pro-heated atomizing assembly

~S

Inert gas /

Spray of atomized /~-.-droplets //

Refractory t ubc

Spray dropl,

\

Figure 7.6. Close-coupled (a) and free fall (b) spray-forming (Singer, 1970, Lavernia and Wu 1996).

structural ODS, sintered aluminum powder (SAP), was developed by Irmann in 1952 (Irmann 1952). In 1970, Benjamin introduced the concept of fabricating ODS using mechanical alloying (Benjamin 1970). Since then, many ODS systems, notably Ni- and Fe-based alloys, have been successfully developed. It should be noted that in either of these conventional techniques, typical powder metallurgy processing procedures, such as powder production, canning, degassing, and consolidation, have to be carried out. In reactive spray-forming, the atomization gas reacts with the droplets, forming inorganic phases such as oxides, carbides, or nitrides. The large surface area associated with the micron-sized droplets greatly facilitates the reaction kinetics. Nevertheless, this is offset to some extent by the limited time of reaction during spray-forming, since the cooling and solidification rates of the droplet are very high. For example, a 20 #m A1 droplet may be cooled down from 800~ to 80~ in about 6 ms. Consequently, only a very thin intermetallic film (0.1-30 nm) may be formed on the surface of the droplet. This ensures that the size of the intermetallics is fine (<0.1 #m), and hence presumably beneficial to the mechanical properties. The presence of coarse intermetallic particles generally has detrimental effect on the mechanical properties. During deposition stage, the thin film will be deformed and fragmented due to the high velocity impact of the droplets with the substrate/deposit, leading to a uniform distribution of the intermetallic particles in the deposit. A uniform distribution of the intermetallic particles in the matrix is another criteflon that must be satisfied for them to be beneficial to mechanical properties. The main advantage of reactive spray-forming over powder metallurgy for the fabrication of dispersion strengthened materials lie in its simplicity: it combines powder production, reaction, canning, degassing, and consolidation into a single step.

Spray-Forming

163

Induction

heated ladle

Tube

.--L

~x- ~ -

~.....

AKomizer

_~_

Recipient Deposition chamber

- - - - Exhaust

Figure 7.7. Spray-forming of tubular materials (Lavernia and Wu 1996, Fiedler et al.

1987).

It is worth noting that even in reactive spray-forming, the protective environment in the chamber mentioned in Section 7.2 remains necessary. In the absence of a protective environmental chamber, the ratio of the reactive to the protective gas will change unpredictably. This unpredictability is to be avoided either in scientific research or scale-up commercial production.

7.3.1.5 Substrate: near net shape spray-forming As stated earlier, segregation (macro- and micro-) is inherently minimized in sprayforming, ensuring microstructural homogeneity in the deposit. Such a homogeneity leads to another important feature of spray-forming: its potential capability for net shape or near net shape forming, which means that the products of spray-forming need no, or limited, further thermomechanical treatment. The importance of this feature lies in the associated cost reduction when applied to a commercial product. However, full utilization of this potential depends on proper design of the spray-forming facility, especially the geometrical shape and the movement of the substrate. Representative examples are shown in Fig. 7.5a and Figs. 7.7-7.10. From these designs, flat sheets (Fig. 7.5a), round tubes (Fig. 7.7), strip ribbons (Fig. 7.8), rolls (Fig. 7.9), and round billets (Fig. 7.10) may be manufactured in near net shape. 7.3.1.6 Spray-forming and co-injection This technique was developed during the course of synthesizing metal matrix composites (MMCs) using spray-forming techniques. Metal-matrix composites may be synthesized exactly the same way as that employed to synthesize the monolithic alloys using the conventional spray-forming technique. In this case, the charge in the crucible is either a pre-synthesized metal matrix composite bulk material or a mixture of the monolithic alloy and the reinforcement. Challenges to this method include:

Bing Li and E. J. Lavernia

164

Liquid metal

Crucible

Close-type atomizer

Inert gas Chamber

Atomized droplet spray

Spray-deposited strip Roll

Roll, substrate Hydraulic ram Hot rolled strip

Figure 7.8. Spray-forming of rolled strips (Singer 1970, Singer 1972, Lavernia and Wu 1996).

Chamber

9

Collector

F~refo~~ _~'-----~-

I//.O'/Z~7/////I'~,,,'$.

i (a)

'

"//.////h (b)

Figure 7.9. Spray-forming of mill rolls: (a) schematic diagram of the equipment and (b) rolls produced (Lavernia and Wu 1996, Ikawa et al. 1990).

165

Spray-Forming

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induction----.--~

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(c) Figure 7.10. Spray-forming of round billets using (a) horizontal substrate plus vertical motion (Lavernia and Wu 1996), (b) inclined substrate plus vertical motion (Lavernia and Wu 1996, Leatham and Lawley 1993), and (c) inclined substrate plus horizontal motion (Leatham 1996).

i

166

Bing Li and E. J. Lavernia

(a) Long time exposure of the reinforcement to the matrix environment at elevated temperatures, leading to detrimental interfacial reactions between the reinforcement and the matrix. (b) The difference in melting temperatures of the matrix and the reinforcement. In many cases, such as A1/SiC, the difference is so large that it is not practical to melt the matrix and the reinforcement simultaneously. Accordingly, the matrix is melted while the reinforcement remains solid. In the extreme, when the volume fraction of the reinforcement is high, the fluidity of the melt becomes an issue. In this case, it will be very difficult for the melt containing the reinforcement to flow through the delivery tube. Therefore, it is difficult to synthesize MMCs with high volume fraction of reinforcement using this method. As an alternative to circumvent these problems, spray-forming and co-injection was developed in the 1980s (Singer and Ozbek 1985, Willis 1988, Gupta et al. 1989, 1990, Wu and Lavernia 1991). Spray-forming and co-injection does not involve modification of any of the key components of spray-forming mentioned earlier, i.e., crucible, atomizer, atomization gas, and substrate. Instead, it incorporates a new component: the particulate injection unit, as illustrated in a schematic representation of the spray-forming and co-injection facility used to fabricate TiA1/TiB2 composites (Fig. 7.11). The particulate supplying compartment is connected to a high-pressure protective gas on one side, and injector tubes the other side. The injector tubes are positioned at a certain distance below the atomizer. In order to obtain uniform distribution of the reinforcement in the composite, these injector tubes are normally positioned symmetrically around the central line of the spray cone. During the experiment, the crucible is charged with the matrix alloy. Upon melting, the molten matrix alloy is released through the delivery tube and subsequently atomized, forming the spray cone. Following atomization, as the partially solidified matrix droplets travel towards the water-cooled substrate under the combined forces of gravity and fluid drag, the injector tubes release high speed gas flow which contains entrained reinforcement particulates. The reinforcement particles penetrate into the matrix droplets, forming composite droplets. The micron-sized composite droplets experience cooling and solidification, and are finally collected on the substrate as a highly dense composite preform.

7.3.2 Nomenclature 7.3.2.1 Spray atomization The term spray-forming is sometimes confused with spray-atomization in the scientific literature. Possible reasons include: (a) The term spray-atomization has a much longer history than spray-forming. Its earliest utilization may be traced back to the 1920s (Lawley 1992). (b) Spray-forming was originally developed on the basis of spray atomization by incorporating a substrate across the spray cone to collect the impinging solid/semisolid/liquid droplets. Hence, in the absence of deposition stage, spray-forming becomes spray atomization.

167

Spray-Forming En vi ron mental ch a

m

b

_

~

Skull crucible ..............{~z].....................i.....................~.. Moltenmelt . ~ ~ I ~ Induction coil (2) ~ ~\ - ) ' ' ~ - ~ ~ ~t

Nozzle ~ Atomizer .... J

".

~

~

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~1 ~

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~

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~

if/

~

IJ1

I~ ...... ~ l~..,.......,~

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"~/ . . . . . . u] l~\ectors ] T i B 2 _ ~

Deposited TiAlFfiB composite Water cooled --deposition substrate

............. I

~ j , / j ~ ~,"\~ .....~ i

HIPOWDER INJECTOR COLUMN -3"-~t ' / ~] llnert gas cxhau,t ;

.- CycJo.e ! collector CYCLONE SEPARATOR Figure 7.11. Schematic diagram showing the spray-forming and co-injection facility used to fabricate TiA1/TiB2 composites (Li and Lavernia 1997).

(c) In many cases, the exterior appearance of spray-forming facility is the same as that of a spray-atomization facility. However, from a processing point of view, spray-forming is completely different from spray-atomization. Spray-atomization is a powder production technology. It is only a method to synthesize powders which will be used in powder metallurgy technology. It is only one of the many steps in powder metallurgy processing technology. In comparison, spray-forming is a technology that parallels conventional casting and powder metallurgy, with many potential advantages over conventional casting and powder metallurgy. 7.3.2.2 Related processing techniques

Other spray-related processing techniques include: plasma spray-forming (Heimann 1996, Sampath et al. 1991, Zhao et al. 1993a), the high-velocity oxyfuel thermal spray process (HVOF) (Parker and Kutner 1991), and the electric arc spray process (Buhrmaster et al. 1988, Fussell et al. 1994, Zhao et al. 1993b). The primary differences between these techniques and spray-forming are summarized in Table 7.1. The most critical difference is in the dimension of the product of these processing techniques. One of the main reasons that spray-forming has attracted increasing attention is that it is capable of fabricating bulk materials, just like the conventional casting and powder metallurgy processing techniques.

168

Bing Li and E. J. Lavernia

Table 7.1 Differences between various spray related processing techniques

Techniques

Plasma

High velocity oxyfuel

Electric arc

spray-forming

thermal spray process

spray process

Melting method

Plasma arc

Internal combustion jet

Electric arc

Induction furnace*

Feedstock

Powders

Powders

Consumable

Bulk material

Coating

electric wires Coating

Tonnage bulk material

Product

Coating t

Spray-forming

* Sometimes electrical resistance heating. t Sometimes three dimensional, but limited in size.

7.4. A P P L I C A B I L I T Y

In principle, spray-forming is applicable to any material system, especially metallic systems. The only limitation lies in the melting of the material when: (a) the material's melting temperature is extremely high, such as Ta and W; (b) the material reacts with the crucible when it is at an elevated temperature or in the molten state, such as Ti and Zr; (c) the material is not electrically conductive, or the material is not in bulk form (such as powders, wires). In these cases, induction heating may not work properly. These problems may be circumvented by modifying the spray-forming facility, especially the melting unit, such as has been performed in the implementation of the induction skull melting/spray-forming technique (Delplanque et al. 1997, Morris et al. 1992, Li et al. 1996a, Li and Lavernia 1997, Li et al. 1997). Spray-forming has been heretofore employed to manufacture a variety of materials, including A1, Cu, Fe, Mg, Ni, Ti, and Zr based monolithic alloys, Ni-A1, Ti-A1, and Fe-A1 intermetallics, and metal/intermetallic matrix composites. These materials are summarized in Table 7.2 for conventional monolithic alloys and intermetallics, Table 7.3 for novel alloys, and Table 7.4 for composites. The motivations for manufacturing these materials using spray-forming include one or several of the following aspects: (a) refined and uniform microstructure generated by spray-forming which may lead to improved mechanical properties; (b) cost effectiveness of spray-forming; (c) eliminated macrosegregation, minimized microsegregation, and extended solid solubility, which may lead to development of novel alloy compositions, such as in A1-Li, A1-Si, A1-Cu, A1-Zn and Cu-Zr alloy systems; (d) rapid solidification effects in spray-forming which promote improved physical properties, such as electric conductivity of Cu based alloys, magnetic properties in Fe-Nd-B systems, as well as damping capacities in A1/graphite/SiCp and A1/graphite composites; (e) superfine and uniformly distributed oxides in reactive spray-forming, which may be exploited as dispersion strengthening phases;

Spray-Forming Table 7.2

169

Spray-formed conventional commercial/experimental alloys*

Category

Specific alloys

References

A1 based alloys

2014A1 2024A1 2090A1 2618A1 A390A1 5083A1 6061A1 7075A1

Annavarapu and Doherty (1995), White et al. (1988) Lavernia et al. (1986) Lavernia and Grant (1988) White et al. (1988), Underhill et al. (1993) Wu et al. (1995) Wu et al. (1997) Wu and Lavernia (1991), White et al. (1988) Grant and Cantor (1990), Lengsfeld et al. (1995), Chu et al. (1998) Lengsfeld et al. (1995) Lengsfeld et al. (1995) Lengsfeld et al. (1995) Hariprasad et al. (1993) White et aL (1988), White et al. (1994), White et al. (1989) Wu et al. (1992) Ebert et aL (1997) Lavernia et al. (1991), Lavernia et al. (1987) Lavernia et al. (1991), Lavernia et al. (1987) Rickinson et al. (1981) Ikawa et al. (1990) Hanlon et al. (1997) Namba et al. (1993) Brooks et al. (1977a), Asano et al. (1991) Megusar et al. (1984) Rickinson et al. (1981), Ucok et al. (1991) Brooks et al. (1977b) Brooks et al. (1977b) Brooks et al. (1977b) Brooks et al. (1977b) Brooks et al. (1977b) Brooks et al. (1977b) Brooks et al. (1977b) Brooks et al. (1977a) Ibrahim et al. (1992) Tanigawa et al. (1986) Lawrynowicz and Lavernia (1994) Dalai and Pilchard (1993), Huron (1993), Dalai and Prichard (1992) Dalai and Prichard (1992), Bricknell (1986) Dalai and Pilchard (1992) Dalai and Pilchard (1992), Benz et al. (1993) Dalai and Pilchard (1992) Annavarapu and Doherty (1995) Dalai and Pilchard (1993), Fiedler et al. (1987) Dalai and Pilchard (1992), Ogata et al. (1986) Evans et aL (1985) Chu et aL (1988) Dalai and Pilchard (1993), Dalai and Pilchard (1992) Dalai and Pilchard (1992) Underhill et al. (1995) Underhill et al. (1995) Liang et al. (1992a), Zeng et al. (1995), Vetter et al. (1990), Lawrynowicz et al. (1997a) Li et aL (1997) Morris et al. (1992), Li et al. (1996a) Morris et al. (1992) Delplanque et al. (1997)

Mg based alloys

Fe based alloys

Ni based alloys

7150A1 7090A1 7475A1 8009A1 8090A1 WeldaliteTM 049 QE22 Mg Mg-5.6Zn-0.3Zr Mg-8.4A1-0.2Zr M15 Fe-2.5C-6V Fe-0.8C-(3 ~ 5)Cr Fe-2.3C-18Cr-2Mo-IV Stainless steel: (10 ~ 34)Cr 9Cr-Mo 316L En. 3A En. 8 En. 16 En. 19 En. 24 En. 31 En. 33 En. 58J SAE1020 Others: Fe- 16Nd-8B Fe3A1? R e n g f a m i l y : Ren6 41

H i g h s p e e d tool steel:

Ren6 80 Ren6 88DT Ren6 95 Ren6 125 Inconelfamily:

IN 625

IN 718 Others: Mer176 Nimonic 115 Astroloy Waspaloy AF2 - 1DA-6 UDIMET 720 MAR-M-200 Ni3A1t

Refractory and reactive alloys

Zircaloy-4 TiA1t TiA13t Ta-14.6Fe

* For detailed information on their nominal compositions of these commercial alloys, see corresponding references as well as (Brandes and Brook, 1992). t Intermetallics, in atomic percent.

c

4

0

Table 7.3 Novel commerciaVexperimenta1alloys developed using spray-forming Category

Alloy composition

High strength alloys

N202: AlL(6.1 6.S)Cu-(1.7 1.9)Mn-(0.4 O.S)Mg-(0.9 l.l)Zr(V, Ti) ES2x: AI-(5 6)Cu-(O.S 0.8)Mn-0.5Mg-0.5Ag-l .OZr 1.6)Ni-(1.0 1.2)MgN403: AL(17.5 19)Si-(4.2 4S)Cu-(l. IOFe(Mn, Cr, Ti, Zr, V) AI-20Si-SFe-2Ni AI-20Si-SFe

-

-

References

-

-

-

-

-

-

-

AI-20Si-3.1Cu-1.3Mg-0.3Fe

-

-

-

N707: Al-( 10.8 11.4)Zn-(2.2 2.S)Mg- 1.0 1.2)Cu-(0.25 (O.O-O.Z)Si(Fe) Eural: AlL(10 13)Zn-(S 6.5)Mg-(Cr,Mn, Zr) Eura2: same as Eural, but Cr/Mn content reduced

-

AI-12.4Zn-1.94Mg-1.9SCu-O.lZr Light weight high strength alloy

High strength high electrical conductivity alloys Magnetic alloys

Al-1 lZn-2Mg-lCu-0.3Zr UL30: A1-4Li-lCr-0.7Mg-0.3Zr UL40: Al-4Li-O.2Zr Al-4Li- 1Mg-OSGe-0.2Zr AI-3.5 1Li- 1.08Cu-0.63Mg-0.37Zr Cu-1SNi-SSn Cu-(0.1 0.8)Zr Fe-1SNd-8B Fe-2OCo- 1 SNd-8B

-

-

0.32)Zr-

Report (19904 White et al. (1994) Report (1990b) Zheng and Kahl(l990) Zhou et al. (1991) Estrada and Duszczyk (1990) Report (1990~)

a

3

R

de Sanctis (1991) de Sanctis (1991) Lengsfeld et al. (1995) Machler et al. (1991) White et al. (1994), White et al. (1993) Palmer et al. (1993) Gupta et al. (1990). Chanda et al. (1991) Zhang et al. (1989) Herman and Moms (1994). Hermann et al. (1997) Singh et al. (1991) Tanigawa et al. (1986), Harada et al. (1991) Chin et al. (1986)

.h

171

Spray-Forming

Table 7.4 Metal/intermetallic matrix composites fabricated using spray-forming Matrix

Pure A1

TiA1

Singer and Ozbek (1985), Willis (1988), Gupta et al. (1989), Gupta et al. (1990), Wu and Lavernia ( 1991), Llorca et al. (1994), Ibrahim et al. (1991) Wu et al. (1992), Zhang et aL (1993a), Gupta et aL (1993a), Llorca et al. (1993), Kahl and Leupp (1990), Carbalho et al. (1992) Kim et al. (1994) Gupta et al. (1993b), Gupta et aL (1991a), Gupta et al. (1991b), Gupta et al. (1992a) Wu and Lavernia (1991), Kahl and Leupp (1990), Gupta et al. (1992a), Perez et al. (1992), Zhang et aL (1994), Zhang et al. (1993b), Perez et al. (1993) Wu et al. (1994), Wu and Lavernia (1992), Srivatsan and Lavernia (1994), Gupta et al. (1992), McLelland et al. (1991) Gupta et al. (1992c) Maher et al. (1990) Noguchi et al. (1995) Majagi et al. (1992), Perez and Morris (1994) Bedford et al. (1993) Baskin et al. (1994) Lawley and Apelian (1994), Liang et al. (1992b), Lawrynowicz et al. (1997b) Li and Lavemia (1997)

A1203, graphite, SiC, TiB2, TiC, VC, B4C, WC, Cr3C 2 fiber: steel Mo

Lawrynowicz and Lavernia (1994) Lawrynowicz and Lavernia (1994)

A1-Cu

A1-Fe A1-Li

A1-Mg-Si

A1-Si

A1-Ti A1-Zn-Mg-Cu Mg alloys Cu alloys Fe alloys CoSi Ni3A1

Reinforcement

particulate:

172

Bing Li and E. J. Lavernia

I non-equilibrium [ conditions I 'refined microstructure: grains, secondary phase particles

,

1

absence of macrosegregation minimized microsegregation

T

improved mechanical properties 1-"

[near net shape I I simplicity I ._1potential cost !_. --[ effectiveness ]-Figure 7.12.

Interrelationshipbetween differentfeatures of spray-forming.

(f) short term exposure of reinforcement to elevated temperatures during fabrication of composites using spray-forming and co-injection, which may ameliorate interfacial deterioration; (g) penetration of reinforcement particles into droplets in spray-forming and co-injection technique, which allows extension of the reinforcement particles distribution into the interior of the matrix powders. This enhances the homogeneity of the particle distribution in the matrix. In conventional powder metallurgy techniques, distribution of the particles is limited to exterior of the matrix powders. While motivations (a) and (b) apply to all material systems, motivations (f) and (g) apply exclusively to metal/intermetallic matrix composites. Motivations (c), (d) and (e) are applicable to other specific material systems as described above.

7.5. NON-EQUILIBRIUM PHENOMENA IN SPRAY-FORMING

In terms of the technique itself, the characteristics of spray-forming include simplicity and potential cost effectiveness. In terms of the product, characteristics of spray-forming include refined microstructure, absence of macrosegregation, minimized microsegregation, improved mechanical properties, and near net shape capability. These characteristics are interrelated to each other, as shown schematically in Fig. 7.12, and are all associated with the non-equilibrium conditions of this technique.

Spray-Forming

173

7.5.1 Non-equilibrium nature

7.5.1.1 Rapid solidification in atomized droplets The liquid droplets formed in the spray-atomization stage transform into solid powders as they travel down to the bottom of the chamber. The solidification experienced by the droplets during flight is generally not an equilibrium process. Equilibrium solidification is a process in which the solidification rate is infinitely slow such that complete diffusion in the liquid and in the solid are ensured. Such a process is idealized and could never be realized in practical situations. However, when the solidification rate is relatively slow, such as in ingot casting, the process may be approximated as equilibrium solidification. The solidification rate in the droplets is so rapid that this approximation is no longer valid. Such a rapid solidification rate results from the high cooling effect and the large undercooling which may be experienced by the droplets.

Large undercooling. The very first step of the solidification process is nucleation, i.e., the formation of at least one solid nucleus. (1) The temperature at which the nucleation can be initiated is the nucleation temperature. This temperature is always lower than the liquidus temperature, and the temperature difference between these two is the so-called undercooling (prior nucleation undercooling), A T. It represents the driving force for solidification. In solidification, the solid/liquid interface moves towards the liquid at a velocity (2) U, which may be directly related to the undercooling as follows (Levi and Mehrabian 1982): U =~AT

(7.1)

where ~ is a constant. Eq. (7.1) states that a large undercooling level will lead to rapid solidification. The magnitude of the undercooling level is determined by the physical/thermo properties of the melt, the cooling rate, and most importantly, the nucleation mode (i.e., homogeneous vs. heterogeneous nucleation) (Flemings 1974, Kurz and Fisher 1986). In homogeneous nucleation, the solid nucleus is formed out of the liquid independently; while in heterogeneous nucleation, the solid nucleus is formed by attaching the solid atoms to the surface of heterogeneous nucleants (Flemings 1974, Kurz and Fisher 1986). The heterogeneous nucleants are generally solid materials that come into contact with the liquid, including the container-wall and inclusions in the melt, as is often the case in conventional solidification techniques. According to classic nucleation theory, the undercooling level achievable in homogeneous nucleation is always larger (sometimes significantly larger) than that in heterogeneous nucleation (Flemings 1974, Kurz and Fisher 1986). Consequently, in order to achieve large undercooling levels, homogeneous nucleation mode should be promoted. During spray-atomization, the melt is disintegrated into micron-sized droplets. A large proportion of these droplets undergo solidification during flight. Accordingly, container(l) A nucleus is a solid embryowith size larger than a specific critical value. (2) The interface velocity determines the solidification rate. The solidification rate is defined in the literature as the fraction of solid formed per second. In general, a larger interface velocity corresponds to a larger solidification rate. But the explicit form for their interrelationship depends on the geometryof the solidifying system (e.g., spherical droplet or rectangular block) and the geometryof the interface (planar or dendritic).

174

Bing Li and E. J. Lavernia

wall induced heterogeneous nucleation is avoided during solidification. Moreover, the probability for a droplet to contain heterogeneous inclusions decreases as the droplet size decreases. This may be illustrated by a simple calculation. Suppose the number of inclusions in 1 gram of molten metal is Nt. As this amount of melt is atomized into droplets with different sizes, the average weight of the droplets will be Wa grams, and the total number of droplets may be approximated as 1/ Wa. It is evident that the finer the droplets are, the smaller the average droplet weight will be, hence the larger the total number of droplets. When the droplet size is finer than a critical value, 1/ Wd will be much larger than iV,. Therefore, there will be a large proportion of droplets that are free of inclusions. These droplets are able to experience homogeneous nucleation, attain large levels of undercooling, and hence achieving high interface-velocities/solidification-rates. This concept was first employed by Turnbull and Cech (1950) to investigate undercooling phenomena. Once solidification begins, the heat of fusion is released by the liquid during the phase change reaction. If the energy released cannot be effectively dissipated, it will lead to an increase in the temperature of the melt, a decrease in undercooling, and hence a decrease of the interface-velocity/solidification-rate according to Eq. (7.1). Therefore, to ensure high solidification rate during the entire period of solidification, it is necessary to maintain a high cooling rate as well. High cooling effects. In atomization, the elevated temperature droplet loses thermal energy to the surrounding gas environment via a combination of convection, and radiation. The amount of energy dissipated by convection during a period of dt may be expressed as"

Qlds -- h ( T - To)S dt

(7.2)

where T and To are droplet and gas temperature, respectively; S is the area of the interface between the droplet and the atomization gas, which is mathematically equivalent to the droplet surface area; h is the convection heat transfer coefficient, which may be estimated as follows (Ranz and Marshall 1952): t~ ~fl/2 MI/3

h - ( k g / D ) ( 2 + 0 .... Re ~'Pr )

(7.3)

where NRe(--Opgt)dg/#g) and Npr(=#gCpg/kg) are the Reynolds number and Prandtl number, respectively; kg and Cpg are the thermal conductivity and specific heat capacity of gas, respectively; # g the gas viscosity; D droplet diameter; pg gas density; and Vctg droplet-gas velocity difference. The amount of energy dissipated by radiation during the time period dt is given by:

Q2ds -- 6~

- T~)S dt

(7.4)

where e is the emissivity of the droplet, and cr the Stefan-Boltzmann constant. Other variables have the same meaning as defined earlier. Compared with the convection term (Qls), energy loss by radiation is small when the droplet temperature is not very high (<1000~ Accordingly, in many situations, the radiation term is neglected.

Spray-Forming

175

The consequences of energy dissipation from the droplet to the gas environment are as follows: (a) Cooling of the liquid droplet; A decrease of dT in droplet temperature during the time period dT corresponds to a thermal energy decrease of: Q1

rl -- --Pd VdCpl dT

(7.5)

where Pd is the droplet density, Vd droplet volume, Cpl the specific heat capacity of the liquid droplet. The negative sign is incorporated since the temperature decreases, i.e., dT <0. (b) Solidification of the droplet when it is cooled down to a certain temperature below the liquidus of the droplet; The solidification of a fraction, d f , releases the following amount of energy: Q21 -- H f (pd Vd d f ) (7.6) where H i is the heat of fusion. If there is a solidification temperature range, as is the case of alloy systems, the temperature may decrease as well. Accordingly, the total energy released by the droplet is Q2t rl -- Hf(pdVd d f ) - pdVdCpm dT

(7.6)'

where Cpm is the specific heat capacity of the partially solidified droplet. (c) Further cooling of the solidified droplet following the completion of solidification; The corresponding thermal energy decrease during this period is:

Q3rl - --Pd VdCps d r

(7.7)

where Cps is the specific heat capacity of the solid droplet. Energy conservation requires:

Q1 Q2 ds 4- ds --

I Q~l 2 2t Qrl(Qrl) Qr31

prior to solidification solidification

(7.8)

after solidification

Although it is almost impossible to derive a general solution to Eq. (7.8), an order of magnitude estimate on each of the variables in this equation can lead to an important finding, that the small size of the droplet (D) is the key to achieving a high cooling rate and high solidification rate. For example, during solidification, if the temperature of the droplet remains unchanged (such as eutectic or peritectic solidification) and the radiation effect is neglected, then the energy balance equation becomes:

Q1ds _ Q2rl

(7.9)

df _ 6 kg (T - T0)(2 + 0. t;Arl/2 M1/sa dt - H f pd D D "~'Re ~'Pr J

(7.9) t

or explicitly,

176

Bing Li and E. J. Lavernia

For many metals and processing conditions, H f

,~

3

105 J/kg, 1/2

Pd

~

103 kg/m 3,

1/3

kg ~ 10 -2 W/mK, (T - To) "~ 10 K, and (2 + 0.6Nge Npr ) ranges from 2 to 100. Accordingly, d f / d t is in the range of [(10 -7 ~ 10-5)/D2]s-1, where D is in the units of

m. For spherical droplets, the interface velocity may be roughly estimated as follows: U = D(df/dt)

(7.10)

Therefore, the interface velocity is in the range of [(10 -7 -~ 10-5)/D] m/s. In the case of droplets of about 500/zm, the interface velocity ranges from 0.2 to 20 mm/s. Such a solidification rate may or may not be considered as rapid solidification, depending on the precise definition. For a droplet about 5/zm, the corresponding interface velocity is in the range of 20 ~ 2000 mm/s. Such a solidification rate will undoubtedly fall in the so-called rapid solidification regime. Effects o f rapid solidification. Rapid solidification permits large departures from equilib-

rium, resulting in: (a) formation of non-crystalline phases or metastable crystalline phases; (b) extension of solid solubility beyond the maximum equilibrium value; For example, the maximum equilibrium solid solubility of Cr in A1 is 0.44 at.% at 935 K. Upon rapid solidification, it may be extended to exceed 5 at.% (Jones 1978), and (c) Retention of disordered crystalline structures in normally ordered materials and intermetallic compounds. Rapid solidification may also change the microstructure characteristics, including: (a) modification of phase morphology; (b) refinement of dendrites, dendrite arms, and dendrite arm spacing; (c) refinement of second phase particles (constituents, dispersoids and precipitates). 7.5.1.2 Transient semi-solid layer in deposition stage

As stated earlier, spray-forming includes two stages: spray atomization and spray deposition. The rapid, and by definition non-equilibrium, solidification kinetics that are inherent to the atomization stage were examined in the preceding section. In this section, the non-equilibrium conditions present in the deposition stage will be discussed. During deposition, liquid, solid and partially solidified droplets continuously arrive on the top surface of the deposit/substrate. Identification of the precise mechanisms that govern the formation of the microstructure during deposition has been the subject of intense research activities for many years (Liang et al. 1992a, Estrada and Duszczyk 1990, Singer 1982, Singer and Evans 1983, Annavarapu et al. 1988, Gutierrez et al. 1988, Lavernia 1989, Mathur et al. 1989, Annavarapu et al. 1990, Zhuang et al. 1990, Feng et al. 1992, Leatham and Lawley 1993, Annavarapu and Doherty 1993, Liang and Lavernia 1994). It is now believed that the impinging droplets form a thin semi-solid layer on the top surface of the deposit while the rest part of the deposit underneath the semisolid layer remains solid (see Fig. 7.13). This layer experiences/completes solidification and becomes the main part of the deposit, while newly impinging droplets form another

177

Spray-Forming

gas ~ . . . . . . . ~

,~ Q

"1 II~ I

liquid

!

!

I II I

I dendrite~

II

" I~ I

I

,

~

I

I

I I I

I I I |

I

\ \

l

\

Semi-solid

'ate

Solidified deposit 1

9

deposit T = 1 -

10 K/s

F i g u r e 7.13. Semi-solid layer on the top of the deposit (Liang and Lavernia 1994).

thin semi-solid layer. Such a process continues until no more droplets impinge, i.e., at the end of spray-forming. The main part of the deposit underneath the semi-solid layer experiences cooling in the solid state, with no solidification involved. Experimental measurements and numerical studies indicate that the cooling rate during deposition is in the range from 1 -~ 10 K/s (Liang et al. 1992a, Lavernia and Grant 1988, Gutierrez et al. 1988, Mathur et al. 1989). Such a cooling effect is higher than that in ingot casting (<1 K/s) (Lavernia et al. 1992) but lower relative to that in rapid solidification (103 ~ 106 K/s) (Liang et al. 1992a, Mathur et al. 1989, Li et al. 1996b, Gutierrez et al. 1989, Grant et al. 1993b). Accordingly, the cooling process in the main part of the deposit is normally not considered as non-equilibrium. However, the semi-solid layer represents an inherently non-equilibrium condition in several senses. Firstly, when the high velocity droplets arrive and then join the relatively quiescent top surface of the deposit/substrate, they lose their kinetic energy. A portion of this kinetic energy is absorbed by the semi-solid layer, resulting in violent convection in the layer. Secondly, some of the kinetic energy loss is absorbed by the droplets when they experience extensive deformation. Such deformation may include (Liang et al. 1992a, Lavernia 1989): (1) change on the geometrical shape of the droplets from sphere to disk; and (2) deformation and fragmentation of the dendrites formed in the droplets during the atomization stage.

178

Bing Li and E. J. Lavernia

These phenomena dramatically alter the solidification process which is experienced by the droplets. Thirdly, recent results indicate that each discrete thin semi-solid layer completes its solidification in a relatively short time (10 -1 ~ 10 -3 s) because of its thinness and the effective cooling effect of the environmental gas (Liang and Lavernia 1994). Corresponding interface velocity is estimated to be approximately 1 mm/s. Moreover, the cooling rate experienced by the thin layer is in the range of 10 ~ 105 K/s, depending on the materials being spray formed and the processing conditions (Liang and Lavernia 1994). Such a cooling rate and solidification rate may be considered, to a certain extent, to be non-equilibrium. 7.5.2 Non-equilibrium related features in the deposit

In Section 7.5.1, the non-equilibrium conditions that are present during spray-forming were examined. Manifestations of such a non-equilibrium state in the deposit include: metastable phases, extended solid solubility, absence of macrosegregation/minimized microsegregation, as well as refinement of grains and secondary phase particles; these will be discussed individually below.

7.5.2.1 Metastable phases Since atomized droplets experience rapid solidification during their flight, it is anticipated that some metastable phases will form prior to impingement. During the deposition stage, especially in the main part of the deposit, cooling is not as effective as in the atomization stage, allowing the material to be exposed to high temperatures for a relatively long period. Accordingly, there is a possibility for the metastable phases to be transformed into thermodynamically stable phases. Whether these metastable phases will be retained in the deposit depends on: (i). the tendency of metastable phase formation in the material; and (ii). the kinetics of the metastable phase transformation. The former determines the amount of metastable phases present in the material prior to a metastableto-stable phase transformation. The latter affects the amount of metastable phases that are left in the deposit. Experimentally, metastable phases have been heretofore noted in several spray-formed alloys, including A14Mn and ot'-A13Zr phases in A1-6.5Mn2.35Li-0.65Zr (Ruhr et al. 1990, Baram 1991), ot'-A13Zr phase in 7xxx A1 (Machler et al. 1991), 6-A14FeSi2 phase in A1-20Si-5Fe (Zhou et al. 1991), amorphous/quasicrystalline particles in 2024 A1 (Song et al. 1994), as well as Cu2Sb intermetallic phase in Pb-12Sn- 11.5Sb- 1.25Cu-0.75Ni-0.3Cd-0.2As (Upadhyaya et al. 1997). 7.5.2.2 E x t e n d e d solid solubility

The formation of extended solid solution is another characteristic phenomenon of rapid solidification, which might occur during the atomization stage and then modified during the deposition stage. Whether the extended solid solubility is retained in the deposit or not depends on the relevant precipitation kinetics. Experimentally, it has been reported that many spray-formed alloys exhibit solid solubility levels that are much higher than the corresponding equilibrium values at room temperature (Lavernia et al. 1991, Gupta et al. 1992c, Liang and Lavernia 1994, Ruhr et al. 1990, Baram 199 l, Juarez-Islas et al. 1994). For example, the room temperature equilibrium solid solubility of Zn in Mg is

179

Spray-Forming Table

7.5 Extendedsolid solubilityin spray-formedalloys

Alloy (wt.%)

Solid solubility in deposit (wt.%)

A1-6.5Mn-2.3Li-0.65Zr A1-Ti A1-Fe

2.1 "~ 2.25 (Mn) 0.8 ,~ 1.1 (Ti) 2.2 (Fe)

Maximumequilibrium References solid solubility (wt.%) 1.82 (Mn) 1.00 (Ti) 0.025 (Fe)

Ruhr et al. (1990), Baram (1991) Gupta et al. (1992c) Liang and Lavernia (1994)

less than 1.0 wt.% (Lavernia et al. 1991). In a study on Mg-5.6wt.%Zn-0.3wt.%Zr alloys, the room temperature solid solubility of Zn in Mg in the spray-formed Mg-Zn-Zr deposit was found to be 5.3 wt.% (Lavernia et al. 1991), considerably higher than the equilibrium value. However, in this study, the solubility (5.3 wt.%) is still lower than the maximum equilibrium solid solubility of Mg-Zn alloy system, which is 6.2 wt.% at the eutectic temperature (340~ (Lavernia et al. 1991). In extreme cases, the solid solubility in the deposit may be extended beyond the maximum equilibrium solid solubility, as summarized in Table 7.5. Such an extended solid solubility is comparable to that normally observed in relatively more conventional rapid solidification techniques, such as spray-atomization and melt-spinning (Jones 1978). However, compared with these techniques, the extension of solid solubility in spray-forming is generally less pronounced. For example, in the study of A1-6.5Mn-2.3Li-0.65Zr alloy, while the solid solubility of Mn in the deposit is approximately 2.1 ~ 2.25 wt.% (Ruhr et al. 1990), the corresponding value under similar processing conditions in the atomized powders is in the range of 4.25 ~ 6.00 wt.% (Ruhr and Baram 1991).

7.5.2.3 Absence of macrosegregation/minimized microsegregation While the presence of metastable phases and the extension of solid solubility in spray formed alloys are material dependent, the absence of macrosegregation and minimized microsegregation are inherent to spray-forming technology. These characteristics are achieved through a combined effect of the non-equilibrium conditions that are present in the atomization and the deposition stages. During spray atomization, the droplet is small and the solidification rate is high. Accordingly, segregation, such as that associated with the formation of dendrites, is limited to a micron size scale. Fragmentation of dendrites during the deposition stage minimizes microsegregation. It also plays an important role in multiple nucleation in the semi-solid layer. Multiple nucleation is important in preventing the appearance of macrosegregation during the deposition stage. Finally, the relatively slow cooling rate experienced by the solidified deposit underneath the semi-solid layer exposes the material to an elevated temperature for a relatively long period, allowing certain homogenization of the microstructure of the deposit. Such a homogenization effect further minimizes the microsegregation present in the deposited materials.

180

Bing Li and E. J. Lavernia

Figure 7.14. Optical micrograph showing fine equiaxed grains in spray-formed 7075A1 (Lavernia 1989).

7.5.2.4 Refinement Equiaxed grains. A salient feature revealed by numerous studies is the presence of fine, equiaxed grains (see Fig. 7.14) in the deposit, as summarized in Table 7.6. Again, the presence of fine, equiaxed grains in spray-formed materials is directly related to the nonequilibrium conditions present both in the atomization and in the deposition stages. As mentioned earlier, extensive efforts have been dedicated to the understanding of the evolution of such fine, equiaxed grains (Liang et al. 1992a, Annavarapu et al. 1988, Lavernia 1989, Grant et al. 1991, Grant 1995). While there is some controversy on the details of the fundamental mechanisms, it is widely agreed that formation of fine dendrites during atomization stage, and fragmentation of these dendrites during deposition stage, play key roles in the formation of the fine, equiaxed grains. The formation of fine dendrites in the droplets is mainly attributable to the large undercooling, the high cooling effect, and the rapid solidification present during atomization; all of them represent nonequilibrium conditions. Fragmentation of these fine dendrites, on the other hand, is itself a non-equilibrium phenomenon, as discussed in Section 7.5.1.2. Secondary phase particles. In conventional ingot casting, solute segregation is coupled with a slow solidification/cooling rate, leading to the formation of coarse secondary phase particles. In spray-forming, solute segregation is minimized; the rapid solidification experienced by the droplet tends to refine the microstructural characteristics; the high cooling rate in the semi-solid layer and the intermediate cooling rate in the solidified deposit limits the time for secondary particles to precipitate and grow. All of these characteristics of spray-forming lead to refinement/suppression of coarse secondary phase particles, as

Spray-Forming

181

Table 7.6 Grain size in spray-formed materials Alloy (wt.%) Sn-38Pb Mg-8.4A1-0.2Zr Mg-5.6Zn-0.3Zr A1-4Cu A1-4Si 2014A1 2618A1 5083A1 6061A1 A390 Fe-15Nd-8B Cu-6Ti IN625 IN718, Rene 95 Rene 80 316L stainless steel 316L stainless steel 0.18% plain carbon steel SAE 1020 steel Ni3AI* TiAI* Zircaloy-4

Grain size (/zm) 4,-~10 20~50 50~70 75 45 17 10~25 15 24~44 12.5 0.1~5 46 15 30 10~25 10~40 100 10,-~40 10~50 11 30~60 175t

References Lavernia and Wu (1996), Bewlay and Cantor (1991) Lavernia and Wu (1996), Lavernia et al. (1991) Lavernia and Wu (1996), Lavernia et al. (1991) Lavernia and Wu (1996), Grant et al. (1991) Lavernia and Wu (1996), Wu et al. (1993) Annavarapu and Doherty (1995) Lavernia and Wu (1996), Underhill et al. (1993) Wu et al. (1997) Wu and Lavernia (1991) Wu et al. (1995) Veistinen et al. (1987) Lavernia and Wu (1996), Annavarapu and Doherty (1993) Annavarapu and Doherty (1995) Lavernia and Wu (1996), Fiedler et al. (1987) Lavernia and Wu (1996), Bricknell (1986) Lavernia and Wu (1996), Ucok et al. (1991) Lavernia and Wu (1996), Lawley and Cantor (1986) lbrahim et al. (1993) Ibrahim et al. (1992) Lavernia and Wu (1996), Liang et al. (1992a) Morris et al. (1992), Li et al. (1996a) Li et al. (1997)

* In atomic percent. t Grain size of ingot counterpart is larger than 1000/zm.

well as concomitant modification to their morphology. This has been verified experimentally in several studies, as summarized in Table 7.7. Meanwhile, many studies reported the presence of superfine secondary intermetallic particles in spray-formed materials (see Table 7.8). The refined characteristics of these phases are beneficial to the mechanical properties of the material.

7.5.3 Effects of non-equilibrium features on mechanical~physical properties In the development of spray forming, extensive efforts have been devoted to the evaluation of properties of spray-formed materials, including yield strength, ultimate strength, tensile elongation, elastic modulus, fracture toughness, creep resistance, fatigue strength, electrical conductivity, magnetic properties, and damping capacity. The impetus behind these efforts is the belief that the non-equilibrium related features may have a significant effect on the physical/mechanical properties of spray-formed materials, either directly or indirectly. Direct effects are the difference in mechanical properties resulting from non-equilibrium features alone, with no modifications to the material's chemical composition. Indirect effects involve the difference in mechanical properties resulting from

Bing Li and E. J. Lavernia

182

Table 7.7 Effect of spray-forming on the secondary phase particles Alloy

Particle

Processing*

Morphology

Size (/zm)

A390

Primary Si

IM

Block

30~ 100

Eutectic Si

SF IM SF

Particulate Flake-like elm. ~t

5 Coarse elm.

IM SF IM SF IM

Needle elm. Faceted Non-faceted Needle

Coarse elm. 50~ 100 3~ 1 10"~30

SF

elm.

elm.

IM

Particulate

15~20

SF

Particulate

5 ~6

7075A1+ 1Ni+0.8Zr

A13Zr/A13Ni

Babbitt alloyt High speed tool steel

SnSb MC2, M6C

MC

References Lavernia and Wu (1996), Wu et al. (1995) Wu et al. (1995) Wu et al. (1995) Wu et al. (1995) Lavernia and Grant (1986) Lavernia and Grant (1986) Upadhyaya et al. (1997) Upadhyaya et al. (1997) Ikawa et al. (1990), Leatham et al. (1995), Leatham et al. (1985) Ikawa et al. (1990), Leatham et al. (1995), Leatham et al. (1985) Ikawa et al. (1990), Leatham et al. (1995), Leatham et al. (1985) Ikawa et al. (1990), Leatham et al. (1995), Leatham et al. (1985)

* SF: spray-forming; IM: ingot metallurgy. t Pb- 12Sn- 11.5Sb- 1.25Cu-0.75Ni-0.3Cd-0.2As. $ elm.: eliminated.

Table 7.8 Superfine secondary intermetallic particles in deposited materials* Alloy (wt.%) 2618A1 2618+(0.5 -v 1.0)(Fe, Ni) A1-11Zn-2Mg- 1Cu-0.3Zr Fe-0.8C-3Cr 7xxx

Particle t

Size

A19NiFe AI9NiFe AI7Cu4Ni A13Zr Fe3C, Cr3C r/-MgZn2

--~3/zm 0.5--~3.0/xm 0.5,~3.0/zm 25~45 nm 134 nm ~0.25/xm

* All in as-spray-formed condition. t The particle morphology is normally equiaxed.

References Llorca et al. (1993) Underhill et al. (1993) Underhill et al. (1993) Machler et al. (1991) Hanlon et al. (1997) Lengsfeld et al. (1995)

Spray-Forming

183

Table 7.9 Effect of spray-forming on tensile mechanical properties Alloy

Proces sing*

cry S

crU T S

(MPa)

(MPa)

Elongation References KIC (MPaml/2) (%)

2024A1

SF+T4

402~417 562~586

15~16

31

2024A1 Weldalite TM 049 Weldalite TM 049

IM+T4 SF+T6 IM+T6

310~329 441~469 638 665 500 540

19~20 12 16

31

A390A1 A390A1

SF+Extr.+T6 IM+Extr. +T6

506 374

532 387

3.0 1.3

6061 A1 6061A1

SF+T6 IM+T6

313 280

352 310

15.2 13

Cu-0.1Zr Cu-0.1Zr

SF+TMP IM+TMP

441 390

462 450

6 12

Lavernia et al. (1986) Lavernia et al. (1986) Wu et al. (1992) Wu et al. (1992) Wu et al. (1995) Wu et al. (1995) Wu and Lavernia (1991) Brandes and Brook (1992) Singh et al. (1991) Singh et al. (1991)

* SF: spray-forming; IM: ingot metallurgy; Extr." extrusion; T6: solution heat treated and then artificially aged; T4: solution treated and aged naturally to a substantially stable condition; TMP: thermo-mechanical processed.

modification to the chemical composition of the material, which is made possible by the non-equilibrium conditions present.

7.5.3.1 Direct effects It has been experimentally noticed that in the case of numerous alloy compositions that spray-formed materials exhibit better mechanical properties when compared to those of the ingot metallurgy counterparts. Representative results of the tensile properties are summarized in Table 7.9. It is evident that the yield strength and the ultimate strength values are significantly increased upon spray-forming, while the ductility is either improved or degraded. The latter situation may be acceptable if the loss of ductility does not lead to brittle behavior (<2%). Such an improvement in mechanical properties may be associated with the refinement in grains and secondary phase particles, as well as the absence of macrosegregation and the minimized microsegregation. Refined equiaxed grains formed during spray-forming may lead to improvement of the tensile properties, especially the yield (ultimate) strength, because of a decrease in grain size. The well-known Hall-Petch relationship predicts: ffs - - O'so -k- X D g

n

(7.11)

where as is the strength; as0 the internal stress; X and n are experimental constants; and Dg is the grain size. Since the grain size of spray-formed materials is considerably smaller than that present in ingot materials, there is a potential for the strength of the deposit to be improved. A reduction in grain size also decreases the dislocation pile-up length and consequently lowers the stress concentration at grain boundaries and grain boundary triple junctions, hence promoting homogeneous deformation and preventing early crack nucleation and premature failure due to strain localization effects (Perlinde and Lutjering 1982, Lavernia et al. 1990).

184

Bing Li and E. J. Lavernia

Refinement of the secondary phase particles (constituents, dispersoids and precipitates) also directly influences the mechanical properties. Depending on the degree of coherency, there are two types of secondary phase particles: coherent and incoherent. The strengthening mechanisms of these two types of particles are different. For incoherent particles, moving matrix dislocations cannot penetrate these particles. Accordingly, these particles will not be sheared along with the matrix. Instead, the moving dislocation is forced by the applied stress to bend around the particle and bypass it through the Orowan looping mechanism. The critical resolved shear stress due to this mechanism is inversely proportional to the particle size (Reppich 1993): 750R -- C1 (Gmb) v/ f p (7.12) r where C1 is a dimensionless constant, Gm the shear modulus of the matrix, b the Burgers vector, fp the volume fraction of particles, and r the particle radius. Therefore, refinement of secondary phase particles will lead to strengthening effect on the matrix for incoherent particles. For coherent particles, the moving dislocation can cut through them, i.e., they can be sheared along with the matrix. The critical resolved shear stress due to the cutting mechanism is given by (Reppich 1993): h3/2 P 2v/fpr ~c -- C2 ~/--~b

(7.13)

where C2 is another dimensionless constant, h p is a hardening parameter, and other variables have the same meaning as defined earlier. According to Eq. (7.13), coarse particles will have large strengthening effect. This is true only when the particle size is small. When the particle size becomes excessively coarse, Eq. (7.13) will break down for the following two reasons: (1) As the particle size is increased to a certain point, the coherent particle will become incoherent (Reed-Hill 1973). Normally, the lattice parameter of a secondary phase particle is different from that of the matrix. The interface is coherent if the crystallographic planes of the matrix may be epitaxially extended to the particle. The lattice mismatch causes certain strain on the crystallographic planes near the interface. However, as the particle size increases, the lattice parameter mismatch may no longer be accommodated by the strain. Accordingly, the crystallographic plane of the matrix will be terminated at the interface, i.e., the interface becomes incoherent. Under such a circumstance, looping mechanism will be operative. (2) For coherent particles, both cutting and looping mechanisms may be operative, depending on the required stress level (Reppich 1993). When the particle is very small, the stress predicted by Eq. (7.12) may be higher than that by Eq. (7.13). Accordingly, the cutting mechanism requires a lower stress level, and is operative. As the particle size increases, the stress required for the looping mechanism decreases, while that for the cutting mechanism increases. At a certain particle size, the stress level required by the looping mechanism will be less than that by the cutting mechanism. In this case, even though the particle is still coherent, the looping mechanism will be operative.

Spray-Forming

185

When looping mechanism becomes operative, the strengthening effect decreases with increasing particle size. Therefore, for coherent particles, there is a critical particle size at which the strengthening effect is at maximum level. For the classic precipitation hardening, this critical particle size corresponds to peak ageing condition (Reed-Hill 1973). In practice, this critical particle size is in the range of 0.1 ~ 10 #m. Such a fine size is exactly what may be achieved through spray-forming technology (see Table 7.7). In slow cooling/solidification, coarse secondary phase particles are generally segregated to the grain boundary region (Lavernia et al. 1990). The brittle nature associated with many of these secondary phase particles embrittles the grain boundary, making it extremely susceptible to crack nucleation. Such brittle grain boundaries favor the low energy intergranular fracture over high-energy transgranular dimple rupture, resulting in low fracture toughness and poor ductility (Lavernia et al. 1990). Meanwhile, such segregation is sometimes accompanied by a precipitate free zone (PFZ) near the grain boundary (Lavernia et al. 1990). The PFZ region has an inherently lower strength, leading to abrupt variation in local mechanical properties. It is also clear from the above discussion that the compositional uniformity (i.e., absence of macrosegregation and minimized microsegregation) in the deposited materials are critical to the mechanical properties as well. These characteristics decrease the chance of localization of flow stresses, leading to improvement on the mechanical properties, especially the crack resistance, the fracture toughness, and the ductility.

7.5.3.2 Indirect effects

In many alloy systems, the mechanical properties may be improved through effectively modifying the content and the type of alloying elements, provided that the alloying elements are uniformly distributed in the matrix. In ingot metallurgy, such alloying modification is hindered by the macrosegregation of solutes and by the formation of coarse primary or secondary phases. Especially, increased alloying is typically precluded. Sprayforming inherently avoids macrosegregation and microstructure coarsening. Moreover, it sometimes leads to extended solid solution in the deposited materials. Accordingly, sprayforming is an effective method in alloying modifications, especially when heavy alloying is involved. Representative examples in this category include A1-Cu, A1-Zn and Cu-Zr alloys. Commercial A1-Cu alloys (2xxx series) contain Cu, Mn, Cr, Si, and other elements from trace to significant amounts. Further increase in the alloying amounts may improve the available mechanical properties of this alloy system. Nevertheless, exploration/utilization of these potential benefits are difficult in ingot metallurgy because of strong solute segregation. In contrast, compositional uniformity in spray-formed materials enhances alloying flexibility, leading to the development of several heavily alloyed A1-Cu alloys, such as N202 alloy {Al-(6.1 ~ 6.5)Cu-(1.7 ~ 1.9)Mn-(0.4 ~ 0.5)Mg-(0.9 1.1)Zr(V, Ti)} and the family of ES2x alloys {Al-(5 ~ 6)Cu-(0.5 ~ 0.8)Mn-0.5Mg0.5Ag-I.0Zr(V, Ti)}. Comparisons of the mechanical properties of the commercial and novel A1-Cu alloys are provided in Table 7.10, Fig. 7.15 and Fig. 7.16. It is evident there that the mechanical properties of A1-Cu alloys were improved upon heavy alloying.

186

Bing

L i a n d E . J. L a v e r n i a

Table 7.10 Tensile properties of commercial and novel A1-Cu alloys Alloy N202 2618 A1 2014 Alt 2024 A1 2618 A1

Processing*

ffYS (MPa)

ffUTS (MPa)

SF+T6 SF+T651 IM+T6 IM+T6 IM+T6

490 418 385~465 395 330

560 445 440,~500 475 430

References Report (1990a) Llorca et al. (1993) Brandes and Brook (1992) Brandes and Brook (1992) Brandes and Brook (1992)

* SF: spray-forming; IM: ingot metallurgy; T6: solution heat treated and then artificially aged. t Depending on the geometry of the product: sheets, bars, rivet stocks, tubes or wires. 500

400

:--.-.7.---'-~ . . . . . . . . . . . . . . x~ "', x~k

-

"

t~

..........

----,

....

2014

AI

2 6 , s A~

ES20 ES21

300 G

.u ~"

100

0

Figure 7.15.

50

100

150 200 250 Temperature (*C)

300

350

Yield strength of ES2x alloy following thermal exposure at testing temperatures for 100 h (White et al. 1994).

400

380 ... 360 ~" 320 300

.~ 2 8 0

~

-

o'.

-

O

260 -

,

ES20. 176~

240

o

2618 AI. 170"C

220-

9

2618 AI. 176~

200 O.l

J

I

O

J

10

J

100

1000

Time to Rupture (hr)

Figure 7.16. Creep rupture strength of ES20 compared with that of 2618 A1 (White et al. 1994).

187

Spray-Forming

Table 7.11 Tensileproperties of commercial and high Zn content 7xxx alloys Alloy

N707 Eurall Eural2 A1-11Zn-2Mg-1Cu-0.3Zr 7075A1

Processing*

SF+T6 SF+T6 SF+T6 SF+T6 IM+T6

cryS

crUT S

Elongation

(MPa)

(MPa)

(%)

760 775 762"~790 798~810 807 819 705 719 583 600

8.0 2.2~4.9 7.1 29 9

References

Report (1990c) de Sanctis (1991) de Sanctis (1991) Machleret al. (1991) Domalavage et al. (1975)

High strength A1-Zn alloys (7xxx series) contain a significant amount of Zn, Mg and Cu which form major strengthening particles such as intermediate r/t and equilibrium r/phases. Further improvement on the tensile properties may be achieved through either refining grain size or increasing the amount of fine precipitates in the matrix. The latter involves increase in the Zn content to the theoretical limit as predicted in the equilibrium phase diagram while simultaneously maintaining proper ratios of Cu and Mg. In conventional ingot metallurgy, the Zn content is normally below 6.2 wt.%. An increase beyond this limit is difficult because of extensive macrosegregation and cracking associated with slow solidification. These problems, however, may be circumvented by spray-forming. Accordingly, several high Zn content 7xxx alloys have been developed using spray-forming, including N707 {AI-(10.8 ~ 11.4)Zn-(2.2 ~ 2.5)Mg-(1.0 ~ 1.2)Cu(0.25 ~ 0.32)Zr-(0.0-0.2)Si(Fe) }, Eural/Eura2 (3) {AI-(10 ~ 13)Zn-(5 --~ 6.5)Mg-(Cr, Mn, Zr) }, and AI-11Zn-2Mg-lCu-0.3Zr. Comparison of the mechanical properties between the commercially available A1-Zn alloys and the high Zn content alloys is provided in Table 7.11. As shown in Table 7.11, a significant improvement in tensile properties was achieved through an increase of Zn content using spray-forming. Cu-Zr alloys are of interest for applications that require high strength and high electrical conductivity. The high strength of this alloy comes from age hardening effects by fine, zirconium-rich precipitates. However, the equilibrium solid solubility of Zr in Cu is relatively low, with the maximum value being 0.124 wt.% at 950~ When the fine zirconium rich precipitates are produced through solution treatment and subsequent ageing, their amount is restricted by such a small solid solubility. Consequently, the achievable strength is limited. This may be further explained as follows. During ingot casting, solute segregation and slow cooling rate lead to the formation of coarse Zrrich particles in the matrix. Such coarse particles are detrimental, instead of beneficial, to the mechanical properties (see Section 7.5.3.1). When the Zr content is low, all of these coarse particles may be dissolved upon solution treatment. Upon ageing, fine particles are precipitated, enhancing the strength of the alloy. However, when the Zr content is beyond the maximum equilibrium solid solubility (e.g., 0.5 wt.% Zr), only a small amount of the coarse particles may be dissolved upon solution treatment, while others (3) The compositions for Eural and Eura2 are similar, but the content of Cr and Mn in Eura2 is less than in Eural.

188

Bing Li and E. J. Lavernia

Table 7.12 Tensile properties of commercial and high Zr content copper alloys (Singh et al. 1991) Alloy (wt.%) Cu-0.8Zr Cu-0.4Zr Cu-0.2Zr Cu-0.1Zr Cu-0.1Zr

Processingt

crys (MPa)

crUTS (MPa)

Elongation (%)

Conductivity (%IACS)*

SF+TMP SF+TMP SF+TMP SF+TMP IM+TMP

511 525 476 441 390

538 552 504 462 450

3 7 8 6 12

75 82 78 88 85

* IACS: Intemational Annealed Copper Standard. t SF: spray-forming; IM: ingot metallurgy; TMP: thermal mechanical processing, including cold rolling (70%), recrystallization (930 ~ cold rolling (90%), and ageing (400~615 ~

remain in the solvent matrix. Upon subsequent ageing, the remaining coarse particles act as pre-existent nuclei, making it unlikely to generate fine particles. Even though some fine particles are precipitated upon subsequent ageing, their beneficial effect is negated by the detrimental effect of the coarse particles, with the overall effect being hardly beneficial. Such problems inherent in conventional ingot casting may be circumvented through several non-equilibrium features present in spray-forming: extended solid solubility, eliminated macrosegregation/minimized microsegregation, and refined secondary phase particles. Table 7.12 summarizes the high Zr content Cu-Zr alloys produced heretofore by spray-forming, and their significantly improved tensile properties. From the preceding discussion, it is evident that the non-equilibrium characteristics of spray-forming have many direct or indirect beneficial effects on the properties of a material. It is to be noted that many of these features are also present in materials synthesized using other relatively conventional rapid solidification techniques, such as spray atomization and melt spinning. As a matter of fact, some of them such as refinement and extended solid solubility are more pronounced in these techniques than in spray forming because of the higher cooling rates associated with them. Consequently, it may be argued that products of another structural processing technique, powder metallurgy, will exhibit better physical/mechanical properties than spray-formed counterparts, since in many cases the powders used in this technique are produced using spray atomization. This may or may not be true, depending on many factors such as material chemistry and consolidation details. It is a challenge to ensure that the non-equilibrium features obtained in spray atomization are not eliminated in subsequent consolidation steps which normally involve long periods of exposure of the powders to elevated temperatures. Moreover, it is worth noting that, besides the high cost of powder metallurgy, oxidation and contamination (such as pick up of coarse ceramic inclusions during screening (Kim 1988)) problems in this technique may also prevent full attainment of the beneficial effects offered by the non-equilibrium conditions. Oxides present in a material may strengthen the material if they are superfine and uniformly distributed, as in ODS materials (see Section 7.3.1.4). However, the oxide films formed on the surface of powders in conventional powder metallurgy technology often leads to less than optimal fracture properties, if not properly dispersed (Kim 1988, Kim et al. 1985).

Spray-Forming 7.6. C O N C L U D I N G

189

REMARKS

Spray-forming is a novel but powerful processing technology suitable for the manufacturing of structural materials. It has many merits over well established processing techniques such as ingot casting and powder metallurgy. These merits include superior microstructures, improved mechanical properties, and cost effectiveness. Noticeably, many of these merits are associated with the non-equilibrium phenomena present in this technology. Compared with many other novel processing techniques which are being currently explored/utilized, the uniqueness of spray-forming is its capability of manufacturing structural materials in tonnage production levels. This enables it the potential to challenge parallel conventional processing techniques such as casting and powder metallurgy, arousing intensive interest in the industrial community. Spray-forming is also of great value to scientific research community. The rapid solidification effect experienced by the atomized droplets, the transient s e m i - s o l i d region on the top surface of the deposit, the spray-forming and co-injection of m e t a l / i n t e r m e t a l l i c m a t r i x composites, and the reactive spray-forming of d i s p e r s o i d - s t r e n g t h e n e d alloys, all represent areas rich with fundamental problems of interest to the scientific community. Over the last three decades, spray forming has evolved from a laboratory curiosity to a potentially revolutionary manufacturing technique. It is also evident, however, that our understanding of the relevant fundamental physical phenomena involved is, in many cases, inadequate, partly because of the complex interactions of the fluid, thermal, and solidification phenomena that are present during spray-forming. Fortunately, this has not hindered the application of spray-forming technology, as evidenced by the development of a variety of materials with unusual combination of microstructure and physical attributes.

ACKNOWLEDGMENTS

This work was financially supported by the Army Research Office (Grants No. DAAG5598-1-0088 and DAAG55-98-1-0141) and the National Science Foundation (CTS 9614653).

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Marinkovich, J. M., Mohamed, E A., Pickens, J. R. and Lavemia, E. J. (1989) JOM, 41(9), 36. Mathur, E, Apelian, D. and Lawley, A. (1989) Acta Metall., 37, 429. McLelland, A. R. A., Atkinson, H. V., Kapranos, E and Kirkwood, D. H. (1991) Mater. Lett., 11, 26. Megusar, J., Lavernia, E. J., Domalavage, P., Harling, D. K. and Grant, N. J. (1984) J. Nucl. Mater., 122/123A, 789. Morris, D. G., Morris, M. A., Gunter, S., Leboeuf, M. and Hollrigl, G. (1992) Scripta Metall. Mater., 27, 1645. Namba, Y., Saito, M., Matsushita, T. and Takigawa, H. (1993) in Spray Forming 2, ed. Wood, J. V. (Woodhead Publ., Cambridge, UK), p. 221. Noguchi, A., Ezawa, I., Kaneko, J. and Sugamata, M. (1995) J. Japan Inst. Light Metals, 45, 64. Ogata, K., Lavemia, E. J., Rai, G. and Grant, N. J. (1986) lnternat. J. Rapid Solid., 2, 21. Palmer, I. G., Chellman, D. J. and White, J. (1993) in Spray Forming 2, ed. Wood, J. V. (Woodhead Publ., Cambridge, UK), p. 385. Parker, D. W. and Kutner, G. L. (1991) Advanced Mater. & Proc., 139, 68. Perez, J. E and Morris, D. G. (1994) Scripta Metall. Mater., 31, 231. Perez, R. J., Zhang, J. and Lavemia, E. J. (1992) Scripta Metall. Mater., 27, 1111. Perez, R. J., Zhang, J., Gungor, M. N. and Lavernia, E. J. (1993) Metall. Trans., 24A, 701. Perlinde, G. and Lutjering, G. (1982) Metall. Trans., 13A, 1283. Picraux, S. T. (1984) Ann. Rev. Mater. Sci., 14, 335. Ranz, W. E. and Marshall, W. R. (1952) Chem. Eng. Prog., 48, 141. Rawers, J. C. and Wrzesinski, W. R. (1992) J. Mater. Sci., 27, 2877. Reed-Hill, R. E. (1973) Physical Metallurgy Principles (2nd edition, Von Nostrand, Princeton, NJ). Report, "Spray Deposited High Performance Materials: N202", Alusuisse-Lonza services Ltd. (contact person: W. Kahl), June, 1990a. Report, "Spray Deposited High Performance Materials: N403", Alusuisse-Lonza services Ltd. (contact person: W. Kahl), June, 1990b. Report, "Spray Deposited High Performance Materials: N707", Alusuisse-Lonza services Ltd. (contact person: W. Kahl), June, 1990c. Reppich, B. (1993) in Plastic Deformation and Fracture of Materials, ed. Mughrabi, H., Vol 6 of Materials Science and Technology (VCH Publ. Inc., Weinheim, Germany), p. 311. Rickinson, B. A., Kirk, E A. and Davies, D. R. G. (1981) Powder Metall., 24, 1. Ruhr, M. and Baram, J. (1991) Metall. Trans., 22A, 2503. Ruhr, M., Lavemia, E. J. and Baram, J. (1990) Metall. Trans., 21A, 1785. Sampath, S., Tiwari, R., Gudmundsson, B. and Herman, H. (1991) Scripta Metall. Mater., 25, 1425. Singer, A. R. E. (1970) Met. and Mater., 4, 246. Singer, A. R. E. (1972) J. Inst. Met., 100, 185. Singer, A. R. E. (1982) Powder Metall., 25, 195. Singer, A. R. E. and Evans, R. W. (1983) Met. Technol., 10, 6. Singer, A. R. E. and Ozbek, S. (1985) Powder Metall., 28, 72. Singh, R. E, Lawley, A., Friedman, S. and Nurty, Y. V. (1991) Mater. Sci. Eng., 145A, 243. Song, G., Shen, J., Jiang Z. and Li, Q. (1994) Mater. Sci. Eng., 179/180A, 249. Srivatsan, T. S. and Lavernia, E. J. (1994) Composite Eng., 4, 459. Steen, W. M. (1985) Met. and Mater., 1, 730. Tanigawa, S., Hara, Y., Lavemia, E. J., Chin, T. S., O'Handley, R. C. and Grant, N. J. (1986) IEEE Trans. Mag., 22, 746. Turnbull, D. and Cech, R. E. (1950) J. Appl. Phys., 21, 804. Ucok, I., Ando, T. and Grant, N. J. (1991) Internat. J. Powder Metall., 27, 237. Underhill, R. E, Grant, E S., Bryant, D. J. and Cantor, B. (1995) J. Mater. Synth. Process., 3, 171. Underhill, R. U., Grant, E S. and Cantor, B. (1993) Mater. Design, 14, 45. Upadhyaya, A., Mishra, N. S. and Ojha, S. N. (1997) J. Mater. Sci., 32, 3227. Veistinen, M. K., Hara, Y., Lavemia, E. J., O'Handley, R. C. and Grant, N. J. (1987) in High Performance Permanent Magnet Materials Symposium, eds. Sankar, S. G., Herbst, J. F. and Koon, N. C. (MRS, Pittsburgh, PA), p. 93. Vetter, R., Zhuang, L. Z., Majewska-Glabus, I. and Duszezyk, J. (1990) Scripta Metall. Mater., 24, 2025. White, J., Mingard, K., Hughes, I. R. and Palmer, I. G. (1993) in Spray Forming 2, ed. Wood, J. V. (Woodhead Publ., Cambridge, UK), p. 355. White, J., Mingard, K., Hughes, I. R. and Palmer, I. G. (1994) Powder Metall., 37, 129.

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White, J., Willis, T. C., Hughes, I. R. and Jordan, R. M. (1988) in Dispersion Strengthened Aluminium Alloys, eds., Kim, Y. W. and Griffith, W. M. (TMS, Warrendale, PA), p. 693. White, J., Palmer, I. G., Hughes, I. R. and Court, S. A. (1989) in Aluminum-Lithium Alloys, eds. Sanders, Jr., T. H. and Starter, Jr., E. A. (MCE Publ., Birmingham, UK), p. 1635. Willis, J. (1988) Met. and Mater., 4, 485. Wu, M., Srivatsan, T. S., Pickens, J. R. and Lavernia, E. J. (1992) Scripta Metall. Mater., 27, 761. Wu, Y. and Lavemia, E. J. (1991) JOM, 43(8), 16. Wu, Y. and Lavernia, E. J. (1992) Metall. Trans., 23A, 2923. Wu, Y., Cassada, W. A. and Lavernia, E. J. (1993) in International Conference on Advanced Synthesis of Engineered Structural Materials, eds. Moore, J. J., Lavernia, E. J. and Froes, E H. (ASM, Materials Park, OH), p. 245. Wu, Y., Zhang, J. and Lavernia, E. J. (1994) Metall. Mater Trans., 25B, 135. Wu, Y., Cassada, W. A. and Lavemia, E. J. (1995) Metall. Mater Trans., 26A, 1235. Wu, Y., Del Castillo, L. and Lavemia, E. J. (1997) Metall. Mater Trans., 28A, 1059. Yule, A. J. and Dunkley, J. J. (1994) Atomization of Melts for Powder Production and Spray Deposition (Clarendon Press, Oxford, UK). Zeng, X., Liu, H., Chu, M. G. and Lavernia, E. J. (1992) Metall. Trans., 23A, 3394. Zeng, X., Nutt, S. R. and Lavernia, E. J. (1995) Metall. Mater Trans., 26A, 817. Zhang, J., Perez, R. J., Gupta, M. and Lavernia, E. J. (1993a) Scripta Metall. Mater., 28, 91. Zhang, J., Gungor, M. N. and Lavernia, E.J. (1993b) J. Mater Sci., 28, 1515. Zhang, J., Perez, R. J. and Lavernia, E. J. (1994) Acta Metall. Mater., 42, 395. Zhang, Y. C., Wu, Y., Zhao, Y. Y., Liu, Z. E and Chen, G. Y. (1989) inAluminium-Lithium Alloys, eds. Sanders, Jr., T. H. and Starter, Jr., E. A. (MCE Publ., Birmingham, UK), p. 95. Zhao, Y. Y., Grant, P. S. and Cantor, B. (1993a) J. Physique, 3, 1685. Zhao, Y. Y., Grant, P. S. and Cantor, B. (1993b) J. Microscopy, 169, 263. Zheng, J. H. and Kahl, W. (1990) in Advanced Aluminum and Magnesium Alloys, eds. Khan, T. and Effenberg, G., p. 317. Zhou, J., Duszczyk, J. and Korevaar, B. M. (1991) J. Mater. Sci., 26, 5275. Zhuang, L. Z., Majewska-Glabus, I., Vetter, R. and Duszczyk, J. (1990) Scripta Metall. Mater., 24, 2025.

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Chapter 8 Ion-Mixing 8.1. 8.2. 8.3.

Introduction Brief Description of Underlying Physics in Ion Mixing Thermodynamics of Alloy Phase Formation 8.3.1 Miedema's Theory and Alonso's Method 8.3.2 Interfacial Free Energy in the Multilayers 8.4. Experimentation of Ion-Mixing 8.4.1 Sample Design 8.4.2 Ion-Mixing Parameters 8.4.3 Characterization Methods 8.5. Amorphous Phase Formation 8.5.1 Glass-Forming Ability of Systems with a Negative Heat of Formation 8.5.2 Glass-Forming Ability of Systems with a Positive Heat of Formation 8.5.2.1 Amorphous Alloys Formed Within Restricted Compositions 8.5.2.2 Amorphous Alloys Formed in a Broad Composition Range 8.5.2.3 Nominal and Intrinsic Glass-Forming Ability 8.6. Formation of Metastable Crystalline Alloys 8.6.1 Structural Classification of the Metastable Crystalline Phases 8.6.1.1 Solid Solutions 8.6.1.2 h.c.p.-I and f.c.c.-I Phases 8.6.1.3 h.c.p.-II and f.c.c.-II Phases Based on b.c.c. Metals 8.6.2 Free Energy calculation of the MX phases 8.7. Interface-Generated Solid-State Vitrification in Systems with a Positive Heat of Formation 8.8. Concluding Remarks Acknowledgments References

197 199 200 200 201 203 203 205 205 205 206 207 207 207 211 213 213 213 213 215 216 218 219 220 220

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Chapter 8 Ion-Mixing B. X. LIU

8.1. I N T R O D U C T I O N

The earliest use of an energetic ion beam generated from low-energy accelerators (several 100-keV or higher) to modify material properties was initiated in early 1960s, i.e. the doping of semiconductors by implanting the desired ion species into silicon wafers (Carter and Grant 1976). Since then, ion implantation has significantly developed and has become a supporting technique in the semiconductor and VLSI technology (Dearnaley et al. 1973). Since early 1970s, ion-implantation has also been employed to deal with metallic materials, e.g., for improving the wear and corrosion resistance, catalytic properties, superconductivity, etc. (Mayer and Lau 1988). As a versatile metal ion source has not been available for some time, it has not been possible to implant various metal species to achieve a high ion concentration frequently required for modifying the surface properties of the metallic materials. An alternative scheme, namely ion-mixing of multiple metal layers, was therefore introduced and a new research area of synthesizing metastable alloys by ion-mixing has boomed since the early 1980s (Liu et al. 1983). Figure 8.1 shows schematically how ion-mixing is realized in multilayered films consisting of metal layers A and B. The A-B multilayered films are prepared by depositing alternately the metals A and B on a substrate. The as-deposited films are then irradiated by a scanning ion beam, e.g., 300 keV xenon ions, to a certain dose to achieve a uniform mixing between metal layers A and B, resulting in forming an Al-xBx alloy. The substrate can be SiO2 or NaC1 acting as a supporting substance or it can be metal A as a matrix for forming a surface Al-xBx alloy coating on metal A. It is known that the first amorphous alloy, also called metallic glass, was obtained by liquid melt quenching (LMQ) in 1960 (Klement et al. 1960). Because amorphous alloys have a non-crystalline structure (with no vulnerable grain boundaries), they are expected to possess unique properties. During the past 30 years, various techniques have been developed to synthesize metallic glasses. From a physical point of view, high cooling rates of 106 K/s or higher are needed to avoid crystallization against amorphization (Davies 1983, Hafner 1981, Giessen and Whang 1980) and cooling rates of up to 107 K/s have been achieved so far. Such developments lead to expand the number and types of metallic glasses, including a variety of metal-metalloid and metal-metal amorphous alloys (Takayama 1976). Laser-quenching has also been applied for producing amorphous alloys (von Allmen 1983, Lin and Spaepen 1984), as many intense laser sources have 197

198

B. X. Liu

amorph~s SiOz

... ..

: NaCI 3,.keVXeions

:~~7

[

cn'statline Figure 8.1. Ion-mixingof multilayered films consisting of metal layers A and B.

become available and they can achieve a cooling rate of 1012 K]s in the material's surface layer (See the chapter on "Laser Processing" by K.E Kobayashi in this book). In comparison, effective cooling rates can reach as high as 1014 K]s (Thompson 1969) in ion-mixing, and therefore this technique can also produce a variety of amorphous alloys. In fact, the cooling rates achieved during ion-mixing are probably the highest among the currently available glass producing techniques. Up-to-now, numerous metastable alloys with either an amorphous or crystalline structure have been obtained by ion-mixing in some 100 binary metal systems (Liu and Jin 1997), and in many alloy systems that are not obtainable by other non-equilibrium processing techniques. The glass-forming ability (GFA) of a system can be considerably improved by ion-mixing because of its high cooling rate. Furthermore, ion-mixing offers a unique opportunity for studying the amorphization mechanism, since the ion dose can be varied at small intervals and so the process of ion-mixing can be studied step-by-step. Based on experimental data, several empirical models have been proposed to predict the GFA and the glass-forming composition range (abbreviated hereafter as GFR) of the binary metal systems. Liu et al. (1983) proposed the first rule, namely the structural difference rule, which predicted a sufficient condition and demonstrated the predominant role of the crystalline structures of the component metals in forming metallic glasses by ionmixing. Further studies led Liu to formulate a two-parameter model. The first parameter is the maximum possible amorphization range (MPAR), which was defined as the total width of the two-phase region, equaling the whole composition range (100%) minus the maximum solid solubilities, observed from the corresponding equilibrium phase diagram (Liu et al. 1987). The second parameter is the heat of formation A H f calculated by Miedema's model (Miedema 1976). Based on a quantitative combination of the parameters, the binary metal systems are classified into readily, possibly, and hardly glass-forming ones (Liu et al. 1987). Alonso and Simozar (1983) constructed a two-dimensional map with

199

Ion-Mixing

the ratio of atomic radii, rA/rB, and the A H f and by summarizing the published data, they separated the binary metal systems into glass-forming and non-glass-forming ones and predicted that no amorphous alloy could be formed in a system with a A H f exceeding + 10 kJ/mol. Ossi (1990) proposed an atomistic model based on collision cascades and the related bombardment-induced surface compositional changes. Though the prediction by the two-parameter model matched best with the experimental data, lack of a relevant thermodynamic interpretation was the main drawback until the early 1990s. It was also noted that metallic glasses were unexpectedly obtained in some alloy systems with a positive A Hu, which have long been recognized as non-glass-forming ones (Liu 1987, Jin et al. 1995, Zhang et al. 1995). The thermodynamic stability of the different alloy phases was evaluated by Miedema, who developed a semi-quantitative model in the mid 1970s to calculate the free energy curves of the alloy phases versus alloy compositions by using the available characteristic parameters of the constituent metals (de Boer et al. 1989), rather than the parameters of the alloy phases, as required in CALPHAD calculation. This chapter attempts to present a brief review of the up-to-date progress on the synthesis of metastable alloys by ion-mixing of multilayered films. It is divided into the following sections: a brief description of the underlying physics in ion-mixing, thermodynamic calculations based on Miedema's model, experimentation and sample design, formation of metallic glasses in systems with a positive or a negative heat of formation, metastable crystalline alloys synthesized by ion mixing, and applicability of the interface concept to solid-state amorphization.

8.2. B R I E F D E S C R I P T I O N O F U N D E R L Y I N G P H Y S I C S IN I O N M I X I N G

This section briefly describes the underlying physics in ion-mixing and the experimental details of ion-mixing will be discussed in Section 8.4. The ion-mixing process is commonly divided into two steps, viz., an atomic collision cascade and subsequent relaxation. After relaxation, one can consider having a delayed step and it has no significant effect on alloy phase formation. When an ion with an energy E1 impacts on the material surface and penetrates inside, it collides with the atoms in the material. If an atom in the material receives an energy E2 exceeding a threshold energy Ed, the atom will leave its own lattice site and a vacancy is created. This displaced atom is referred to as knock-out atom and it will, if E2 is greater than Ed, trigger secondary collisions, then tertiary ones, and so on, and this is named an atomic collision cascade. To describe atomic displacements and defect creation caused by ion irradiation, a parameter "dpa", (abbreviation for displacement p e r atom) is defined, which can be calculated through atomic collision theory, as

do-

E2 (max)

d p a -- Nd = ~ . Crc

No

rJed

v(E2)

9dE2 ~

(8 1) "

whereNd and No are the numbers of displaced and total atoms per volume, respectively; r is the ion flux, ere is the collision cross-section; Ed and E2 are the displacement threshold energy and the energy received by the knock-out atom, respectively; v(E2) is the number

200

B. X. Liu

of displaced atoms created by the knock-out atom with an energy of E2, and dot/dE2 is the differential collision cross-section at E2. The atomic collision cascades trigger intermixing of the metal layers A and B in the multilayered films. As the energy of the ions is on the order of several 100 keV and the binding energy of the solids is typically 5-10 eV, atomic collision is a highly dynamic process of far-from equilibrium. By receiving an adequate irradiation dose, the layered structure of the multilayered films is smeared out and metals A and B mix into a uniform mixture, which is in a highly energetic state with a disordered atomic configuration. At the termination of the cascade, the highly energetic mixture should somehow relax towards equilibrium, and at this stage, the equilibrium thermodynamics comes into play in governing the direction of relaxation. Whether the mixture can reach the equilibrium state or not, however, depends on the temperature and time conditions available during this period. During the atomic collision cascades, all the atoms are involved in violent motion, it is hard to imagine the formation of any kind of alloy phase. Consequently, it is commonly considered that the structure of the alloy phase is to be formed during the relaxation period. As the relaxation period is very short (10-1~ -9 seconds, as we will discuss soon) and the mixture cannot straightforwardly go to the equilibrium state in most cases, it frequently resides in some intermediate states, corresponding to the metastable states of either an amorphous or a simple-structured crystalline phase. From an energy point of view, a thermal spike concept was introduced to discuss the ion-mixing process. Suppose a target is irradiated by ions of 300 keV energy and a current density of 1-4/zA/cm 2 for a certain time period, a huge amount of energy would be deposited into the material surface within a very thin layer. Estimation showed that in a local region, e.g., using a cylinder geometry, with a diameter of 5 nm and a length of 100 nm, the temperature could be as high as 103-104 K, while the surrounding material remains close to the ambient temperature of 300 K. According to thermal conduction calculations, the temperature drops to 300 K within approximately 10-1~ -9 second, which is the duration of relaxation. Apparently, during relaxation the effective cooling rate can be as high as 1014 K/s (Thompson 1969).

8.3. T H E R M O D Y N A M I C S

OF A L L O Y P H A S E F O R M A T I O N

As described above, the process of ion-mixing, as a whole, takes place under conditions that are far from equilibrium. However, during the relaxation period, the highly energetic state relaxes towards equilibrium. At this stage, equilibrium thermodynamics plays a role in governing the possible evolution path, thus determining the nature (amorphous or crystalline) of the alloy phase formed. It is therefore necessary to calculate the free energy of the competing phases in the alloy system to understand and predict alloy phase formation upon ion-mixing.

8.3.1 Miedema's theory and Alonso's method In this section, the discussion is focused on the amorphous phase formation and the thermodynamics for metastable crystalline phase formation will be discussed later.

201

Ion-Mixing

Generally, the Gibbs free energy A G of a phase is calculated by A G = A H -- T A S, where A H and AS are the enthalpy and entropy of formation, respectively. The entropy usually plays only a secondary role in solid phase formation, and its expression can be taken, as a first approximation, as that of an ideal solid solution, i.e., AS = - - R [ x A ln(xa) + XB ln(XB)], where R is the gas constant, and XA and XB are the atomic concentrations of metals A and B, respectively. For a substitutional solid solution of transition metals, the enthalpy of formation, according to Miedema's model and Alonso's method (Alonso et al. 1990, Niessen and Miedema 1983, Lopez and Alonso 1985), can be considered as a sum of three contributions A H = A H c + A H e + A H s,

(8.2)

where A H c, A H e and A H s correspond to the chemical, elastic, and structural contributions, respectively, and they can be calculated according to the well documented theory and methods. For an amorphous phase, the enthalpy of formation due to both elastic and structural terms is absent. Following the suggestion of Niessen et al. (1988), the enthalpy of the amorphous phase is then given by

(8.3)

AHamorphous-- A H c @OI(XATm,A @ XBTm,B),

where ot is an empirical constant equal to 3.5 J/(mol 9K) and Tm,i is the melting point of the component i. 8.3.2 I n t e r f a c i a l f r e e energy in the multilayers

Since ion-mixing (IM) begins with multilayered films, which definitely consist of a number of interfaces, the effect of interfacial free energy on the alloying behavior should be taken into account. To include the effect of interfaces, the Gibbs free energy of the initial state of the A-B multilayered films was calculated by adding the interfacial free energy to the ground state of a mixture of A and B in bulk form. As pointed out by Turnbull (1955), the interfacial free energy between two solid phases can be split into geometrical and chemical contributions and can be written as ysAB m

.geo .chem r s s + Yss 9

(8.4) geo For two solids with quite different lattice structures, the geometrical term Yss stands for the energy of a large-angle grain boundary. It is estimated that the grain boundary enthalpy is about 30 percent of the surface energy at 0 K (Murr 1975) and the entropy contribution is neglected. Thus .

.chem , y~sAsB - - 0 . 1 5 x (ys,o + YBB,0) + YSS

(8.5)

where yjS,0, ys,0 are the surface energies of metals A and B, respectively, and Yss'chem is given by Miedema (1978)

B. X. Liu

202 ----d

---'B

--dA'-~ ~ d B - ~ d A

B

A

I I I I I I

iI I I I I I I

A

Figure 8.2. Schematicconfiguration of A-B multilayered films.

ychem

SS

i12/3

-- AnOinB/(CO, a ),

(8.6)

where Co is a constant being 4.5 • 108, A H ~ in B is the heat of solution of metal A in metal B at infinite dilution (Miedema 1978) and VA is the atomic volume of metal A. According to Gerkema and Miedema (1983), the interfacial free energy of the multilayers, AGmul was calculated as follows"

A Gmul - S f • ysAsB,

(8.7)

where Sf is the surface area occupied by one mole of interfacial atoms. Figure 8.2 shows schematically the configuration of the A-B multilayers. Let us assume that NA and N8 are the numbers of layers of metals A and B, dA,i and dB,i are the thicknesses of the ith layers of metals A and B, respectively; nA and n8 are numbers of interfaces (A on B) and (B on A); AdA,i and AdB,i are the thicknesses of the ith interfaces (A on B) and (B on A), respectively. Obviously, the relationship of n A+ n 8 + 1 = NA +N8 always holds for these four parameters. The interfacial free energy of the multilayers can then be calculated as

(8.8)

AGmul -- OtASf AYBA + OIBSfBYAB,

where SfA and SfB are the surface areas occupied by one mole of A and B atoms, respectively, lYa and OrB are the ratios of interfacial atoms A and B to the total number of atoms in the multilayered films, respectively, and are given by (Zhang and Liu 1995)

S nA AdA,i OlA -- XA Z NA dA,i

Z j B mdB,j ,

or8 - x 8

~NB NB,j

,

(8.9)

Z ; A AriA, i / ~-~NA dA,i and )--~.8 A d s , j / ~--~8 ds,j are the ratios of interfacial A atoms to the total number of A atoms and the ratio of interfacial B atoms to the total number of B atoms, respectively. These terms, multiplied by XA and x s, respectively, give the fractions of A and B interfacial atoms versus the total number of (A + B) atoms where

Ion-Mixing

203

in the multilayers. According to experimental studies on thin films (Clemens 1986) and for simplicity of calculation, the thickness of each interface is assumed to have the same value of 0.5 nm. It should be pointed out that in experimental studies one can always add a few extra interfaces in a sample design to ensure a higher energetic state of the multilayers than that of the phase of interest, so that the assumption of the effect of this thickness can be ignored. Figure 8.3 shows two free energy diagrams calculated by the above methods for two representative systems with negative (Ni-Nb) and positive (Y-Nb) AHf to show a significant difference due to a different sign of the heat of formation. One sees that the free energy curve of the amorphous phase for the Ni-Nb system is concave, i.e., it is always lower than that of an equilibrium mixture of two crystalline metals, suggesting a favored condition for forming glass. The free energy curve of the amorphous phase in the Y-Nb system is convex and is always higher than that of the equilibrium state without considering the interfacial free energy. It is noted that taking into account the interfacial free energy, the energetic sequence of the respective states in the Y-Nb system is changed significantly, while the effect of interface is negligible for the Ni-Nb system. The details of this will be further discussed together with the experimental results later.

8.4. E X P E R I M E N T A T I O N

OF I O N - M I X I N G

8.4.1 Sample design In designing the multilayered samples, the total thickness of the multilayers should match with the range of the irradiating ions. Inert gas ions, such as Xe, Kr, Ne, or Ar, are frequently chosen to be the irradiating ions. The advantages of employing inert gas ions are that the chemical impurity effect is negligible, if the ion dose is below 1016 ion/cm 2, and that the gas ion sources are frequently available easily for the common implanters. As heavier ions induce more efficient mixing than the lighter ones when the other conditions are kept the same, Xe + ions are therefore mostly used. The typical energy of the irradiating ions is several 100 keV. It was found experimentally that if the total thickness of the multilayered film was designed to be the projected range Rp plus the projected range straggling ARp of the irradiating ions, uniform mixing between the layers could be achieved well at an ion dose range of 1014-1016 ion/cm 2. The Rp and ARp for elemental metals can be calculated easily by the LSS theory (Dearnaley et al. 1973). For an alloy target of AxB 100-x, where x refers to the atomic percentage of metal A, the projected range of the alloy can be calculated by the following equation (Townsend et al. 1976):

Rp (Ax B 100-x ) --

Rp (A) Rp (B) yRp(B) + ( 1 0 0 - x)Rp(A)'

(8.10)

where y refers to the weight percentage of metal B in the alloy and can be obtained from x easily. The calculation of the projected range straggling ARp for the alloys is rather complicated. However, considering the errors involved in the calculation itself, as well as in controlling the thickness during deposition, the above equation is also good enough to estimate ARp of the alloy, by substituting ARp and Rp. The total thickness is divided into the thicknesses of metals A and B, according to the desired atomic ratio of the alloy being

204

B. X. Liu

Nb

Y ,

,

(a)

.

,

Muitilayers(17 interfaces)

+60

Amorphous +40

+20

0

&

+20

&

I

9

Multilayers(7 interfaces) (17 interfaces) --"-'--'"'"'""-7~.~ --

-20

----7,

Amorphous

(b) Ni

I

20

9

I

1

40 60 80 Nb concentration (at. %)

J

Nb

Figure 8.3. Calculated free energy diagrams of the Y-Nb (a) and Ni-Nb (b) systems showing the different effects of interfacial free energy on the energetic sequence in the positive and negative heat of formation systems.

studied. Typically, the total thickness of the films of transition metal alloys is about 60100 nm under 300 keV xenon ion-mixing. Another parameter is the fraction of interfacial atoms, which determines the number of metal layers. Each individual layer is always designed to be 5-8 nm or even thinner. The multilayered films are prepared in a vacuum system to deposit alternately the metals A and B onto an inert substrate supporting the film or on a metal matrix. The supporting materials can be SiO2, A1203 or NaC1 single crystals. The system is often a multi-crucible electron gun evaporation system with a vacuum level better than 10 -5 Pa.

Ion-Mixing

205

8.4.2 lon-mixing parameters It was found experimentally that the ion dose required to homogenize the discrete layered films is in the range of 1014 to 1016 ions/cm 2 and varies from system to system. The ionbeam current density should be controlled generally around 1/~A/cm 2, or at most up to 4/zA/cm 2 to avoid overheating of the multilayers during ion irradiation. In operating the implanter, the ion dose, which is measured by the electrical charge, is easy to adjust and control. Ion-mixing thus offers a unique method for investigating the phase transformation process step-by-step. Another important parameter in ion-mixing is the temperature at which the experiment is conducted. As ion-mixing often forms metastable alloy phases, ion-irradiation is generally performed at low temperatures, e.g., at room temperature or down to liquid nitrogen temperature. Nevertheless, the temperature mentioned for the experiments is often the nominal temperature measured at the sample holder instead of the effective temperature of the irradiated films. The effective temperature is an important parameter in the process induced by ion irradiation, and it should be considered when discussing the alloy phase formation mechanism. The stable temperature of the expecting alloy should be considered, i.e., if one expects to form an amorphous alloy with a recrystallization temperature Ta-c (from amorphous-to-crystalline), the effective temperature should of course be kept at a temperature far below Ta-c.

8.4.3 Characterization methods 1-2 MeV He + ion Rutherford backscattering spectrometry (RBS) (Chu et al. 1978) is frequently used to analyze the amount of mixing in the multilayered films. Figure 8.4 shows a typical example for Fe-Mo multilayered films before and after ion-mixing and one can clearly see the information given by the RBS spectra. Besides, XPS, SIMS and Auger spectroscopy are employed to analyze the composition depth profile of the ion mixed films. The structure of the phase is identified mainly by X-ray diffraction and TEM. Fourpoint probe resistivity measurement is a supplemental method to give evidence that the amorphous phase is formed by a crystal-to-amorphous transformation, because the resistivity of the amorphous alloys is often considerably greater than that of their crystalline counterparts.

8.5. A M O R P H O U S

PHASE FORMATION

The maximum possible amorphization range (MPAR) is a decisive parameter in determining the GFA or GFR. When a system has an MPAR exceeding 20%, the amorphous phase can probably be formed in the system. If the MPAR is less than 20%, the related two-phase region is too narrow and can be further decreased by the extension of the solid solubilities from two metal ends under ion-mixing. In other words, it is difficult to form an amorphous alloy (Liu 1987). For an extreme value of MPAR = 0, the phase diagram of the system is of continuous solid solution type with no two-phase region, and GFA and GFR are zero. The following discussion will show that the empirical prediction agrees

206

B. X. Liu

~.~

Mo

2.0MeV

SiO2 ~ ~~70.

He

Fe

+

Mo

.,4 ~a 4~

Ix 1016Xe,./ c m

2

\

as-deposited'

|

. ~ _

500

i_ =.........

600

Channel. Number Figure 8.4. Backscattering spectra of Fe-Mo multilayered films before (solid line) and after room temperature ion-mixing to a dose of 1 • 1016 Xe+/cm 2.

reasonably well with the experimental results for the systems featuring either negative or positive A H f 8.5.1 Glass-forming ability of systems with a negative heat of formation A large body of data confirmed that an amorphous alloy is most likely to form by IM in the two-phase region (Liu et al. 1983), which can be attributed to the competition of crystallization between two crystalline phases, resulting in a frustration of nucleation and growth of the crystalline phase. It is also easy to form amorphous alloys at or in the vicinity of intermetallic compound compositions. This is even more so when the intermetallics have complex crystal structures and therefore cannot form due to a very restricted kinetic condition available in IM (Liu and Jin 1997). Consider the Ni-Nb system (A H i = - 4 2 kJ/mol) as a representative example, whose calculated free energy diagram is shown in Figure 8.3. One sees clearly that the energetic sequence of the phases involved is thermodynamically favorable for the formation of a glassy phase. As the formation of intermetallic compounds is usually hindered kinetically, they need not be considered when discussing the GFA. The GFR was therefore regarded as the composition range in which the free energy of the amorphous phase is lower than

Ion-Mixing

207

that of the solid solutions. The points of intersection between the free energy curve of the amorphous phase with those of the two solid solutions are around 20 and 85 at. % Ni. The GFR in Ni-Nb is therefore from 20 to 85 at.% Ni, which is in excellent agreement with the IM results, i.e. 20-85 at.% Ni, observed by Zhang et al. (1993), as well as very close to the defined MPAR of the system.

8.5.2 Glass-forming ability of systems with a positive heat of formation For the Y-Nb system with a positive A H i, shown in Figure 8.3, the free energy curve of the amorphous phase is usually higher than that of an equilibrium mixture of crystalline metals A and B in bulk form, leading to a situation unfavorable for amorphization. However, as the interfaces in multilayered films are generally regarded as transient layers of mixed A and B atoms in a highly disordered configuration, they contribute extra free energy to the films. With an increasing fraction of the interfacial atoms, it is possible to have the free energy curve of the multilayers intersect with that of the amorphous phase in some specific composition ranges, making it possible to form a glass. These composition ranges depend on the designed fraction of the interfacial atoms of the films (Zhang et al. 1995). Some examples of experimental results are given in the following paragraphs.

8.5.2.1 Amorphous alloys formed within restricted compositions Figure 8.5 shows the calculated free energy diagram of the Ag-Mo system with a A Hf of +56 kJ/mol, the second largest A H f for transition metal binary systems. One sees that the free energy curve of the multilayered films containing 12 layers (or 11 interfaces) intersects with that of the amorphous phase, dividing the diagram into three regions along the composition axis. Obviously, in regions I and III, the amorphous phase is possible to be formed, while in region II, amorphization is hardly to be achieved. Ion-mixing by 200 keV xenon ions at liquid nitrogen temperature did result in the formation of an amorphous alloy at near 80 at.% Mo, and the formation of another amorphous phase co-existed with a new Mo-based f.c.c, phase around 40 at.% Mo (Jin et al. 1995). Such compositional discontinuity in glass formation has also been observed in other systems with a positive A H f (Jin and Liu 1994).

8.5.2.2 Amorphous alloys formed in a broad composition range With a further increase in the fraction of interfacial atoms versus the total number of atoms in the multilayered films, the two separated composition ranges favoring amorphization may become broader and meet each other to become a continuous one. The calculated free energy diagram of the Y-Mo system with A H f = +35 kJ/mol, shown in Figure 8.6, includes three free energy curves for the multilayered films with 9, 11 and 19 layers, corresponding to the fractions of interfacial atoms of 8.1%, 10.0% and 18.1%, respectively. Table 8.1 lists the predicted composition ranges favoring amorphization for these three cases. The Y-Mo multilayered films were prepared accordingly for ion-mixing investigation and the experimental results are in excellent agreement with the prediction (Zhang et al. 1995). These results suggest that through appropriately designing the multilayers, it is possible to synthesize metallic glasses in systems with positive A Hf even with a composition near the equiatomic stoichiometry, where the amorphous phase frequently has the

208

B. X. Liu

_~-

O

6O

>'

40

"

i

~

I

'

I

'

I

E

30 ii

~

2o

~

10

L_

,

f.c

0

12 L a y e r s

)

(.9

I

~

v

K,.,

'

-

rphous

1

l

Mo 10

1

,

1

30

50

9

1

70

j.

~'~

I

90

Ag

Ag C o n c e n t r a t i o n Figure 8.5. Calculated free energy diagram of the A g - M o system.

Table 8.1 Predicted and experimental results of amorphous phase formation in the Y-Mo multilayered films upon 200 keV xenon ion mixing at room temperature, n" number of interfaces in the multilayers; o~" fraction of interfacial atoms in the multilayers. Composition

Y80 Mo20

Y60 Mo40

Y50 Mo50

Y40 Mo60

Y20 Mo80 Yes

n = 18

Predicted

Yes

Yes

Yes

Yes

(a = 18.1%)

Ion mixing

Yes

Yes

Yes

Yes

*

n = 10

Predicted

Yes

Yes

No

No

Yes

(or = 10%)

Ion mixing

Yes

Yes

No

No

Yes

n =8

Predicted

Yes

No

No

No

Yes

ot = 8.1%)

Ion mixing

Yes

No

No

No

Yes

* Experiment not conducted.

highest free energy and is most difficult to form. It is worth mentioning that the IM results listed in Table 8.1 include both "yes" and "no" and that they strongly depend on the correlation between interfacial fraction and alloy composition. These results also support the argument that IM induced amorphization can really be attributed to the interfacial free energy, but not to the irradiation energy. The irradiation energy has exactly the same effect on all the multilayered films and does not depend on alloy composition. In the past 15 years, nearly 100 binary metal-metal systems, according to published results, have been studied by ion-mixing of multilayers. Table 8.2 lists these systems in alphabetical order of the chemical symbols of the constituent metals for convenience of searching.

Ion-Mixing

209

Table 8.2 Binary metal-metal systems that have so far been studied by ion mixing AHf

Any

<0

System

Ref.

Remarks

A1-Au A1-Co A1-Fe A1-Mn A1-Mo A1-Nb A1-Ni A1-Pd A1-Pt A1-Ta A1-Ti A1-Y A1-Yb Au-Ta Au-Ti Au-V Co--Gd Co-Hf Co-Mo Co-Nb Co--Ta Co-Tb Co-Ti Co-Y Co-Yb Co-Zr Cu-Y Cu-Zr Er-Ni Fe-Gd Fe-Mo Fe-Nb Fe-Ta Fe-Tb Fe-Ti Fe-Y Fe-Zr Hf-Ni Mo-Ni Mo-Ru Mo-Zr Nb-Ni Ni-Ta Ni-Ti Ni-Y Ni-Yb Ni-Zr Ru-Ti Ru-Zr

Schmid and Ziemann (1985) Hung and Mayer (1985) Rauschenbach and Hohmuth (1987) Knapp and Follstaedt (1987) Meissner et aL (1987) Liu et al. (1983) Hung et al. (1983) Hung et al. (1983) Hung et al. (1983) Meissner et al. (1987) Cheng (1986) Liu et al. (1991) Ding et al. (1992) Pan et aL (1996) Liu et al. (1982) Tsaur et aL (1991) Yan et aL (1986) Anderson et al. (1990) Liu et al. (1983) Zhang and Liu (1994) Zhang and Liu (1994) Yan et al. (1987) Hung and Mayer (1985) Liu et al. (1991) Liu et aL (1991) BCttiger et al. (1987) Liu et aL (1991) Liu et al. (1991) Liu et al. (1983) Yan et al. (1987) Liu et aL (1987) Zhang and Liu (1994) Zhang (1995) Yan et al. (1987) Brenier et al. (1987) Liu et al. (1991) BCttiger et al. (1987) van Rossum et al. (1983) Liu et al. (1983) Liu et al. (1983) Jin and Liu (1995) Zhang et al. (1993) Liu and Zhang (1994) Cheng et aL (1984) Liu et al. (1991) Liu et al. (1991) BCttiger et al. (1989) Cheng et al. (1984) Liu (1985)

Several glasses have been formed in many systems. Or, at least one glass has been formed in some systems.

B. X. Liu

210 Table 8.2 Continued AHf

System

Ref.

Remarks

Any < 0

Au-Cu Cu-Rh Mo-Nb Fe-W

Rauschenbach (1986) Peiner and Kopitzki (1988) Liu (1982) Goltz (1982)

No glass has so far been formed.

Cu-Ir Ti-Zr

Peiner and Kopitzki (1988) Liu (1982)

Ag-Cr Ag--Cu Ag-Mo Ag-Nb Ag-Ni Ag-Ta Ag-V Au-Co Au-Ir Au-Mo Au-Os Au-Ru Au-W Co-Cu Cu-Fe Cu-Mo Cu-Nb Cu-Ta Fe-Nd Hf-Nb Hf-Ta Mo-Y Nb-Ti Nb-Zr Ta-Ti Ta-Y Ta-Zr Ti-Y Y-Zr

Liu et al. (1987) Liu et al. (1986) Jin et al. (1985) Jin and Liu (1996) Liu et al. (1986) Andersen et al. (1990) Andersen et al. (1990) Tsaus et al. (1981 ) Peiner and Kopitzki (1988) Pan and Liu (1995) Peiner and Kopitzki (1988) Meissner et al. (1987) Meissner et al. (1987) Liu et al. (1987) Huang and Liu (1987) Chen and Liu (1997) Andersen et al. (1990) Tsaur (1980) Yan et al. (1987) Jin et al. (1995) Liu et al. (1996) Zhang et al. (1995) Chen and Liu (1988) Jin and Liu (1994) Chen et al. ( 1991 ) Zhang and Liu (1995) Jin and Liu (1997) Liu et al. (1991 ) Liu et al. (1998)

Several glasses have been formed in many systems. Or, at least one glass has been formed in some systems.

Ag-Co Ag-Fe Ag-Ir Ag-Os Ag-Rh Ag-Ru Ag-W Au-Fe Au-Ni Au-Rh Cr-Cu Cu-Os Fe-Yb

Liu et al. (1998) Andersen et al. (1990) Peiner and Kopitzki (1988) Peiner and Kopitzki (1988) Peiner and Kopitzki (1988) Peiner and Kopitzki (1988) Hiller et al. (1989) Tsaur et al. (1981 ) Tsaur et al. (1981) Peiner and Kopitzki (1988) Shin et al. (1983) Meissner et al. (1987) Liu et al. (1991)

No glass has has so far been formed.

AHf = 0

AHf = 0

AHf

> 0

AHf>O

Several glasses have been formed. No glass has so far been formed.

Ion-Mixing

211

18.1%

+40

*20

10.0% 8 1%

1

0 Bulk L..

,I,,

Y

20 Mo

State 9

40

J

,

60

Concentration

,,

&

80

Mo

(at%)

Figure 8.6.

Calculated free energy diagram of the Y-Mo system. The numbers of 8.1%, 10.0% and 18.1% are for the fractions of interfacial atoms versus the total atoms in the multilayers, respectively.

8.5.2.3 Nominal and intrinsic glass-forming ability From both practical and theoretical points of view, the GFA and its prediction are of vital importance. The originally defined GFA was only to classify the systems into two categories of glass-forming and non-glass-forming ones in a qualitative way and the classification mostly depended on glass preparation technique employed. However, by employing powerful techniques, e.g., IM, many new metallic glasses have been obtained in those systems considered as non-glass-forming ones by LMQ. Apparently, the definition or the physical meaning of GFA needs to be extended. Firstly, GFA of a system should reflect qualitatively whether a metallic glass can be formed or not in the system. As mentioned above, the GFA determined by more powerful techniques frequently possessing higher cooling rates is expected to be closer to the real ability of the system. In other words, from a physical point of view, there exists an intrinsic GFA reflecting the capability of the system itself to form a metallic glass and it does not depend on the technique employed to produce the glass. Furthermore, it is necessary to have a quantitative measure of GFA to compare the capabilities of glass formation among the glass-forming systems. The experimentally determined GFR, in the author's view, can be taken as a quantitative measure of GFA. Apparently, the larger the GFR, the greater the GFA. It is true that if one employs a technique with a higher cooling rate, one may obtain the glassy phase in a broader composition range than that observed by a technique with a lower cooling rate. In this sense, the experimentally observed GFR can only be considered as a nominal GFA. Obviously, the greater the GFR observed by a specific technique, the closer the GFR is to the intrinsic GFA. Consequently, to approach the intrinsic GFA of a system, one should

B. X. Liu

212

Table 8.3 Some examples of metallic glasses formed in binary metal systems with a negative heat of formation by ion-mixing No.

System

AHf (kJ/mol)

MPAR (%)

Glass-forming range

1 2 3 4 5

Zr-Ru Ni-Nb Co-Nb Nb-Fe Co-Mo

-86 -44 - 38 -23 -7

90 86 86 85 76

25-75 at.% Ru 20-85 at.% Nb 20-77 at.% Nb 20-78 at.% Fe 35-77 at.% Mo

Table 8.4 Some examples of metallic glasses formed in binary metal systems with a positive heat of formation by ion-mixing. (Total thickness of the multilayered films is 40-50 nm.) No.

System

AHf (kJ/mol)

MPAR (%)

Number of interfaces

Glass-forming range

1 2 3 4 5

Zr-Nb Ag-Mo Ti-Ta Y-Nb Y-Mo

+6 +56 +3 +44 +35

70 100 100 100 100

7 11 15 17 18

at 12 or 60 at. % Nb at 30 or 80 at.% Mo 74-89 at.% Ta 28-44 at.% Nb 20-80 at.% Mo

look for a technique capable of generating cooling rates as high as possible. At present, IM can be considered as the technique capable of revealing the m a x i m u m GFA a m o n g the currently available glass producing techniques. The next issue is how to correlate the n o m i n a l G F R observed by IM and the intrinsic GFA of the system. Based on IM study, a parameter, namely MPAR, has been proposed by the author to predict the m a x i m u m possible G F R of a system. According to the definition, M P A R = 1 0 0 % - ( o r + 13), where c~ and/~ stand for the m a x i m u m solid solubilities at the two metal ends (Liu et al. 1987, Liu 1987). As IM can frequently form some supersaturated solid solutions, which would actually reduce the total width of the two-phase regions (Liu 1986), an empirical formula is therefore suggested for correlating the possible G F R and MPAR, i.e. G F R = 0.7MPAR, and it is reasonably compatible with the IM results in either positive or negative A H f systems, as shown by those examples listed in Tables 8.3 and 8.4. W h e n MPAR=0, corresponding to a system with a continuous solid solution phase diagram, G F R = 0 and no glass has so far been obtained in such systems, which are therefore considered as non-glass-forming ones.

213

Ion-Mixing

8.6. FORMATION OF METASTABLE CRYSTALLINE ALLOYS

8.6.1 Structural classification of the metastable crystalline phases Up-to-now, a number of metastable crystalline phases (MX) have been formed by ionmixing in both negative and positive A H i systems and they are classified into 5 structural categories (Liu et al. 1995), i.e., supersaturated solid solutions, the h.c.p.-I phases in h.c.p.- or f.c.c.-based alloys, the f.c.c.-I phases in h.c.p.-based alloys, the h.c.p.-II and f.c.c.-II phases in b.c.c.-based alloys. 8.6.1.1 Solid solutions

Ion-mixing of multilayered films can extend the solid solubility of the binary metal alloys resulting in the formation of a supersaturated solid solution always having the same simple crystal structure as the major alloying metal. Most of these alloy phases are identified as solid solutions. Their lattice constants agreed quite well with those predicted by Vegard's law, and they are therefore substitutional solid solutions (Liu and Jin 1987).

8.6.1.2 h.e.p.-I and f.e.e.-I phases A metastable crystalline phase with a similar h.c.p, structure has been formed by ionmixing in five binary (A-B) metal systems (Co-Au, Ti-Au, Co-Mo, Ni-Mo, and Ni-Nb) (see Table 8.5) where A refers to the first entry in the alloy designation. The MX phases were formed in A-rich multilayered films with an overall composition near A3B. These MX phases were named as h.c.p.-I, since they have all about the same spacing (dcpp) of the close packed planes (Table 8.5). Their formation was attributed to the valence electron effect. For a close-packed hexagonal structure, the maximum number of electron states per atom, n, in the Jones zone (including doubling for spin) can be calculated by the following equation (Jones 1960):

n=2-~

(a)2[ 1(a)2] c

1-~

c

'

(8.11)

where a and c are the lattice constants of the h.c.p, structure. According to this calculation, the n values of these five phases are almost identical, i.e., n -- 1.73 to 1.74, which is very close to the value of 7/4 = 1.75, corresponding to the well-defined Hume-Rothery 7/4 electron compound. These A3B alloy phases can therefore be considered as metastable electron compounds (Liu 1983). Table 8.6 shows that MX f.c.c.-I phases have formed in some h.c.p.-metal-ricia multilayered films by ion-mixing. As both h.c.p, and f.c.c, structures are built up by the stacking of close-packed atomic planes differing only slightly in the stacking order, one can easily transform into another. For example, sliding of atoms on the planes parallel to the (0002)hcp plane along (1 i00)hcp directions by a vector of 1/3(1 i00)hcp can transform the h.c.p, structure into f.c.c., and the reverse procedure transforms the f.c.c, structure to h.c.p (Liu and Zhang 1994a).

B. X. L i u

214

T a b l e 8.5 Structure data of the h.c.p.-I phase formed in five binary systems by ion-mixing. These phases were considered as metastable electron compounds. M X phase System Co-Au

Constituent metals

lattice constants

dcpp

lattice constants

dcpp

(nm)

(nm)

(nm)

(nm)

h.c.p,

d002 = 0.260

Co h.c.p,

d002 = 0.2023

a = 0.324, c = 0.521

a = 0.254, c = 0.407

c/a = 1.61

c/a = 1.59

Ref. Liu and Nicolet (1983)

Au f.c.c.

Ti-Au

h.c.p,

d002 = 0.263

a = 0.408

dl 11 = 0.2355

Ti h.c.p,

d002 = 0.234

a = 0.327, c = 0.526

a = 0.259, c = 0.469

c/a = 1.61

c/a = 1.59

Liu et al. (1982)

Au f.c.c.

Co-Mo

h.c.p,

d002 = 0.261

a = 0.408

d l l l = 0.2355

Co h.c.p,

d002 = 0.2023

a = 0.327, c = 0.522

a = 0.254, c = 0.407

c/a = 1.60

c/a = 1.59

Liu et al. (1983b)

M o b.c.c.

Ni-Mo

h.c.p,

d002 = 0.261

a = 0.3147

d110 = 0.2225

Ni f.c.c,

d l l l = 0.2034

a = 0.325, c = 0.521

a = 0.352

c/a = 1.60

Mo b.c.c. a = 0.3147

Ni-Nb

h.c.p, a = 0.328, c = 0.522

d002 = 0.261

Liu et al. (1983b)

d110 = 0.2225

Ni f.c.c. a = 0.352

d l l I = 0.2034

Kung et al. (1983)

Nb b.c.c.

c/a = 1.59

a = 0.3307

d110 = 0.233

Table 8.6 f.c.c.-I M X phases formed in Co-, Y-, Zr-, and Hf-rich multilayered films by room temperature 200 keV xenon ion mixing. System and composition

Mixing dose

Lattice parameter (nm)

Co94Mo6 ~ Co82Mo18

3 x 1015 X e + / c m 2

0.355 ~ 0.360

Co98Ta2 ~ Co80Ta20

3 • 1015 X e + / c m 2

0.355 ~ 0.365

Y73Mo27 ~ Y83MoI7

7 • 1015 X e + / c m 2

0.510

Y85TaI5 ~ Y75Ta25

1 x 1015 X e + / c m 2

0.518

Y25Zr75 ~ Y85ZrI5

3 • 1015 X e + / c m 2

0.467

Zr68Nb32 ~ Zr88Nb12

7 • 1014 X e + / c m 2

0.480

Hf75Nb25 ~ H f 9 0 N b l 0

5 • 1014 X e + / c m 2

0.455

Ion-Mixing

215

Table 8.7 h.c.p.-II and f.c.c.-II MX phases formed in Nb-, Ta-, Mo-rich multilayered films by room temperature 200 keV xenon ion-mixing. A-B

f.c.c.-b.c.c.

h.c.p.-b.c.c.

b.c.c.-b.c.c.

System and composition

h.c.p.-II

f.c.c.-II

Ni23Mo77 Ni20Nbs0 Ni35Ta65 Ag30Mo70 Ag33Nb67 Co20Mo80 Co23Nb77 Co30Ta70 Y22Mo78 Y25Ta75 Zr7Mo93 Zrl9Nbs1 Zr20Ta80 Hf17Nb83 Hf30TaT0 Fe24Mo76 Fe20Nbs0 Fe20Ta80

Yes

Yes

Yes No No No Yes Yes No Yes Yes Yes Yes No No No Yes

Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes

Yes

Yes

No

Yes

8.6.1.3 h.c.p.-H and f.c.c.-H phases based on b.c.c, metals In some Nb-, Ta-, and Mo-rich multilayered films, two MX phases, labeled h.c.p.-II and f.c.c.-II, were formed by ion-mixing, and their compositions were around AB3-4, in which B stands for the b.c.c metal. (see Table 8.7). Some of the systems have negative AHf, while combining with metals like Fe, Co and Ni and some have positive AHf when alloying with Ag, Y, Zr and Hr. Carefully performed experiments revealed that the h.c.p.II phase was observed as an intermediate phase in the process of forming the f.c.c.-II phase. Thus a two-step transition of b.c.c. --~ h.c.p. -+ f.c.c, was proposed and named as reverse martensitic transformation. First b.c.c, phase transforms to a h.c.p structure through a shearing mechanism. During shearing, the (111)bcc acted as the habit plane, while [ l l l]bcc (or [2110]hcp) acted as the shearing axis. The second step is to transform the h.c.p, phase to the f.c.c, phase. The shear can be achieved by sliding the atoms on planes parallel to the (0002)hcp along (1100)hcp directions by a vector of 1/3 (l l00)hcp. Although the relaxation period of ion-mixing is very limited, it is thought that it could meet the requirements of a shearing mechanism. The lattice parameters of h.c.p.-II and f.c.c.-II phases were deduced according to the two-step transition and these are in good agreement with the experimental results (Zhang 1995, Liu and Zhang 1992a).

B. X. Liu

216

8.6.2 Free energy calculation of the MX phases A metastable crystalline phase is of ordered structure, and so its entropy term can be neglected. The free energy of the MX phase is therefore A GMX = A H~I X -~- A Hi,ix -~- A H ~ x .

(8.12)

The expression of the chemical term is as follows A HI,cix = A HampX A I/" " A2/3 f A B ,

(8 13)

in which VA is the atomic volume of component A in the alloy, A Hamp is an amplitude concerning the magnitude of the electron redistribution interaction and is a constant for one specific metal system, faB is a function which accounts for the degree to which atoms of type A are surrounded by atoms of type B. In the calculation of faB, an empirical constant y is used to describe the short-range-order difference of solid solutions, amorphous phases, and ordered compounds, for which it is usually taken to be 0, 5, and 8, respectively. As an MX phase, except the solid solutions, is commonly considered as compoundlike, the value of y is therefore taken as 8 in the calculation. The elastic term of the MX phase can be derived from the equation A H ~ I X - - X A X B [ X A A H Me X ( B i n A )

where

-'F X B A

e , HMX(AinB)]

(8 . 14)

is the elastic contribution to the heat of solution of i in j, which can be calculated from the elastic constants of the constituent metals. The expression for the structural term is written as AOMx(iinj)e

A H ~ x = E(Z)

-

xAE(ZA)

--

xBE(ZB),

(8.15)

where E(Z), E(ZA) and E(ZB) are the structural stabilities of the alloy phase and the pure A and B metals, and Z, ZA and ZB are the mean number of valence electrons of the alloy phase and the number of valence electrons of pure metals A and B, respectively. For an MX phase at a composition in the vicinity of an intermetallic compound, the elastic term should be included. This term was from the difference in structures between the MX phase and that of the intermetallic compound. This is for the case that the MX phases were formed at A3B stoichiometry. Another situation is that the MX was formed at compositions near the eutectic point, where no equilibrium compound exists. In this case, the MX phase can be considered as ordered, so as the elastic term can be relaxed down to zero because of ordering. This situation corresponds to the formation of the MX phases with a composition of AB3. In short, with the above assumption, the formula for calculation for the A3B MX phase AGMx -- A G ~ x 4- A G ~ x + A G ~ x ,

(8.16)

and for the AB3 MX phase AGMx-

AG~x + AG~x.

(8.17)

Ion-Mixing

217

~" 12001-1 bcc(Mo) o.. I]

i) 91(P~

~ .

80(,[]

~

MoN

~- 6o0[/

i3

Y

4(t(t U

. . . . . . . . . 2(t 40 6(t

Mo

8()

Ni

(a)

~

"~ +8 E

bcc solution

fcc solution

..x +4

=

0

r

hcp-ll ~

~,, -4

~kJ "~

fcc-ll --

Mo

(b)

,A ....

9

20

9

J

40

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,~

60

,

9

9

84)

9

Ni

Ni concentration (at.%)

Figure 8.7. Equilibrium phase diagram of the Ni-Mo system; (b) Calculated free energy diagram of the Ni-Mo system.

The free energy curves of the MX phases formed in the Ni-Mo system are calculated, using the above equations, as an example. Figure 8.7(a) presents the equilibrium phase diagram. Figure 8.7(b) shows the constructed free-energy diagram of the system, in which the free-energy curves of all the concerned phases are included. One sees that the diagram is a little complicated because some newly formed MX phases are also included. Before employing the diagram to interpret the alloy phase formation, an important question to be considered is if the calculated results are relevant, since the model used in the calculations is still semi-quantitative. The author has proposed an alternative way for checking the relevance of the calculation, i.e., the steady-state thermal annealing of the multilayered films at some important compositions. The idea was to see whether the phase evolution in

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B. X. Liu

the films upon thermal annealing matches with that deduced from the calculation. Let us consider two examples in which the phase appearance in the samples is as follows:

300 ~ 450 ~ 650 ~ 800 ~ Niv5Mo25 30min > Am~ min > hcp-I20 min > Ni + hcp-II30 mi n > Ni + fcc-II

350~ 500~ 600~ Ni25Mo75~ > Amor.~ > hcp-II~ > fcc-II 60 min 30 min 30 min Apparently, the phase evolution sequences of these two reactions are in good agreement with that deduced from the calculated free energy diagram, thus confirming the applicability of the Miedema's model in this respect (Zhang and Liu 1994a). The author's group has also performed similar annealing experiments for several positive A H f systems and these results are also in support of the calculations (Chen and Liu 1996). The use of calculated free energy diagrams can explain the alloy phase formation upon IM, as discussed in the previous sections by some representative examples.

8.7. I N T E R F A C E - G E N E R A T E D

S O L I D - S T A T E V I T R I F I C A T I O N IN S Y S T E M S W I T H A

P O S I T I V E H E A T OF F O R M A T I O N

In 1983, Schwarz and Johnson reported for the first time that solid-state reaction (SSR) of multilayers consisting of two different crystalline metals (La and Au) can lead to the formation of an amorphous alloy. They argued that spontaneous vitrification was expected to take place in metal systems with a negative A H i and one metal diffusing anomalously fast in the other. However, the author's group has found very recently that in some binary metal systems with positive A H f , vitrification could be achieved by SSR, e.g., in the HfTa (Liu et al. 1996), Y-refractory metals (Chen and Liu 1997) and Cu-Ta (Pan et al. 1998) systems. Consider the Y-Nb system as an example, Figure 8.8 shows the X-ray diffraction patterns of the Yv2Nb28 multilayered films before and after thermal annealing. It should be emphasized that similar solid-state vitrification was also observed by author's group very recently in some other systems with a positive A H f (Chen and Liu 1997a). To test whether a large size difference is a decisive prerequisite in solid-state amorphization, the Au-Ta system with a negative A H f o f - 4 9 kJ/mol (Pan et al. 1995) was selected as the atomic radius ratio and the volume ratio between Au and Ta are 1.008 and 0.936, respectively. When the Auz3Ta77 multilayers were heated up continuously at a rate of 5 K/min to 200~ partial amorphization was observed. After maintaining at 200~ for 30 min, the annealing temperature was raised to 350~ and kept for 30 min. The structure of the multilayers turned entirely into an amorphous state. When the temperature was further raised to about 650~ and kept for 20 min, diffraction lines from metals Au and Ta emerged again, indicating crystallization of the Auz3Ta77 amorphous alloy. These results suggest that solid-state vitrification can also be achieved in a system with a small atomic size difference.

Ion-Mixing

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J

.1500

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.

.

.

0i 0

.,,.i

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~0

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Figure 8.8. X-ray diffraction patterns of the Y72Nb28 multilayered films with 121 layers and a total thickness of 350 nm before and after thermal annealing: (a) as-deposited; (b) amorphized by annealing at 300~ for 1.5 hours.

8.8. C O N C L U D I N G R E M A R K S

(1) In the past 15 years, study of IM of multilayered films scheme has provided extensive data, which promoted significantly the non-equilibrium processing of materials. IM has so far synthesized a great number of new amorphous as well as new MX alloys. It is anticipated that this technique is capable of synthesizing more new metallic glasses in those not yet studied systems known as non-glass-forming ones. (2) Miedema's theory was employed for free energy calculation and steady-state thermal annealing results showed that the calculated energetic sequence of the related phases was relevant, at least in its outline, suggesting the applicability of Miedema's theory in this respect. (3) Taking into account the interfacial free energy, a unified thermodynamic model was proposed for predicting the possibility of metallic glass formation by IM in binary metal systems with either negative or positive AHz, and for a reasonable explanation for the formation of MX phases. (4) The interface concept was also tested by solid-state reaction experiments in some positive A Hf systems and it turned out that metallic glasses were indeed obtained by thermal annealing in the properly designed multilayered films with enough fraction of interfacial atoms. These results offer a neat evidence demonstrating the decisive role of the interfacial free energy in amorphization taking place in the multilayered films.

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B. X. Liu

(5) Kinetically, due to the very restricted conditions available in IM for a crystalline phase to nucleate and grow, only simple-structured phases can be formed and they feature a close crystallographic correlation with their matrix. (6) Based on the IM and SSR studies, a new glass-producing technique is developed, i.e., the multilayer technique, in which the multilayered films are firstly designed according to the interfacial free energy calculation and then either by thermal annealing or ion irradiation. (7) Further studies are still needed in this area, especially in the systems with a positive A Hf. For instance, detailed growth of the amorphous phase, local microstructures and the abnormal properties of the new metallic glasses formed in such systems are highly interesting topics to be investigated by IM as well as SSR technique.

ACKNOWLEDGMENTS

The author is grateful to his current and previous students for their continuous contributions to this research project, especially to Mr. Y.G. Chen for his substantial assistance in writing and finalizing this chapter. Special thanks go to the Editor of this book, Professor C. Suryanarayana, for his critical reading of the manuscript of this chapter and for many helpful suggestions. The continuous financial support from National Natural Science Foundation of China under several projects and by the Administration of Tsinghua University are greatly appreciated.

REFERENCES Alonso, J. A. and Simozar, S. (1983) Solid State Comm., 48, 765. Alonso, J. A., Gallego, L. J. and Simozar, S. (1990) Nuovo Cimento, D12, 587. Andersen, U. A., BCttiger, J. and Dyrbye, K. (1990) Nucl. Instru. Meth. Phys. Res., B51, 125. Bcttiger, J., Pampus, K. and Torp, B. (1987) Nucl. Instr. Methods Phys. Res., B19/20, 696. Bcttiger, J., Dyrbye, K., Pampus, K. and Poulsen, R. (1989) Phil Mag., A59, 569. Brenier, R., Thevenard, P., Capra, T., Perez, A., Treilleux, M., Romana, L., Dupuy, J. and Brunel, M. (1987) Nucl. Instr. Methods Phys. Res., B19/20, 691. Carter, G. and Grant, W. A. (1976) Ion Implantation of Semiconductors. (Edward Arnold Ltd., London). Chen, Y. G. and Liu, B. X. (1996) Appl. Phys. Lett., 68, 3096. Chen, Y. G. and Liu, B. X. (1997a) J. Alloys and Compounds, 261, 217. Chen, Y. G. and Liu, B. X. (1997b) Nucl. Instr. Meth. in Phys. Res., B127/128, 145. Chen, Y. G. and Liu, B. X. (1997c) J. Appl. Phys., 82, 3815. Chen, Y. G. and Liu, B. X. (1988) Mater Sci. Eng., B ( to be published). Chen, Y. G., Zhang, Q. and Liu, B. X. (1997) Nucl. Instr. Methods Phys. Res., B124, 523. Cheng, G. A. (1986) Thesis, Tsinghua University, Beijing, China. Cheng, Y. T., Johnson, W. L. and Nicolet, M.-A. (1984) in Proceedings of MRS Fall Meeting, Boston, MA. Chu, W. K., Mayer, J. W. and Nicolet, M.-A. (1978) Backscattering Spectrometry. (Academic Press, New York). Clemens, B. M. (1986) Phys. Rev., B33, 7615. Davies, H. A. (1983) in Amorphous Metallic Alloys, ed. Luborsky, EE. (Butterworths, London), p.13. de Boer, F. R., Boom, R., Miedema, A. R., Niessen, A. K., and Mattens, W. C. M. (1989) Cohesion in Metals: Transition Metal Alloys. (North-Holland, Amsterdam). Dearnaley, G., Freeman, J. H., Nelson, R. S., and Stephen, J. (1973) Ion Implantation. (North-Holland Publishing Company, Amsterdam). Ding, J. R., Che, D. Z., Zhang, H. B., Tao, K. and Liu, B. X. (1992) Appl. Phys. Lett., 60, 944. Gerkema, J. and Miedema, A. R. (1983) Surf. Sci., 124, 351. Giessen, B. C. and Whang, S. H. (1980) J. Physique, C8, 95.

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Goltz, G., Fernandez, R., Nicolet, M.-A. and Sadana, D. K. (1982) in Metstable Materials Formation by Ion Implantation. eds. Picraux, S. T. and Choyke, W. J. (North-Holland, Amsterdam), p. 227. Hafner, J. (1981) in Metallic Glasses, eds. Gtintherodt, H. J. and Beck, H. (Springer-Verlag, Berlin), p. 91. Hiller, W., Buchgeister, M., Eitner, P., Kopitzki, K., Lilienthal, V. and Peiner, E. (1989) Mater Sci. Eng., All5, 151. Huang, L. J. and Liu, B. X. (1987) Nucl. Instr. Methods Phys. Res., B18, 256. Hung, L. S. and Mayer, J. W. (1985) Nucl. Instr. Methods Phys. Res., B7/8, 657. Hung, L. S., Nastasi, M., Gyulai, J. and Mayer, J. W. (1983) Appl. Phys. Lett., 42, 672. Jin, O. and Liu, B. X. (1994) J. Phys.: Condensed Matter, 6, L39. Jin, O. and Liu, B. X. (1995) J. Phys.: Condens. Matter, 7, 2667. Jin, O. and Liu, B. X. (1996) Nucl. Instr. Meth. Phys. Res., Bl14, 56. Jin, O. and Liu, B. X. (1997) J. Non-Cryst. Solids, 211, 180. Jin, O., Ye, Y. and Liu, B. X. (1995a) J. Phys. C: Condens. Matter, 7, 7959. Jin, O., Zhang, Z. J. and Liu, B. X. (1995b) Appl. Phys. Lett., 67, 1524. Jones, H. (1960) in The Theory of Brillouin Zones and Electronic States in Crystals. (North-Holland Publ. Co., Amsterdam), p. 191. Klement, W., Willens, R. H. and Duwez, P. (1960) Nature, 187, 869. Knapp, J. A. and Follstaedt, D. M. (1987) Nucl. Instr. Methods Phys. Res. B19/20, 611. Kung, K. T.-Y., Liu, B. X. and Nicolet, M.-A. (1983) Phys. Stat. Sol., (a) 77, 355. Lin, C. J. and Spaepen, E (1984) in Rapid Solidification of Metastable Materials, eds. Kear, B. H. and Giessen, B. C. (North-Holland Publishing Co., Amsterdam). Liu, B. X. (1982) California Institute of Technology, Pasadena, CA, unpublished research. Liu, B. X. (1983) Phys. Stat. Sol., (a) 75, K77. Liu, B. X. (1985) Nucl. Instr. Methods Phys. Res., B7/8, 547. Liu, B. X. (1986) Phys. Stat. Sol., (a) 94, 11. Liu, B. X. (1987) Mater. Lett., 5, 322. Liu, B. X. and Jin, O. (1997) Phys. Stat. Sol., (a)161, 3. Liu, B. X. and Nicolet, M.-A. (1983) Thin Solid Films, 101, 201. Liu, B. X. and Zhang, Z. J. (1994a) J. Mater. Res., 9, 357. Liu, B. X. and Zhang, Z. J. (1994b) Phys. Rev., B49, 12519. Liu, B. X., Nicolet, M.-A. and Lau, S. S. (1982) Phys. Stat. Sol., (a) 73, 183. Liu, B. X., Johnson, W. L., Nicolet, M.-A. and Lau, S. S. (1983a) Appl. Phys. Lett., 42, 45. Liu, B. X., Johnson, W. L., Nicolet, M.-A. and Lau, S. S. (1983b) Nucl. Instr. Meth., 209/210, 229. Liu, B. X., Clemens, B. M., Gaboriaud, R., Johnson, W. L. and Nicolet, M.-A. (1983) Appl. Phys. Letters, 42, 624. Liu, B. X., Huang, L. J., Li, J. and Ma, E. (1986) in Thin Films-Interfaces and Phenomena, eds. Nemanich, R. J., Ho, P. S. and Lau, S. S. (North-Holland, Amsterdam), p. 215. Liu, B. X., Ma, E., Li, J. and Huang, L. J. (1987a) Nucl. Instr. Methods. Phys. Res., B19/20, 682. Liu, B. X., Ma, E., Li, J., and Huang, L. J. (1987b) NucL Instr. Methods. Phys. Res., B19/20, 682. Liu, B. X., Ding, J. R., Zhang, H. B. and Che, D. Z. (1991) Phys. Stat. Sol., (a) 128, 103. Liu, B. X., Che, D. Z., Zhang, Z .J., Lai, S. L. and Ding, J. R. (1991) Phys. Stat. Sol., (a) 128, 345. Liu, B. X., Zhang, Z. J., Jin. O. and Pan, E (1995) Nucl. Instr. Methods Phys. Res., B106, 17. Liu, B. X., Jin, O. and Ye, Y. (1996) J. Phys. C: Conden. Matter, 8, L79. Liu, B. X., Zhang, Z. J., and Jin, O. (1998) Mater. Lett. (submitted). Lopez, J. M. and Alonso, J. A. (1985) Z. Naturforsch., A40, 1199. Mayer, J. W. and Lau, S. S. (1988) Electronic Materials Science: For Intergrated Circuits in Si and GaAs. (Macmillan Publishing Company, New York). Mayer, J. W., Tsaur, B. Y., Lau, S. S. and Hung, L. S. (1981) Nucl. Instr. Methods, 182/183, 1. Meissner, J., Kopitzki, K., Mertler, G. and Peiner, E. (1987) Nucl. Instr. Methods Phys. Res., B19/20, 669. Miedema, A. R. (1976) Philips Tech. Rev., 36, 217. Miedema, A. R. (1978) Z. Metallkde., 69, 287 and 455. Murr, L. E. (1975) Interfacial Phenomena in Metals and Alloys. (Addison-Wesley, London). Niessen, A. K. and Miedema, A. R. (1983) Ber. Bunsenges. Phys. Chem., 87, 717. Niessen, A. K., Miedema, A. R., de Boer, E R. and Boom, R. (1988) Physica, B151, 401. Ossi, P. M. (1990) Phys. Star. Sol., (a) 119, 463. Pan, E and Liu, B. X. (1995) J. Phys.: Condens. Matter, 8, 383.

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Pan, E et al., (1998) unpublished data. Pan, E, Chen, Y. G. and Liu, B. X. (1995) AppL Phys. Lett., 67, 780. Pan, E, Chen, Y. G., Zhang, Z. J., and Liu, B. X. (1996) J. Non-Cryst. Sol., 194, 305. Peiner, E. and Kopitzki, K. (1988) Nucl. Instrum. Methods Phys. Res., B34, 173. Rauschenbach, B. (1986) Appl. Phys., A40, 47. Rauschenbach, B. and Hohmuth, K. (1987) Nucl. Instr. Methods Phys. Res., B23, 323. Schmid, A. and Ziemann, P. (1985) Nucl. Instr. Methods Phys. Res., B7/8, 581. Schwarz, R. B. and Johnson, W. L. (1983) Phys. Rev. Lett., 51, 415. Shin, S. M., Ray, M. S., Rigsbee, J. M. and Greene, J. E. (1983) Appl. Phys. Lett., 43, 249. Takayama, S. (1976) J. Mater. Sci., 11, 164. Thompson, M. W. (1969) Defects and Radiation Damage in Metals. (Cambridge University Press, Cambridge). Townsend, E D., Kelly, J. C. and Hartley, N. E. W. (1976) Ion Implantation, Sputtering and their Applications. (Academic Press, London), p. 31. Tsaur, B. Y. (1980) Thesis, California Institute of Technology, Pasadena, CA. Tsaur, B. Y., Lau, S. S., Hung, L. S. and Mayer, J. W. (1981) Nucl. Instr. Methods, 182/183, 67. Turnbull, D. (1955) Impurities and Imperfections. (Amer. Soc. Met., Cleveland), p. 121. van Rossum, M., Shreter, U., Johnson, W. L. and Nicolet, M.-A. (1983) in Proceedings of MRS Fall Meeting, Boston, MA. von Allmen, M. (1983) in Laser-Beam Interaction with Materials, ed. Carl, J. (Washington International Society of Optical Engineering, Washington), p. 140. Yan, Z. H., Liu, B. X., and Li, H. D. (1986) Phys. Stat. Sol., (a) 94, 483. Yan, Z. H., Liu, B. X., and Li, H. D. (1987) Nucl. Instr. Methods Phys. Res., B19/20, 700. Zhang, Z. J. (1995) Ph.D. Thesis, Tsinghua University, Beijing, China. Zhang, Z. J. and Liu, B. X. (1994a) J. Appl. Phys. 76, 3351. Zhang, Z. J. and Liu, B. X. (1994b) J. Appl. Phys., 75, 4948. Zhang, Z. J. and Liu, B. X. (1994c) J. Phys.: Condens. Matter, 6, 2647. Zhang, Z. J. and Liu, B. X. (1994d) J. Phys.: Condens. Matter, 6, 9065. Zhang, Z.J. and Liu, B. X. (1995) Phys. Rev., B51, 16475. Zhang, Z. J., Bai, H. Y., Qiu, Q. L., Yang, T., Tao, K. and Liu, B. X. (1993) J. Appl. Phys., 73, 1702. Zhang, Z. J., Jin, O. and Liu, B. X. (1995) Phys. Rev., B 51, 8076.

Chapter 9 Physical Vapor Deposition 9.1. Introduction 9.2. Development of PVD 9.3. Deposition Methods 9.3.1 Evaporation 9.3.2 Molecular Beam Epitaxy (MBE) 9.3.3 Sputtering 9.3.4 Ion-Assisted Deposition 9.3.5 Magnetron Sputtering 9.4. Influence of Energy on Coatings 9.4.1 Film Morphology and Density 9.4.2 Nucleation and Adhesion 9.4.3 Metastable Phases 9.4.4 Microstructure 9.5. Applications of PVD Coatings 9.5.1 Optical Coatings 9.5.2 Corrosion Protective Coatings 9.5.3 Hard Coatings 9.6. Future Trends Acknowledgment References

225 225 228 228 229 229 231 235 238 238 239 240 240 243 243 245 246 248 249 249

This Page Intentionally Left Blank

Chapter 9 Physical Vapor Deposition JOHN S. COLLIGON

9.1. INTRODUCTION Physical Vapor Deposition (PVD) is the general name given to coating processes where the transport of material to the substrate is effected by a physical driving mechanism. Such mechanisms include evaporation, sputtering, ion-plating, and ion-assisted sputtering. Often chemical reactions do take place to form coatings of oxides, nitrides and carbides, but the transport process does not generally rely on chemical reactions. The coated samples can usually be maintained at a temperature between ambient and 100~ during the deposition and the working pressure is below 10 -2 Pa. PVD thus differs quite markedly from Chemical Vapor Deposition (CVD) which occurs at atmospheric pressure, or just below, and at temperatures normally in excess of 500~ CVD also requires the use of complex gaseous species (such as fluorides, organometallics) which need to be handled very carefully whereas PVD requires either no gas, for evaporation, or, relatively simple species (Oxygen, Nitrogen) for most of the other PVD processes. Recent improvements in PVD techniques, in particular the use of plasmas and the unbalanced magnetron for the deposition process, have led to its wide application in industry. Such coatings are used for corrosion protection, to provide hard surfaces for engineering applications, to form decorative coatings on a range of household and industrial products, for energy-saving window glass, for media storage systems such as compact discs and for a range of microelectronics applications. Table 9.1 lists some of the conference series that cover this huge field. The present chapter starts with a brief historical review of the subject. It then describes the main PVD methods and continues with a discussion of the relation between deposition conditions and film properties. Some of the main applications of PVD coatings are then described and the chapter concludes with an outline of some of the possible future trends in this important industrial field.

9.2. D E V E L O P M E N T

OF P V D

The first recorded mention of a PVD process was by Grove (1852) in 1852 who observed the formation of a deposit in a glass tube in which a discharge had been running and identified this as being the same material as one of the electrodes; this was the first reported evidence of the ejection of atoms from an ion-bombarded material, now known as sput225

226

J. S. Colligon

Table 9.1. Established conference series containing papers on PVD. 9 9 9 9 9 9 9 9 9 9 9 9 9 9

InternationalConference on Metallurgical Coatings and Thin Films (ICMCTF) AnnualTechnical Conference of the Society of Vacuum Coaters E-MRSSpring Meeting InternationalConference on Crystal Growth and International Conference on Vapor Growth and Epitaxy InternationalVacuum Congress and International Conference on Solid Surfaces InternationalConference on Thin Films InternationalConference on Nanometer-scale Science and Technology (Nano) InternationalConference on Ion Beam Modification of Materials (IBMM) InternationalConference on Surface Modification of Metals by Ion Beams (SMMIB) InternationalConference on Advances in Surface Engineering EuropeanConference on Diamond, Diamond-like Materials, Nitrides and Silicon Carbide (Diamond) InternationalConference on Plasma Surface Engineering (PSE) NationalSymposiumof the American Vacuum Society MRS Fall Meeting

tering and one of the main PVD processes. A few years later Faraday (1857) exploded metal wires in vacuum and thereby produced an evaporated film. Many years then passed before the concept of ionising some of the depositing species was proposed by Berghaus (1938) and, subsequently exploited by Mattox (1963) in the process known as ion-plating. Ion-plated coatings were seen to have superior adhesion properties which were identified with the additional energy imparted to some of the evaporated atoms which had been ionized and to a general increase in energy of support gas and other species bombarding the growing film. The process operated at a relatively high pressure (0.5 Pa) however, so that many collisions occurred in the gas phase which tended to reduce the energies of ionized species and increase those of neutral particles. The many collisions allowed depositing species to be bounced around shaped substrates to the rear side, giving the technique so-called high throwing power, but the parameters were not easy to control. Weissmantel (1976) proposed the use of ion beams in the deposition process where material was either evaporated onto a surface being simultaneously bombarded by ions, or was sputtered by one ion beam and the deposited film bombarded by a second beam. Here the precise control of the added energy and ion to atom arrival ratio could be achieved. A review of the early work was given by Weissmantel et al. (1979). This type of system, further exploited by Harper et al. (1982a, 1982b, 1983), Sainty et al. (1984), Martin et al. (1983, 1984), Colligon (1995), Kheyrandish et al. (1994) and many others (Rossnagel and Cuomo 1988), has allowed controlled studies of the influence of added energy on film composition, stress, density, microstructure, hardness, wear and durability to be made. An alternative approach has been proposed by Yamada and Takagi (1995) who developed an ion cluster-beam method. In this method, illustrated schematically in Fig. 9.1, a beam of clusters of atoms from 2 - 2 0 in each cluster is created by expanding a vapor of the metal through a nozzle. Some atoms in the clusters are then ionized by an electron beam so that the whole cluster may be accelerated toward the negative substrate. How-

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ce material Figure 9.1. Schematic diagram of ionized cluster-beam system

ever, all the above ion beam systems have a common limitation in that they do not allow the production of coatings over large areas (normally a maximum of a few square centimeters) nor do they allow high growth rates to be attained (normally maximum rates are 0.2-1 nm/s). There are some developments in the cluster-beam method by Gspann (1993) who has recently produced cluster-beam sources with much improved deposition rates but the treated areas are still of order of a few square centimeters. The major breakthrough in industrial processing has however come from the new generation of unbalanced magnetron systems (Window and Savvides 1986a, b, Savvides and Window 1986) which have, in the 1990s, replaced many simple evaporation methods and ion-plating techniques which had been the main coating methods during the period from the 1960's. Film deposition rates of 10-500 nm/s are now possible and the sample area can be several square meters. Similar rates of deposition, but not such large area capability, have been found for recent plasma-based systems, including the pulsed plasma methods (Conrad et al. 1990). All these systems can be developed to produce reliable coatings and, in some cases, can produce metastable phases of alloy materials because the energetic atoms are suddenly trapped into rapidly cooling surroundings. It is not always possible to relate the film properties to the deposition parameters because there are complex interactions between magnetron power, system pressure, ion/atom arrival ratio and deposition energies. Fundamental studies on new coatings are, therefore, often best undertaken on ion-beam systems.

J. S. Colligon

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9.3. D E P O S I T I O N M E T H O D S

This review will outline the main features of the various PVD methods but will concentrate more on the more recent ion-assisted deposition (IAD) techniques where additional energy is provided to the growing film by ion-bombardment. Earlier studies of traditional evaporation and sputtering methods have been fully reviewed in several textbooks by Holland (1956), Chopra (1969), Vossen and Kern (1978) and Bunshah (1982).

9.3.1 Evaporation In this process materials are vaporised by heating using direct resistance-heating, radiation, eddy current heating, electron-beam heating, or running an arc discharge. One advantage of the technique is that the evaporation can be carried out in a UHV chamber so that contamination of the substrate by residual gases is minimal. For best purity an electron-beam heating system should be used as the material can then be evaporated from the center of the crucible in which it is held and reactions with the crucible walls are minimised. A schematic diagram of a typical e-beam coating system is shown in Fig. 9.2. Fundamental equations for evaporation rates and tables of evaporation temperatures and procedures are given in Chapter 4 of the book edited by Bunshah (1982). Some decomposition of chemical compounds always occurs and special procedures have to be followed if fully stoichiometric films are to be formed. The energies of species arriving at the substrate are only of the order of 0.05-0.1 eV which means they have limited energy for surface migration to find good nucleation sites. Substrates must be free of contaminants to allow reasonable levels of film adhesion to be attained so considerable attention has to be paid to cleaning procedure. Because of the lack of surface mobility

Physical Vapor Deposition

229

there is a high probability that the coating will contain voids and will be less dense than the bulk material. This leads to the formation of stress in the coating, usually tensile, which often results in the film delaminating and curling up.

9.3.2 Molecular beam epitaxy (MBE) A special technique of evaporation has been developed, primarily for the semiconductor industry, in which component species are injected into a UHV chamber from gas effusion cells and form the desired film on a substrate. This is known as Molecular Beam Epitaxy. In some systems organometallic gases are used which react at the high temperature substrate to form the desired deposit and a residual volatile species. Following the work of Giinther (1958) an alternative three-temperature method was developed which allowed III-V semiconductor compounds to be deposited from the pure constituent elements; usually the type-III element cell was run at a higher temperature than that for the more volatile type-V element and the substrate temperature was chosen to be high enough to re-evaporate excess type-V species. A review of fundamental aspects of MBE has been published by Foxon and Joyce (1981). In later work an important correlation between the oscillations in intensity of RHEED signals and layer coverage was discovered, allowing the number of layers deposited to be calibrated in situ (Neave et al. 1983). The whole field is well reviewed by Herman and Sitter (1989) and Chow (1991) and a comprehensive bibliography of early data relating to MBE of III-V compounds (from 1958-1983) has been published by Ploog and Graf (1984). Later improvements have allowed the deposition conditions to be improved by applying reactive gas in pulses (Goodman and Pessa 1986) which allows growth at lower temperatures and hence reduces diffusion to produce sharper boundaries. The choice of gas species and effusion cell temperature is very important as the structure of the molecule forming the film (e.g., dimer or trimer) is of major importance. The MBE process has been linked to other stages in the semiconductor production line as described, for example, by Migauchi and Hashimoto (1986) and by Melngailis (1987).

9.3.3 Sputtering When an ion strikes a material its energy is transferred by a series of collisions to the atoms of the first few layers of that material. Some of these atoms will move toward the surface and, if their energy is greater than the binding energy they will escape. The ratio of number of ejected atoms per incident ion is known as the sputtering coefficient and, for an incident ion of energy about 1 keV, it has a value for most materials ranging between 0.1 and 2 atoms per ion. A full review of the mechanisms and major trends in the sputtering phenomena was published by the author in 1968 (Carter and Colligon 1968, Chapter 7). The sputtering coefficient increases, after an initial quadratic behavior below 100 eV, almost linearly with ion energy as the ion range increases and more and more atoms participate in the energy exchange process. However, when the ion energy increases such that the ion range exceeds several atomic layers, the energy is too deep to be transferred back to the surface and the sputtering coefficient begins to decrease again.

230

J. S. Colligon

The fundamental understanding of the process has not changed significantly since the late 1960s but the theory of the process has been fully developed, the model by Sigmund (1969, 1971) proving to be very successful in predicting sputtering coefficients. Of particular importance to thin film deposition is the energy of the sputtered species which has a peaked distribution with a maximum at about 4-7 eV and a long tail down to over 50 eV. These energies are more than an order of magnitude greater than those of evaporated species and allow the depositing species more mobility on the surface to form denser, well-bonded films. One problem with the method is the change in composition of the surface of a sputtered compound material which arises from preferential sputtering of components. This results primarily as a result of the different energy transfer to different mass species. Fortunately, after a period of bombardment, the system is selfcompensating since the number of atoms of the species sputtering most readily is reduced and the rate of further sputtering of this component returns to the stoichiometric level. It is always important therefore to sputter a compound target for a period of time and to cover the substrate during this stabilizing period. Another problem from compound targets is that the spatial distributions of component ejected species can be different. Thus a substrate at a specific position in space facing the target may not collect component atoms in the correct stoichiometric proportions, even though the global total sputtering of species is stoichiometric. Clearly it is preferable, where possible, to move substrates so that they sweep through these spatial distribution variations. Apart from this problem, the process of sputtering is very versatile and control of the deposition rate is relatively simple. The first commercial systems were based on so-called planar diode sputtering and had the form illustrated in Fig. 9.3. Electrons are accelerated towards the anode and have ionizing collisions with gas atoms. These ions then bombard the cathode, causing sputtering and the liberation of further electrons which cause the discharge to be sustained. There are also many ion-gas collisions leading to charge exchange so that the target is bombarded by a mixture of ions and neutrals with energies below the total applied voltage. Typical operating conditions for dc metal sputtering are; cathode current density: 1 mA/cm 2, discharge voltage: 3 kV, argon pressure: 10 Pa, cathode-to-substrate distance: 40 mm, deposition rate: 40 nrn/min. For sputtering of insulating materials a radio-frequency (rf) supply has to be used, usually at the allowed 13.56 MHz frequency. These rf systems run at lower pressures, in the range 1-2 Pa. Although materials can be deposited using planar sputtering systems, the deposition rate is rather low, the ionization efficiency in the plasma is also low and the substrate becomes heated. These limitations have been overcome by the introduction of magnetron sputtering systems which are described in the next section. Deposition rates in these systems are as high as several microns per minute and can be controlled by varying the power density supplied to the target region, the target area, coating pressure and target-substrate distance. Although the substrates are near to the ionized plasma region the degree of ion bombardment of the growing film is low in these systems. The film quality is thus similar to that obtained for other sputtering methods but with a slight improvement in film density arising from the lower operating pressure; atoms sputtered from the target do not suffer many collisions in transit and therefore maintain, on average, a higher energy.

231

Physical Vapor Deposition

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9.3.4 Ion-assisted deposition

Evidence that additional ion bombardment of a surface being coated improved the adhesion and film properties was demonstrated in the early 1960s by Mattox (1963, 1965) using a system known as ion-plating, which was patented much earlier by Berghaus (1938). A typical ion-plating system is illustrated schematically in Fig. 9.4. The coating material is evaporated using an electron-beam and a discharge is struck in the upper region using a support gas such as argon or, for nitrides, nitrogen. Operating pressures are generally in the region 0.1-1 Pa which means that ions and atoms in the discharge make many collisions in passing to the negatively biased substrate. The substrate is thus bombarded by a variety of atomic, ionized and molecular species, including ions and atoms of the coating material with energies from thermal values up to the applied voltage of several keV. A further effect is the irradiation of the substrate and growing film by photons from the discharge. These energy-assisted deposition conditions have been shown to produce coatings with improved adhesion (Mattox 1965, Rigsbee et al. 1986, McLeod and Mah 1974, Enomoto and Matsubara 1975, Takagi et al. 1974, Teer 1983, Ahmed and Teer 1984), increased density (Mattox and Kominiak 1972), improved abrasion resistance (Lardon et al. 1978, Broszeit et al. 1985), and enhanced optical reflectivity (Henderson 1977, Varasi et al. 1985). However the film-growth parameters cannot be controlled independently; for example, a change in pressure also effects gas collisions and therefore energies of arriving particles.

232

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An altemative plasma-based ion-assisted deposition system is the arc-plating system. Here a discharge is run in the abnormal glow region of the current-voltage characteristic for a diode system. Extremely high (many amperes) currents flow in the discharge which forms an arc with the current entering the target at a small (few mm dia) spot. Target material is vaporized and, in the intense discharge, a high fraction of this material is then ionized. A biased substrate will therefore collect both atoms and ions of the film species (together with neutral species). The arc is struck with an igniter (used to produce fieldemission of electrons) and starts with additional gas in the chamber (such as argon) but, the arc once struck, the argon can be removed and the discharge maintained in the vapor. In operation the arc moves randomly over the target surface but this can be controlled by applying a movable magnetic field (steered arc system). Film quality is as good as for ion-plating with the exception that there is a tendency for large droplets of target material to erupt from the surface and fall on the substrate, thus spoiling film uniformity. This problem can be minimized by running at lower arc currents (but with associated lower deposition rates) or by moving a magnet below the cathode to "steer" the arc (Ramalingam et al. 1987, Erttirk and Heuvel 1989). Later versions have however introduced an arc deflecting column to entrain the arc around a 90 degree bend so that the substrate may be

233

Physical Vapor Deposition

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mounted out of the line of sight from the target. This is known as the filtered arc system and is shown in Fig. 9.5 (Vyskocil and Musil 1990). In this situation only the ionized fraction of the ejected material will reach the substrate but this can be high (approaching 100% in some systems) so even this geometry provides high deposition rates. To provide controlled energy-assisted deposition for fundamental studies of the process, ion beam systems were developed in the 1980s. As mentioned earlier, the first of these was the apparatus devised by Weissmantel and his colleagues (Weissmantel 1976, Weissmantel et al. 1979) but many other groups produced similar systems (Harper and Cuomo 1982, Harper et al. 1982a, b, 1983, Sainty et al. 1984, Martin et al. 1983, 1984, Colligon 1995, Kheyrandish et al. 1994). Figure 9.6 shows some of the various possibilities for ion-assisted deposition and includes an e-beam system for evaporation plus partial ionization (PIB, Fig. 9.6(a)) (Bai 1990)), ion-beam deposition (IBD, Fig. 9.6(b)) (Colligon et al. 1976, Lawson et al. 1975, Armour et al. 1986, Sharples et al. 1987), dual ion-beam deposition (DIB, Fig. 9.6(c)) (Harper and Cuomo 1982, Harper et al. 1982a, 1982b, 1983, Sainty et al. 1984, Martin et al. 1983, 1984, Colligon 1995, Kheyrandish et al. 1994), single-beam plus evaporation (SBE, Fig. 9.6(d)) (Weissmantel 1976, Weissmantel et al. 1979), ion-beam-induced chemical deposition (IBC, Fig. 9.6(e)) (Harriott and Vasile 1988) and ionized cluster-beam deposition (ICB, see Fig. 9.1) (Gspann 1993, Takagi et al. 1977, Yamada and Takagi 1995, Yamada 1994). In many of the ion-beam systems the Kaufman source was used (Kaufman et al. 1982) which produces broad beam gaseous ion species with a relatively low energy spread over an energy range 50-2000 eV and with intensity of order 1-5 mA/cm 2. Most of the above techniques are well known and the principle of operation is clear from the diagrams. In IBD the ions are usually extracted from an ion-implanter source at energies of 10-20 keV and, after magnetic separation, are slowed down to about 10 eV using carefully designed ion-optical lenses to prevent beam divergence and to steer the ions

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away from the direct line of sight where fast neutrals, created in charge-exchange collisions, would bombard and probably remove the deposited species (Colligon et al. 1976, Lawson et al. 1975, Armour et al. 1986, Sharples 1987). Such systems allow epitaxial growth of materials and can even grow isotopically pure films (Colligon et al. 1976, Lawson et al. 1975). In IBC a suitable organometallic vapor has to be directed at the substrate surface and the ions simply disrupt this to give a deposit plus a volatile species. Harriott

Physical Vapor Deposition

235

Table 9.2 Principal ion-beam-assisteddeposition methods. Method

Main features

Main application

PIB

Evaporated beam containing about 1% ionized species Direct deposition of 1-10 eV ions Few cm2, low rate One beam sputters One beam ion-assists Several cm2, low rate Precise control of IAD Evaporation plus ion-bombardment Several cm2, low rate Ion-beam dissociates chemical Vapor species at substrate Several cm2 Clusters of 100-1000 atoms About 1% are ionized Arrival energy about 10 eV/atom Several cm2

Epitaxial layers

IBD DIB

SBE IBC

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Epitaxial layers isotopic films Novel compounds Metastable films

Novel compounds Metastable films Metal contacts Hard coatings Mask repair Metal contacts Epitaxial layers

and Vasile (1988) have used a fine-focused liquid-metal ion source with spot size just below 1 # m dia to accurately lay down carbon films and hence repair minute sections of damaged photolithographic masks. ICB systems operate by forcing an evaporant through fine nozzle which produces clusters of atoms of 2-5 nm size containing some 500-2000 atoms. Atoms in such clusters appear to have a binding energy lower than that in the bulk material (Anderson 1976) and the clusters have a low melting temperature (Buttet and Borel 1976). An e-beam crosses the path of the clusters and causes ionization of one or more atoms in the cluster which is then accelerated by applying a negative voltage of about 1 keV to the substrate. On impact the energetic atoms in the cluster disperse across the substrate surface and can produce epitaxial growth even on substrates where there is a large lattice misfit with the film material (Lawson et al. 1975). Later improvements in the temperature control of the nozzle have allowed higher deposition rates of 10-100 nrn/s to be attained for Zn (Gspann 1993), Ag (Gspann 1993) and Mg (Hagena et al. 1994). A summary of these ion-assisted methods and their main application is given in Table 9.2.

9.3.5 Magnetron sputtering All the above ion beam processes are important for fundamental studies of the energyassisted coating process and can be used for small-area coating where film thickness is well below 1/zm. A much quicker coating technique is based on the magnetron sputtering system which has now been adapted to provide a degree of ion bombardment of the growing film. A series of magnetron systems are illustrated in Fig. 9.7. In the conventional dc

236

J. S. Colligon

magnetron, shown schematically in Fig. 9.7(a), electrons are constrained to follow spiral orbits about the lines of magnetic flux which follow paths parallel to and above the target surface. As these electrons pass through the region above the target they have collisions with gas atoms, some of which ionize. If the target is biased to a sufficiently negative voltage, some of these ions will be attracted to it and cause sputtering. Sputtered atoms pass back through the ionizing region and, again, some of these are also ionized. Others escape and will fall onto the substrate to form the new film. When insulators need to be sputtered an rf supply is used. However, the rate of deposition is lower for rf and an alternative technique has been described by Sproul et al. (1983), Yashar et al. (1997) in which a pulsed 70 kHz positive voltage is superimposed on the negative dc supply for approximately 4 #s of the 14.2/zs cycle. This technique prevents charging and arcing when sputtering insulating materials. Another variation of the magnetron sputtering technique was first described by Window and Savvides (Window and Savvides 1986a and b, Savvides and Window 1986) and is known as the unbalanced magnetron. Here the magnetic fields are arranged so that some flux lines flow out towards the substrate. This can be achieved by using magnets of different area or strength (Fig. 9.7(b)), or, by applying an additional electromagnetic field using current-carrying coils (see Fig. 9.7(c)). Some electrons will now spiral along these field lines towards the substrate. Ions can be formed along the electron path and, also, because there is a net movement of negative charge from the magnetron region, positive ions will also be attracted to the substrate. The substrate will thus be bombarded by energetic ions and electrons, depending on its potential. A better practical system is to arrange for two (or four, or more) magnetrons to face each other in pairs having opposite magnetic polarities (Fig. 9.7(d)) and place the substrate in between them. Such systems, often described as closed field unbalanced magnetron sputtering systems (CFUBMS) have been described by Musil et al. (1990), Kadlec et al. (1990) and Teer (1990), who showed that there is a high probability of ionization of atoms in such configurations and that the substrate will be bombarded by ions of both the gas and film material. Under typical operating conditions, with chamber pressures of order 0.2-1 Pa, substrate ion bombardment currents of up to 100 mA at energies 10-100 eV are attained in such systems. Magnetron systems can give reasonable control of the energy-assisting contribution of depositing species and are capable of high rate deposition of the order of microns per minute. They do not allow independent control of these parameters however since the operating pressure influences the discharge conditions and, at higher pressures, collisions of particles with gaseous species can occur which change the energy of particles falling onto the substrate. More recent investigations by Ivanov et al. (1994) show that there is also additional bombardment of the substrate by particles reflected from the magnetron cathode. In spite of this, these authors have also shown that the presence of an external coil, which provides an additional magnetic field, can give reasonable control of deposition conditions. Improved adhesion of unbalanced magnetron sputtered coatings has been obtained by Miinz et al. (1989, 1991) using a preliminary arc bond sputtering process (ABSTM). In this system the coating process starts with a magnetron-assisted arc discharge where either Ar ions or a metal (Cr) ion first bombard the substrate at 800-1200 eV to give a

Physical Vapor Deposition

237

subsfrafe(floating or-V) I/Jat

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preliminary etch. The main deposition then continues in the CFUBMS mode producing a coating which, for Cr-nitride, withstands a critical load over twice the value for coatings produced without the arc etch stage (Hurkmans et al. 1997). When sputtering chemically active species a problem can arise as a result of target poisoning; this is known as the hysteresis effect. It is particularly important for sputtering of titanium to form titanium nitride films since the Ti itself combines with, and pumps,

238

J. S. Colligon

the nitrogen and the speed of this additional pumping depends on the degree of poisoning of the Ti target. The problem can be overcome by using a high-speed pump on the system, by pulsing the nitrogen gas supply, or by using an optical monitor tuned to a Ti emission line and a feedback loop to the gas flow control. In the latter case, any reduction in Ti can be counterbalanced and stoichiometric TiN will continue to form.

9.4. I N F L U E N C E OF E N E R G Y ON COATINGS

The additional energy and the sudden cooling of the atoms forming thin films during ion-assisted deposition has the following effects: (1) film morphology changes and film density increases, (2) residual gases are desorbed from the substrate surface, chemical bonds may be broken and defects created, and hence nucleation processes are modified and film stress and adhesion changes, (3) for films containing two or more elements, the film composition can be chosen to produce a range of novel metastable phases which have application as hard coatings and low friction films, and (4) film microstructure is altered.

9.4.1 Film morphology and density One of the earliest correlations between coating morphology and substrate temperature (or adatom diffusion energy) was made by Movchan and Demchishin (1969). This model was developed further by Thornton (1974, 1977) to include the effect of chamber pressure. In the usual deposition situation, the majority of the condensing atoms arrive from one direction at the substrate. Depending on their own energy and the substrate temperature (T) they will have an opportunity to diffuse across the substrate surface to find a nucleation site. Subsequent atoms may join the earlier ones, or, if their energy is low or the temperature is low, may start new nucleation sites. At low energy and temperature, inverted pyramidal structures will grow toward the source of condensing atoms with large voids between them, forming the so-called Zone I structure. This occurs when the ratio of substrate temperature to the melting point of the condensate (Tm) is between 0.1 and 0.3. As the substrate temperature rises some additional surface diffusion can occur and the coating becomes more dense. For higher T~ Tm the growth process becomes dominated by the adatom surface diffusion and columnar crystalline grains separated by dense boundaries form. At even higher T~ Tm values bulk diffusion can also play a part leading to the formation of equiaxed grains. The additional energy provided to the condensing atoms during ion-assisted deposition can produce similar growth zone regimes but at lower substrate temperature since this transferred energy provides the adatoms with the ability to move further across the surface and to adjacent sites in the bulk. This is immediately obvious from the increase in film densities reported; for example Martin et al. (1984) found an increase in ZrO2 density of 14% when additional 600 eV oxygen ion bombardment was used compared to the density

Physical Vapor Deposition

239

for films formed by evaporation of Zr without bombardment. Computer simulations of Mtiller (1989) have predicted the increase in density for this same system extremely well.

9.4.2 Nucleation and adhesion

A review of ion-induced adhesion effects by the present author shows that there are many examples of improved adhesion resulting from energetic ion irradiation of films (Colligon and Kheyrandish 1989). There is no clear evidence however as to the prime cause for this improvement which could be a result of atomic mixing across the interface, or enhanced surface mobility leading to better bonding. Certainly, cross-sectional microscopy studies of films show that additional energy provided by ion bombardment can reduce and, even, eliminate columnar-type growth as shown by Petrov et al. (1989). This group varied the bias step by step during deposition in a dc planar magnetron system when sputtering TiN onto HSS plates. A substrate bias greater than 120 V was sufficient to produce fully dense films for substrate temperatures of 300~ but a bias exceeding 240-280 V was required and substrate temperature of 900~ before columnar growth was interrupted. It thus appears that it is a combination of substrate temperature and added energy, both leading to increased surface mobility of atoms, which will influence the film nucleation and subsequent morphology. If the ion/atom ratio is rather high then, as Rossnagel et al. (1982) and Robinson and Rossnagel (1984) have shown, not only do the atoms migrate on the surface to form island clusters, but also these clusters are more stable to re-sputtering so that conical features can form. At the nucleation stage, both Choi et al. (1990) and Tsai and Atwater (1992) have demonstrated improved growth resulting from a combination of increased surface mobility and dissociation of islands during the initial stages of deposition of semiconductor materials. This trend supports earlier work of Marinov (1977) who found that energetic condensation produced larger clusters of atoms owing to the break-up of smaller clusters and the increased surface mobility of the re-liberated atoms. As films become thicker, however, volume recovery processes start to dominate over surface energy, as seen by Ziemann and Kay (1983). They show that fcc metal films can have (111) faces parallel to the surface initially but, as a result of strain energy being reduced to a minimum, the orientation will change towards (100). In studies of the growth of epitaxial TiN films on MgO (100) substrates, Hultman et al. (1988) conclude that the following factors are significant: (1) at low temperatures point defects are not annealed out and may form dislocation loops, (2) for low energy ion-assistance this defect density is reduced as a result of enhanced atom mobility; and (3) for higher ion energy, defects are created by the energetic particles at a higher rate than they are annihilated and remain in the film. The competition between the various processes leads to an interplay between substrate temperature and ion energy, with the added proviso, noted by Sundgren et al. (1992), that the ion energy should only be sufficient to cause atom movement without introducing defects or gas bubbles. The use of higher energies for ion-assistance can lead to the im-

240

J. S. Colligon

plantation of ions which may introduce additional stress, form bubbles, and, according to Ziemann and Kay (1982), migrate to grain boundaries and influence grain size.

9.4.3 Metastable phases Some of the earliest experiments on the control of stoichiometry of ion-assisted films were reported by Hentzell et al. (1985). Using a dual ion-beam system, A1 atoms sputtered by one ion-beam were deposited onto a substrate which was simultaneously bombarded by 100-500 eV nitrogen ions. Depending on the arrival ratios of nitrogen to aluminium species the resulting A1Nx compound was formed. The quantity x was seen to increase with N/A1 ratio up to unity; for higher ratios the compound was always A1N. Similar studies reported by Rossnagel and Cuomo (1988) for 100 eV oxygen bombardment of a growing film of copper (in this case evaporated) showed that, as the O/Cu ratio increased, apparent stationary composition regions, corresponding to Cu20, Cu205 and CuO were produced. Further X-ray analysis indicated that the Cu20 and CuO samples were normal polycrystalline phases and, apparently, the Cu205 was a new metastable form of copper oxide, hitherto unseen. The appearance of a similar metastable form of TiNx has been seen by Kim (1991) using a dual-beam system to sputter Ti onto a sample simultaneously bombarded by N2 ions at 100-400 eV.

9.4.4 Microstructure With the potential control in nucleation, growth and atom mobility provided by the additional energy available in ion-assisted deposition processes it is not too surprising that some control of crystal structure and grain size in the films is now possible. Hultman et al. (1988) have shown that the defect structure in TiN grown epitaxially onto MgO(100) substrates is a function of the bombarding energy and target temperature. This defect level decreases as the ion energy increases in the region 100-200 eV but then increases again at higher energies as a result of the increased damage rate. A similar trend, but for fixed substrate temperatures (about 50~ was reported by Kheyrandish et al. (1994). In this work the least damaged crystalline grains, judged from the TEM patterns, were obtained for 400 eV N + ion bombardment energies (i.e. also 200 eV per N atom). Roy et al. (1997) have proposed that the surface packing density of a material dictates the crystal phase that will develop during energetic condensation; for Ni deposition they found that it was the (220) texture that becomes dominant during ion-assisted growth. Chadwick and Miller (1967) showed that the surface packing density is 0.8 for a random Ni hard-sphere liquid model, 0.556 for the (110), 0.907 for (111) and 0.785 for (100) planes. Thus, assuming the deposition process normally gives the random dispersion, some energy is needed to displace atoms to the (110) distribution. There is thus a correlation between added energy and grain size which may be exploited to tailor the properties of the coating, since there is also a correlation between grain size and hardness from the Hall-Petch relationship (Hall 1951, Petch 1953). In general, for grains of size above 100 nm, this relationship shows that alloys with smaller grain size exhibit higher strength. The added energy per atom, discussed above, has been shown by Harper et al. (1985) to be a useful guide for the prediction of deposition conditions leading to improved film

Physical Vapor Deposition

241

properties. Harper et al. plotted the results of a large selection of data from various sources on a diagram showing ion/atom arrival ratio against ion energy and found that most data came within the interval 1-100 eV/atom. A similar diagram has been presented by Martinu et al. (1994) but with data for SiO2, other dielectrics, metals and semiconductors identified separately. These average values, calculated from all atoms and all ions, although a good guide, do not always give a good prediction of the experimental result. This is shown by Petrov et al. (1993) for deposition of Ti0.5A10.5N alloys in a UHV unbalanced magnetron system. Figure 9.8, based on Petrov's work, shows that preferred orientation, composition and lattice parameter for the films was quite different when the ion-assist energy changed, even if the energy per atom was the same. This reinforces the trends already discussed by Ziemann and Kay (1982) and Roy et al. (1988). Roy et al. (1988) measured the effect on stress produced in 5 # m copper films as a function of added energy using 62 eV, or, 600 eV Ar + ions and found a marked difference in the effects for the same added energy when using high and low energy ions. The major problem with the concept of average energy per atom is that (a) the condensing particle energy is not all delivered to the top plane of atoms, and (b) even if it were, not all the atoms in the plane share this energy. Clearly there is an urgent need for molecular dynamics simulations of the process but, unfortunately, for high ion energies, this requires lengthy computer time. Mtiller's molecular dynamics calculation, mentioned earlier for ZrO2, was also applied to Ni films during ion-assisted deposition (MUller 1986a, 1986b, 1987) and shows, for 50 eV Ar + ions at 30 ~ to the normal and an ion/atom ratio of 0.16 on a growing Ni film, that an almost fully dense film is produced. Attempts to obtain a better average energy (but not necessarily at the surface) have been made by Brighton and Hubler (1987), who considered that, to produce dense films, each component film atom must be involved in a collision cascade. Hirsch and Varga (1978) attempted a similar calculation based on movement of atoms in a region influenced by a thermal spike created by the energetic particle, and Grigorov et al. (1988) proposed that, for any ion-assisted deposition situation, there is an optimum flux of ions required to produce best film quality and this flux corresponds to that required to produce one displacement per condensing atom. The situation is even more complicated by the fact that the energy deposited at the surface of the film has been shown by Berg and Nender (1990) and Berg and Katardiev (private communication) to depend strongly on the mass and density of the substrate material. This effect is seen to be particularly important at the start of film growth (or, during the start of growth of each component in a multilayer). The ultimate model must therefore take into account, for deposition at a given substrate temperature: (1) the energy deposited at the uppermost (and, perhaps, second) layer of the growing film; (2) the effect of lower layers (which may have different mass and density) on this deposited energy; and

J. S. Colligon

242

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0.417 -

/

0.i,16 0.4~5

0

:

I 50

I

I 100

(Ed) (eV/mefa[ afom}

Figure 9.8. Effect of added energy on Ti-A1-N films (from Petrov et al. 1993).

(3) the number of atoms receiving an energy above some threshold level. This could be a threshold for breaking bonds to allow surface diffusion, or, for stimulation of some chemical reaction. In the meanwhile, the best practical approach is to aim to use a high ion/atom ratio and condensing ion and atom energies as low as possible so that defects are not created

Physical Vapor Deposition

243

deep in the film. Fortuitously this is exactly the conditions experienced at the substrate in most unbalanced magnetron systems.

9.5. A P P L I C A T I O N S OF P V D C O A T I N G S

PVD coatings are used for a wide range of applications which include modification of the following properties: Optical: for laser optics, compact discs, anti-reflective coatings (for spectacles), optically absorbing coatings, mirrors, architectural glazing, selective solar absorbers. Chemical: corrosion-resistant coatings, catalytic coatings, sensor electrodes. Mechanical: hard coatings for cutting tools, low friction coatings, wear-resistant coatings. Electrical: electrical contacts and interconnects, solid-state devices, solar cells. Decorative: to provide attractive finish on many products.

Looking at the coating industry as a whole it is seen that PVD is not the main coating process and, when all coating methods are considered such as painting and electroplating, it is estimated that PVD represents about 4% of the total coating market (Gorham Newsletter). However this is 4% of a market in excess of 100 billion US$ per annum and still represents a large turnover. In addition, the ion-assisted processes are able to produce dense, well-bonded films and the process can run at relatively low substrate temperatures so that it allows coatings to be applied to materials with low melting points such as plastics. Perhaps the most important advantage of energy-assisted PVD is the ability to create novel phases of binary and ternary compounds and to produce materials with novel microstructural (and nanocrystalline) features. These unique film properties arise as a result of the rapid "cooling" of the fast condensing atoms forming the film which can approach 1014 K/s. Some typical examples of these special features in coatings with novel optical, chemical and mechanical properties are now described. Coatings applied for their electrical and electronic properties are normally specific to a certain small-scale requirement, such as mask repair, contacts, insulating layers and barrier layers and ion-beam methods are often employed. Coatings for decorative purposes have all the same requirements of durability and corrosion resistance as those described below but, because very large areas are to be treated (many square metres for glass and plastic roll-coating systems) a compromise has to be reached between the high throughput requirement and the controlled conditions described below for optimum quality of the coated product. In these systems ion-plating, arc ion-plating and magnetron sputtering are normally used.

9.5.10pticalcoatings

Methods for producing optical coatings have been reviewed by several authors (Martin and Netterfield 1986, Colligon 1992, Martin 1986, Gibson 1987) and it is clear that coating stability and performance is enhanced if they are fully dense and defect-free. Optimum conditions for forming some of the main optical materials, compiled from data in references (Allen 1983, McNeil et al. 1984a, Martin et al. 1983, McNeil et al. 1984b, McNally

244

J. S. C o l l i g o n

Table 9.3 Energy-assisted thin film optical coating conditions (measurement at 550 nm except where indicated). Refraction indix SiO2

1.46

TiO2

2.5

ZrO2

2.05

Ta205

2.25 (350 nm) 1.68 (350 nm) 1.8 (351 nm) 2.4 2.11 (350 nm) 1.39

A1203 Y203 CeO2 HfO2 MgF2

Ion

Energy(eV)

Substrate temperature(~

Ref.

O 30 Ar (oxygen background) O 300

Various

Ar 600 (oxygen background) O 500 (oxygen background) O 300-1000

20

McNeil et al. (1984) McNeil et al. (1985) Martin et al. (1983) Allen et al. (1983), McNeil et al. (1984), McNeil et al. (1985), Martin et al. (1983)

275

McNally et al. (1985)

275

McNally et al. (1985)

< 80

Cole et al. (1984)

20 275

Netterfield et al. (1985) McNally et al. (1985)

O 300 (oxygen background) O 1200 O 500 H2O

700

200

20

Martin and Netterfield (1985)

e t al. 1985, Cole e t al. 1984, Netterfield e t al. 1985 and Martin and Netterfield 1985) are presented in Table 9.3. One of the problems with the energy-assisted optical coating is that, if the ion/atom ratio is too high, or the ion energy is high, there is significant re-sputtering of the film and this can selectively remove one of the components. In addition the ion species itself can become buried in the film. Figure 9.9 illustrates this problem for CeO2 films produced using 1.2 keV oxygen ion-assistance. As the oxygen ion current density is increased the films have a higher refractive index and are more stable to exposure to air. However, beyond a critical current density value, the refractive index falls again owing to either oxygen incorporation, or, preferential removal of Ce. Proper selection of process parameters is thus critical for optical films. The ion-assist technique not only produces superior quality optical coatings, but also, presents a new opportunity to "tailor" the films by sputtering two or more materials at the same time. Figure 9.10 shows data by Cevro (1995) for a single layer anti-reflective coating for 1050 nm wavelength radiation. This was, in fact, produced by ion-beam sputtering of a Ta205/SiO2 mixed layer but, with the geometry used, some reflected ions also bombarded the growing films. Similar data has been reported by Laird and Belkind (1992) and Vretenar (1990) for TiO2/SiO2 films, by Denetale and Harker (1987) for TiO2/A1203 and by several groups (Pond e t al. 1989, Russak e t al. 1989, Gilmore e t al. 1987) for ZrO2/SiO2 films. Further

245

Physical Vapor Deposition 2.5

j

2.t~

I

AD CeOz

2.3 x

2.2

.c_ ~

2.1

r~

2.0

_> D._

~

1.9 1.8 I

films not stable

I

I

1

I

250 500 750 1000 1250 Oz ion current density (pA cm-2) ,

pgssible incorporation w

or oxygen or voua creation

Figure 9.9.

Variation in refractive index of CeO2 a a function of oxygen ion-assist energy. Films marked (x) are not stable on exposure to air. (From work of Netterfield et al. (1985)).

controlled studies of this type of new optical coating should lead to many new composite films with unique properties.

9.5.2 Corrosion protective coatings Another area where film density is important is that of corrosion protection. Ion-beam methods for producing such coatings have been reviewed by Wolf (1992) who notes that stress is an important factor in this type of application. Tensile stress can be reduced by increasing the ion/atom ratio in many cases, for example in Cr films using Ar ion assisted deposition (Roy et al. 1989a, b). Again the novelty of energy-assisted PVD is that new combinations of materials can be studied and their corrosion properties evaluated. Work reported by Hsieh et al. (1992), evaporating Nb and Cr with 250 eV Ar + ion-assistance at an ion/atom ratio of 0.68, has shown that Nb-50%Cr films are the most dense and column free but, because of its ability to form a Nb-rich oxide, Nb-30%Cr films had the best corrosion resistance. The present author has produced a range of Fe-Cr alloys for experiments which showed that the Cr concentration had to exceed 16% before complete passivation occurred (Fujimoto et al. 1993). The ability to "seed-in" additional elements has been exploited by Donahue et al. (1997) who showed that the addition of yttrium to magnetron-sputtered films of Ti-A1Cr-N drastically reduced the oxidation rate at elevated temperature.

246

J. S. Colligon

-

a~ ./Taz

0 s film

2.2

2.11 "0 e-

>

....

1.8

t.l

tU

"

1.6

1.~

x,,,~.,.,~,...~~ Si 0 z f i l m I , , , 250

Ta 1~%1 Si196%] I

I I 500

I

I

wavelength

I

I I 750

i

,

,

1 1000

-

(nm)

Figure 9.10. Refractiveindex of mixed SiO2 "Ta205 filmsas a function of wavelength (from work of Cevro (1995)).

9.5.3 Hard coatings Some of the main coatings used to produce durable hard layers include the nitrides of Ti, Zr, Hf and Cr, more complex nitrides which include Ti-A1-N, Ti-Zr-N, cubic BN, Ti-Nb-N, Ti-A1-O-N, diamond and diamond-like carbon, carbides of Ti, W and Ta and oxides. Hard coatings based on carbon, often referred to as diamond-like carbon (DLC) coatings, have been investigated using ion-assist methods since the early work of Aisenberg and Chabot (1971) who noted that diamond-like films were formed by directing a 40 eV beam of positive carbon ions at a substrate. Weissmantel (1981) later showed that similar films could be formed using ion-assisted techniques. More recently He et al. (1994) have shown that hard films with a high sp 3 content can be formed by sputtering graphite with Ar + ions onto silicon or AISI 52100 bearing steel. The growing film is bombarded with ions formed by introducing CH4 into a second source and running at energies from 0.2 to 25 keV. The highest sp 3 content occurred at the lowest (200 eV) energy where a Knoop hardness of 5100 kgf/mm 2 was registered using a 20 g load. A similar energy requirement was seen by Yugo (1994) in a microwave plasma running in a CH4/H2 mixture but here a two-step process was required, first a high level (40%) of CH4 was used to develop a carbon cluster and then, after 30 minutes, the CH4 level was reduced to only 0.5% allowing production of sp 3 bonds. This extensive field of DLC coatings has been reviewed by Ensinger (1996) who, in other work (Ensinger and Wolf 1989), has shown that these films provide good corrosion protection when applied to alloy steels. Silicon-carbide coatings have also attracted con-

247

Physical Vapor Deposition

1.2 O

-

o m

to C O

0.8

x

C:Si

~ ~ . .

~, A r ' S ( j X

v

"G 0.~ 0 a. E 0

0

, . . . . .

-4 . . . . .

0

100

-r . . . .

200

~mbard~ng ion energy

't-300 (eV)

~00

Figure 9.11. Compositionof ion-assisted Si-C films as a function of CH4 ion energy (from work of He et al. (1995)).

siderable interest (Learn and Haq 1970, Seaward et al. 1986, Yamada 1989, Berezhinskii et al. 1990, Yen and Murray 1990, He et al. 1995). Figure 9.11 shows data by He et al (1995) for dual beam ion-assisted growth using CH4 gas in the ion-assist source and Ar + ions to sputter silicon onto the substrate. It can be seen that the ion-assist energy must exceed 300 eV before the SiC compound forms. At this energy there was considerable re-sputtering of material and hence a low net rate of film growth. Cubic boron-nitride (c-BN) is an extremely useful coating for industrial applications owing to its high hardness, chemical and thermal stability at extremely high temperatures and low reactivity with ferrous alloys. It also has a large band-gap (>6 eV) and can be doped p-, or, n-type. Unfortunately, to form c-BN using conventional methods, pressures of over 11.5 GPa and temperatures in excess of 2000 K are required, which would destroy the properties of most substrates. However, ion-assist methods allow these films to be formed in vacuum at 460~ (see, for example Park et al. (1997)). In all cases a mixture of argon and nitrogen ions is required with an energy of about 300 eV and an ion/atom ratio of about 1.5. According to a model of the process by McKenzie (1993) and simulation experiments by Dworschak et al. (1995) the compressive stress, induced by the ions colliding with the film atoms, and the formation of interstitials is assumed to trigger the conversion of hexagonal-BN to c-BN. One of the most interesting aspects of current PVD methods is their ability to produce coatings with grain sizes smaller than those previously possible. This offers the possibility of producing nanoscale composite metastable alloys and pure materials which have unique properties including high hardness and high ductility. When the dimensions of components of the film become less than 100 nm the term/3 in the Hall-Petch equation

H = Ho + fld -1/2

248

J. S. Colligon

where H is the hardness, H0 a material constant and d the grain size, no longer applies since, according to Carsley et al. (1995) the material now has a significant fraction of its atoms in a grain boundary phase and so must be considered as a two-phase material. By considering typical grain boundary widths and grain sizes the fraction of each phase can be estimated and the overall flow stress of the composite material calculated, using bulk values of flow stress for the grains and amorphous material values for the boundary regions. The predicted values for hardness were compared by Carsley et al. with experimental data of Hughes et al. (1986) and E1 Sherik et al. (1992) for nickel and show good agreement. Similar agreement was also seen for iron and also copper. The maximum hardness is predicted to occur for grain sizes of about 4 nm for these materials. Vep~ek et al. (1995, 1996), Vep~ek and Reiprich (1996) and Christiansen et al. (1998) have predicted that mixed phase nitride materials should produce hard coatings and have demonstrated significant increases in hardness for crystalline TiN in amorphous Si3N4 formed by plasma-assisted CVD. Maximum hardness of 50-70 GPa was found for grain sizes of order 6-7 nm. Other combinations of nitrides are predicted to exhibit similar high values of hardness and Musil and Vlcek (1998) have noted that there are two types of mixed-nitride coatings: nanocrystalline metal nitride/hard matrix (such as c-TiN/a-Si3N4) and nanocrystalline metal nitride/soft matrix (such as c-TiN/Ni, Cu, Co, Ag, Au, rare earth elements). The metal nitride can be of several forms, including TiN, A1N, ZrN, NbN and WN. The former types of mixed nitride coatings are very hard but are brittle whereas the latter are of similar hardness but are ductile (Musil and Regent 1998, Musil and Zeman 1998, Nakayama et al. 1994). Musil et al. have shown that these mixed nanocomposite coatings can be formed using magnetron sputtering if the substrate is heated sufficiently to allow the second, amorphous phase, to migrate to the grain boundaries of the first phase (Musil and Vlcek 1998, Nakayama et al. 1994). It is this segregation that is believed to be responsible for limiting the growth in grain size. The value of substrate temperature required for formation of nanocomposite ZrN/Cu was found by Musil and Vlcek (1998), Musil and Regent (1998) to be of order 600~ and it is hoped that this can be reduced by improving the technique, for example, by incorporating a small quantity of a third element such as yttrium, which is known to reduce the grain size in metals (Liu et al. 1997). Another method for producing hard coatings is to use multilayer films A-BA-B . . . . , where the Young's modulus of one material is lower than that of the other. Seitz and Koehler (1956) predicted that the hardness would also increase for these "superlattice" structures and this prediction has been confirmed by Helmersson et al. (1987), for TiN/VN multilayers, by Kitagawa et al. (1998) for TiN/Ti multilayers and by Wu et al. (1997) for CNx/ZrN multilayers. 9.6. F U T U R E

TRENDS

The energy-assisted deposition processes are thus now producing a range of novel coatings which are well-bonded and can increase the lifetime of the coated product by significant factors. Coatings are formed mainly by magnetron sputtering, arc-ion-plating or other plasma-based techniques and it is expected that these methods will further improve to provide more efficient operation (for example, better utilisation of targets in magnetron systems) and better control of deposition parameters (for example addition of

Physical Vapor D e p o s i t i o n

249

electromagnetic coils to control the degree of ionization in magnetron sputtering). However the main further development appears to be in the production of new metastable alloy coatings (Bates and Arnell 1994, Monaghan and Arnell 1991) and nanocomposite coatings (Hughes et al. 1986, E1-Sherik et al. 1992, Vep~ek et al. 1995, 1996, Vep~ek and Reiprich 1996, Christiansen 1998, Musil and Vlcek 1998, Musil and Regent 1998, Musil and Zeman 1998, Nakayama et al. 1994) According to Musil and Vlcek (1998), ions bombarding a surface during coating impart sufficient energy by momentum transfer to allow high temperature phases to form within the collision cascade. The surrounding atoms however provide fast cooling at rates of order 1014 K/s so that the high temperature phase is frozen into the solid. Thus the presence of (a) an additional small quantity of several additional elements in the film, (b) additional energy to provide the effective temperature and pressure conditions within the film to form the high temperature phase, and (c) high cooling rates, will allow many phases to be formed at relatively low temperatures. Ion-assisted coatings can provide all the above parameters and the immediate work needs to study what types of coating can be formed and how stable they are in working conditions. The possible control of composition, crystal structure and stress in the coatings will lead to some exciting new possibilities in the near future, perhaps even the formation of the theoretical hard beta-phase cubic C3N4, predicted by Liu and Cohen (1989), Cohen (1994) and Yu et al. (1994) and currently occupying the attention of many research groups (Sj6str6m et al. 1996, Vep~ek et al. 1995, Riviere et al. 1995, Boyd et al. 1995).

ACKNOWLEDGMENT

The author wishes to thank many colleagues for providing information for this review.

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Chapter 10 Chemical Vapor Deposition 10.1. Introduction 10.1.1 Presentation 10.2. Gas-Phase Transport and Reactivity 10.2.1 Non-Reactive Fluid Flow 10.2.1.1 Geometrical Effects 10.2.1.2 Influence of Operating Conditions 10.2.2 Reactive Flows 10.2.2.1 Thermodynamic Approach 10.2.2.2 Gas-Phase Mechanisms and Kinetics 10.2.2.3 Simulations of the Gas-Phase Composition 10.2.2.4 Nucleation and Growth of Solid Particles from the Gas Phase 10.3. Solid Phase Formation 10.3.1 Solid-Gas Thermodynamic Equilibrium Approach 10.3.2 The Structure of CVD Layers 10.3.3 The Driving Force for Crystal Growth from the Vapor Phase 10.3.4 Surface Mechanisms and Kinetics 10.3.4.1 Rate of Elementary Surface Processes 10.3.4.2 Estimation of the Kinetics of Surface Processes in CVD 10.3.5 Heterogeneous Nucleation 10.3.5.1 Nucleation Theories 10.3.5.2 Two- and Three-Dimensional Nucleation 10.3.5.3 Epitaxial Nucleation and Growth 10.3.5.4 Multi-component nucleation and chemical reactions 10.3.5.5 Nucleation Enhancement 10.3.5.6 Selective Vapor Deposition 10.3.5.7 Theoretical and Experimental Studies of Nucleation in CVD 10.3.6 Crystal Growth 10.3.6.1 Mechanisms of Crystal Growth 10.3.6.2 Classification of Crystal Faces 10.3.6.3 Crystalline Morphology 10.3.6.4 Morphology and Normal Growth Rate of Crystal Faces 10.3.6.5 Relation between Preferred Orientation and Morphology of Polycrystalline Films 10.3.6.6 Theoretical Studies of Crystal Growth from the Vapor Phase 10.4. Conclusions References

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Chapter 10 Chemical Vapor Deposition E TEYSSANDIER and A. DOLLET

10.1. I N T R O D U C T I O N

In vapor deposition, a solid material is deposited from gaseous molecules onto a substrate to form a coating. The process is nowadays classified into two basic families according to the initial proposition of Blocher (1967): physical vapor deposition (PVD) when coatings are prepared from pure condensation processes, and chemical vapor deposition (CVD) when deposits are formed by chemical reactions only. These two routes were developed separately and many modifications have since been introduced, so that this distinction is no longer strictly valid. A great variety of processes related to CVD now exist. Some were developed to lower the deposition temperature by introducing (i) a glow discharge (plasma-activated CVD, PACVD, or plasma-enhanced CVD, PECVD) or (ii) less thermally stable organometallic precursors (organometallic CVD, OMCVD). Some were developed on the basis of a specific devicewa laser beam can be used, either for gas phase excitation or substrate heating in the laser-enhanced CVD process (LECVD). Certain denominations refer to specific types of coatings and related processes, such as epitaxially grown films for electronic or optoelectronic applications by the metalorganic vapor phase epitaxy process (MOVPE), or infiltration of ceramic materials in fibrous preforms by the chemical vapor infiltration (CVI) process. Distinction is usually made between two basic types of reactors: hot-wall and coldwall reactors (Fig. 10.1). In cold-wall reactors, the parts to be coated are heated by various means: electromagnetic induction or joule effect for metallic parts; tungsten-halogen lamps (RTPCVD, rapid thermal processing CVD) for non-conductive materials. In the hot-wall case, the whole reactor, usually tube-shaped, is heated by means of a classical furnace. Though the same phenomena take place in both cases, they present characteristic features: deposition can occur at hot walls thus leading to gas phase depletion, and the hot part in a cold-wall reactor is responsible for large thermal gradients and significant buoyancy-driven convection. These points will be further detailed. A lot of industrial coatings are currently produced by the CVD process, in the field of electronic, optical, optoelectronic or mechanical applications, and a considerable body of literature as well as scientific journals are devoted to CVD. The reader can also refer to the non-exhaustive review articles already published on the subject (Bryant 1977, Hess et al. 1985, Jasinski et al. 1987, Jensen 1987, Besmann et al. 1989, Jensen 1989b, Wahl 1989, Kleijn 1991, Hampden-Smith and Kodas 1995). The present review will report on 257

F. Teyssandier and A. Dollet

258

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the thermally-activated process that excludes any kind of physical activation either in the gas phase or on the surface. The basic mechanisms involved in the CVD process will be reviewed and illustrated.

10.1.1 Presentation Difficulties in obtaining a comprehensive view of the growth process arise from the number of interdependent parameters such as temperature, pressure, composition of the initial gas phase, flow rate, type of reacting chamber, etc. and the complexity of the various steps. The steps of the individual mechanisms involved, which are summarized in Fig. 10.2, can be divided into two highly interdependent parts: (a) phenomena occurring in the gas phase, such as bulk transport of the species in a hydrodynamic flow with diffusion mechanisms due to concentration gradients and possibly homogeneous reactions or gravity effects; in a cold wall reactor one must furthermore take into account thermal diffusion (Soret effect), and (b) the surface phenomena including adsorption, diffusion, reaction and desorption. In this case a distinction has to be made between the nucleation period when interactions with the substrate take place, and the steady-state growth of the coating from the gas phase.

259

Chemical Vapor Deposition

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Figure 10.2. Basic mechanisms in the chemical vapor deposition process.

The range of variation of the different parameters is as follows. In the thermally activated process, the temperature domain usually ranges from about 100 up to 2000~ The lowest deposition temperatures can be reached only by the use of organometallic precursors, since inorganic precursors, which are much more thermally stable, require at least 800~ to be decomposed. The total pressure in the reactor extends from atmospheric pressure down to a few tenths of a Torr (1 Tort = 133 Pa), so, according to the large characteristic dimensions of the reactors a continuum description is generally appropriate. Furthermore, since the gas velocities are low (Reynolds number < 100) and the Mach number is small (Ma < 1) the flows are laminar and the gas is assumed to be incompressible. The wide variety of relevant scientific domains is one of the difficulties of the CVD process and the various above-mentioned phenomena will be detailed in this review. The presentation order corresponds to the occurrence of the phenomena, from the introduction of the reactive mixture at the reactor entrance to the growth on the surface to be coated.

10.2. G A S - P H A S E T R A N S P O R T AND R E A C T I V I T Y

In CVD, the various chemical elements required to grow the film are transported by molecules in a forced flow from the entrance of the reactor towards the hot surface. The multicomponent, diffusive and reactive fluid flow can be an extremely complex system, and both flow phenomena and homogeneous chemical reactions affect the film growth. The fundamental equations describing the transport properties in continuous media are well known (Bird et al. 1960) and they will not be reviewed here. For a comprehensive review of transport phenomena in CVD see Holstein (1992), Jensen (1994). We will first focus on the fluid flow, ignoring its potential chemical reactivity which will be subsequently discussed.

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10.2.1 Non-reactive fluid flow 10.2.1.1 Geometrical effects The reactor has a critical influence on the uniformity of the fluid flow. The direction of the flow with respect to the gravitational field as well as the reactor geometry induce specific flow patterns such as recirculation cells. Under mass-transfer-controlled conditions the local flux of active species and hence the thickness uniformity and structural homogeneity of the coating is highly dependent on the flow pattern. Transport phenomena were described for a long time by simple analytical models based on boundary-layer similarity and film theory. The validity of the boundary layer hypothesis is questionable in many CVD reactors (Rosenberger 1987) and the absence of such a layer has been demonstrated in several cases by 3D numerical modeling since its thickness eventually exceeds the dimensions of the reactor for typical operating conditions. A stagnant layer theory was also encountered in the CVD field. This non-physical theory, which allowed decoupling convective and diffusive phenomena in several CVD analytical models, originates from flow visualization experiments. Eversteyn et al. (1970) interpreted a TiO2 particle free layer adjacent to the susceptor as evidence for a stagnant layer. Laser Doppler velocimetry experiments (Talbot et al. 1980), as well as modeling (Fotiadis and Jensen 1990) have revealed that thermophoretic forces were responsible for driving the seed particles away from the substrate. None of the above-mentioned approaches gives a satisfactory description of the physical processes involved in CVD, but they have provided simple models which have successfully predicted qualitative trends. The same film theories have also been developed for heat transfer (thermal boundary-layer), and for mass transfer (diffusion boundarylayer) (Hitchman 1980). All these theories are now encompassed in the approach which solves the coupled partial differential equations by numerical integration using finitedifference, finite-volume or finite-element methods. 1D to 3D simulation models are required depending on the type of reactor encountered. Both vertical and horizontal reactors are used. In the absence of gravitational effects, flow patterns are equivalent. This is not the case in CVD reactors on Earth, where natural convection resulting from essentially thermally induced density gradients is superposed on forced convection. The influence of natural convection on the process may be quite different according to the disposition of the reactor orientation. Vertical axisymmetric reactors such as impinging jet, stagnation point flow or rotating disk reactors have been widely investigated on the laboratory scale. The ideal stagnation point and rotating disk flows have uniform vertical gas velocities producing a boundary layer of constant thickness in which the gases move radially outwards. Due to the axisymmetric reactor geometry, 1D simulation provides a reliable description of the flow phenomena which agrees very well with flow visualization experiments (Wahl 1977, Fotiadis et al. 1987, Wang et al. 1995). Forced and natural convection in vertical reactors are coaxial phenomena. They have the same direction in the upward forced-flow case which stabilizes the axial density gradient. Recirculations may however be present in this case because of radial temperature gradients. Downward non-isothermal forced-flow results in counter-current influences that usually generate strong recirculation cells (Wahl

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261

1977). These cells are often eliminated by increasing the inlet flow-rate thus improving the thickness uniformity. Flow in horizontal reactors is much more complex, and any horizontal density gradient in a gravitational field results in buoyancy-driven free convection without threshold (Rosenberger 1987, Gokoglu et al. 1989). Both longitudinal and transverse rolls are then generated and 3D modeling is in this case required (Ouazzani and Rosenberger 1990). Under mass-transfer-controlled conditions the thickness uniformity is highly dependent on the flow pattern. The deposition profiles calculated in Fig. 10.3 (Jensen 1989a, Moffat and Jensen 1986) for two different flow patterns show that growth rate is enhanced or reduced respectively in the domains facing downward or upward secondary flow directions. The same phenomena can take place at the edge of substrates, frequently leading to their rounding off. Mass-transfer inside a reactor depends on both the reactor enclosure and the operating conditions. The latter can be modified from one experiment to another, but the reactor geometry is selected once for all. The design of reactors can be optimized by modeling a priori the flow and heat transfer. These calculations are invaluable in obtaining a better understanding of their behavior for various operating conditions. Several basic options can thus be checked: the shape, distance between inlet and susceptor, inlet tube diameter, introduction of baffles. The influence of these reactor features has been extensively illustrated in the case of a cold-wall vertical reactor (Fotiadis and Jensen 1990).

10.2.1.2 Influence of operating conditions Once the reactor has been selected, the experimenter has to determine the most appropriate operating factors. Several parameters can be chosen and these include temperature, total pressure, nature and proportion of precursors, and total flow rate. The nature and shape of the parts to be coated are the imposed parameters. Our understanding of the basic phenomena involved in the CVD process provides us guidance for selecting or optimizing the operating parameters. Analytical models were first used for that purpose. The influence on the growth rate of the basic parameters such as total pressure or temperature can be easily schematized and therefore memorized thanks to fairly simple models, such as the one proposed by Van den Brekel (1977) which successfully predicts qualitative trends. To quantify the relative influence of gas phase diffusion resistance transfer against the surface reactions kinetic limitation, he introduced the so-called CVD number, which is in fact a simplification of the Damk6hler (surface) number when only first-order surface reactions are considered (Hess et al. 1985). The general behavior of a CVD reactor can also be deduced from dimensionless groups that arise from the scaling of the governing transport equations (Jensen 1989). These groups are also helpful in determining the influence of a specific parameter, such as temperature, on the physical properties of gases and thus the transport properties in CVD reactors (Rosenberger 1987). 10.2.2 Reactive flows In many situations, the residence time and gas temperature are high enough to allow reactions to occur in the gas phase prior to reaching the surface. The new species so

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Figure 10.3. Flow profiles and corresponding growth rate of GaAs calculated with a 3D simulation including a one species (TMGa) diffusion model with a fast surface reaction (presented in Jensen 1989a after Moffat and Jensen 1986). Left: insulation of side walls, right: cold side walls.

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263

produced are transported towards the substrate at variable rates according to their physical properties. Therefore, a large variety of species having different reactivities can be present in the vicinity of the deposition surface. The main CVD processing domains are traditionally reported in terms of growth rates and their dependence on temperature. Surface kinetics or gas phase mass-transfer control operating domains are thus deduced from the overall activation energy. To determine the apparent activation energy, it is usual practice to plot the logarithm of the growth rate against 1/T. Three temperature domains can thus be observed: (a) in the low-temperature domain, the deposition rate increases rapidly with temperature. The steep slope which is indicative of a high overall activation energy is associated with a growth rate governed by the surface phenomena. For very short residence times and/or very low temperature of the gas mixture, almost no intermediate species are formed, and the flow is assumed to be chemically frozen. (b) at intermediate temperatures, surface kinetics become fast and the ability of the process to supply reactive species to the surface becomes the growth-limiting factor. In this range of temperature, convection in the bulk and diffusion in the vicinity of the surface become the rate-limiting transport phenomena. Their weak dependence on temperature is responsible for the low overall activation energy. Higher gas phase temperatures allow homogeneous chemical reactions to develop under kinetic control. Near-equilibrium conditions are reached in the gas phase for long residence times and/or high temperatures. (c) at high temperatures, a decrease in the growth rate is often encountered. Several phenomena may account for such a behavior: the nucleation and growth of particles in the gas phase that deplete the amount of reactive species on the surface, modifications of the fluid flow dynamics or desorption of species that exceed the number of molecules impacting on the surface. In the intermediate and high temperature domains (cases (b) and (c)) gas phase phenomena and especially mass transport phenomena are the factors that control growth rate uniformity. The transition from the surface kinetics rate limiting domain to the mass transport limiting domain is of prime importance in the process as it leads to very specific features of the growth mechanisms. Gas-phase transport-limiting process is often responsible for surface roughening, whereas surface kinetics limitation tends to smooth out the surface (Van den Brekel 1978, Palmer and Gordon 1988, Palmer and Gordon 1989) (see also crystal growth section in part 3).

10.2.2.1 Thermodynamic approach Because the detailed reaction mechanisms and corresponding kinetic data are rarely available, chemical equilibrium has been the first traditional approach to CVD process modeling. Thermodynamic calculation has been applied to a great variety of chemical systems and for various purposes which will be detailed in the "Solid-gas thermodynamic equilibrium approach" section in part 3. Application of thermodynamic calculations to gas phase in the absence of solid state materials leads to pseudo-thermodynamic equilibrium involving a supersaturated gas phase. This allows determining bulk compositions and identifying key species, as well as conditions under which they might form (Allendorf 1993). Though this equilibrium situation, which is representative of a gas phase composition that would be reached at infinite residence time by a reactive mixture flowing at constant temperature (as long as homogeneous nucleation does not take place), is never

264

F. Teyssandier a n d A. D o l l e t

encountered in CVD reactors, near- equilibrium states are often approached in hot-wall reactors. Accurate modeling and, ultimately, control of the CVD process requires knowledge of both the identity of the species formed during gas phase transport and their rates of formation. When these are known, the flux of reactive species toward the surface can be predicted. For that purpose, gas phase reaction mechanisms have to be identified and their corresponding rate constants must be measured or estimated.

10.2.2.2 Gas-phase mechanisms and kinetics In the dilute medium encountered in the CVD process, it is assumed that there are only two types of kinetic processes: unimolecular bond dissociation reactions and bimolecular reactions. For further details on experimental and theoretical development of gas-phase kinetics the reader can refer to classic texts (Benson 1976, Forst 1973, Gilbert and Smith 1990, Robinson and Holbrook 1996). Only a brief review of the basic principles will be given here. In a unimolecular reaction, an isolated molecule undergoes a chemical change. Reaching an internal energy sufficient for the reaction to occur requires excitation, which most frequently is obtained from collision with a bath gas at thermal equilibrium. Unimolecular reaction rates can be obtained experimentally from static (Neudorfl et al. 1980) and flow pyrolysis (Ring et al. 1970), shock tubes (Mick et al. 1996) and laser powered (Jasinski and Estes 1985) homogeneous pyrolysis techniques. They may also be estimated from theoretical models. The energy-dependent rate constant k ( E ) can be computed from the RRKM model (Rice-Ramsperger-Kassel-Marcus) which is the application of the transition state theory to a microcanonical ensemble of excited reactant molecules. This model gives an exact upper bound to the rate coefficient. A simpler model (QRRK) (Westmoreland et al. 1986) in which a molecule is treated as if all its vibrations can be represented by a simple average frequency (usually the geometric mean) can be used instead of the RRKM model. Three categories of bimolecular reactions (A + B --+ products) may be distinguished: displacement, association and metathesis reactions (in which an atom is transferred from A to B). The bimolecular adaptation of the unimolecular QRRK model (Dean 1985, Westmoreland et al. 1986) can be used to predict the pressure and temperature dependence of addition or recombination reactions that initially form excited intermediates. Once the various gas phase reaction mechanisms are identified and their rate measured or estimated, the kinetic data can be introduced in full-scale models or simpler (plug-flow or well-stirred) time-dependent models such as the Sandia CHEMKIN/SENKIN package (Kee et al. 1989, Lutz et al. 1995). These tools are useful in interpreting the evolution of gas mixtures (Teyssandier and Allendorf 1998).

10.2.2.3 Simulations of the gas-phase composition Even the simplest chemical systems may involve numerous gaseous species. The model proposed by Coltrin et al. (1989) for silicon deposition from silane, which includes only two elements (Si and H), is composed of 27 chemical reactions involving 17 gas phase chemical species. Adding carbon to allow SiC deposition, increases the gas phase chemical model up to 83 reactions involving 41 species (Allendorf and Kee 1991). Some of these species are not important in the deposition process and such systems can be simpli-

Chemical Vapor Deposition

265

fled for specific total pressure and temperature conditions (Wang et al. 1994). However the rejection of some species or chemical reactions according to concentration criteria can be misleading. As a fact, the selected subsystem must constitute a closed system with respect to the chemical reactions considered, i.e., the important reaction pathways for the selected species must not be ignored. Furthermore, due to the large difference of surface reactivity between closed-shell molecules and radicals, the latter species, even when present at low concentration, may have a significant influence in the surface process. For instance, though Sill2 forms only at high temperatures and is thus usually present at the interface with the surface substrate in far lower concentrations than the precursors introduced at the reactor entrance (silane or disilane), its unity reactive sticking coefficient (RSC) at the surface is at least three orders of magnitude larger at 1000 K than those of the saturated silanes (Sill4, Si2H6, Si3Hs, etc.). The contribution of silylene to silicon growth rate (at least at high temperatures) can be for that reason significant. Full-scale modeling including gas-phase reactions allows determining the contribution of the various gaseous species to the deposition rate. The partial deposition rates of Sicarrier species at the surface facing a circular impinging jet (Fig. 10.4) reveals that (i) the deposition rate increases rapidly towards the edge of the impingement surface as a result of the increasing amount of highly reactive intermediate species formed, (ii) the contribution of the initial precursor Si2H6 is constant on the surface that reflects its supersaturation which is a consequence of its low surface reactivity, and (iii) Si3H8 is responsible for the step variation of the deposit thickness observed at the edge of the surface.

10.2.2.4 Nucleation and growth of solid particles from the gas phase The non-controlled introduction of solid particles in a CVD reactor is detrimental to the properties of the films. These particles are usually introduced into the reactor as impurities or can be formed by agglomeration of smaller particles nucleated in the gas phase. They act as extrinsic seeds for heterogeneous nucleation, deplete the species necessary for the heterogeneous reactions and create undesirable microstructures including porosity that lower the mechanical properties of the deposits. The elementary processes involved in the formation and growth of particles from gaseous precursors include gas phase chemical reactions, homogeneous nucleation, heterogeneous condensation, coagulation and coalescence or fusion (Fig. 10.5). Both high temperatures and long residence times are favorable to homogeneous nucleation of such particles. This situation is especially encountered at the center of recirculation cells (infinite residence time) in hot-wall reactors. In contrast, the risk of solid particle formation is considerably reduced under low total pressure and/or in diluted reactive mixture. Because of the large temperature gradients found in cold-wall CVD reactors, the thermophoresis forces may cause the solid particles to deviate significantly from the fluid particle trajectories. This phenomenon is the preponderant mass-transfer mechanism used in the modified chemical vapor deposition process (MCVD) to manufacture optical waveguide preforms (Kim and Pratsinis 1988, Lin et al. 1990). Since the thermophoretic coefficient increases with decreasing pressure, its influence is expected to be enhanced with decreasing pressure (Fotiadis and Jensen 1990). Accordingly, low pressures, which reduce buoyancy-driven flows, favor rapid transport of particles away from the deposition zone and reduce particle contamination. The carrier gas has also a significant influence on

F. Teyssandier and A. Dollet

266

C..arr~ ~

N2 :

~,,,,= S.TS'/E--m.~,,,= O.O00E+O0 e..,=~ too~+oo, e,,,,,= o.oooE+oo

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0

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2

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4

5

6

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8

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~/~o Figure 10.4. Simulation of the CVD of Si from Si2H 6 under reduced pressure in a circular impinging jet reactor, (a) deposition rate profile on the whole substrate surface, (b) streamlines (right-hand side) and isotherms (left-hand side), (c) total and partial deposition rate of Si-carrier species after Wang et al. (1994). (T = 973 K, P = 1333 Pa, V0 = 100.6 m/s).

Chemical Vapor Deposition PRODUCT

SEED PARoTICLES .- ......

J , . oo o

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@

_

9

-

267

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the thermophoretic transport, and under the same operating conditions the thermophoretic coefficient for nitrogen is smaller than that for hydrogen (Fotiadis and Jensen 1990).

10.3. S O L I D P H A S E F O R M A T I O N

Once introduced into the CVD reactor, the reactants may experience some chemical transformations during their transport towards the substrate. When reactants and reaction products reach the solid-gas interface, the gaseous chemical species undergo new complex physical and chemical interactions that may lead to film growth provided that some thermodynamic and kinetic criteria are satisfied. We will first discuss the energetics of the solid-gas transformation, and subsequently, the kinetics of the elementary heterogeneous processes. The nucleation and crystal growth phenomena will be finally addressed.

10.3.1 Solid-gas thermodynamic equilibrium approach The final equilibrium state of a thermally-activated CVD system, in which various chemical reactions are involved in the formation of a solid material from a gas phase, is imposed by thermodynamics whether or not this state is reached for kinetic reasons. In other words, when quantitative predictions are difficult to obtain because of the lack of mechanistic and kinetic information, thermodynamic calculations may provide reliable qualitative trends. Several review articles have already been devoted to this subject (Bernard and Madar 1990, Bernard and Madar 1991, Spear and Dirkx 1990). When all gaseous and solid species are taken into account, the "closed box" (nonflowing closed system) thermodynamic calculation at equilibrium provides valuable qualitative information on the nature of the different phases formed, on their amounts,

268

F. Teyssandier and A. Dollet

as well as on the influence of the various parameters. Two methods have been developed for the calculation of chemical equilibria in complex chemical systems: one based on mass action relationships, and the other based on minimization of the Gibbs free energy of the system. The latter is the method currently used in commercial computational programs. In the calculations it is implicitly assumed that (i) thermodynamic equilibrium is reached (no kinetic control), (ii) transport phenomena can be neglected, (iii) all the species that may form at equilibrium are taken into account in the calculation, and (iv) the Gibbs free energy of the species is known with good precision. When the residence time of the gas mixture in the hot zone of the reactor is short and/or the temperature of the gas mixture is low, kinetic limitations cannot be ignored and a detailed mechanistic investigation of CVD reactions is required.

10.3.2 The structure of CVD layers One of the major advantages of the CVD technique is that each of the operating parameters can be selected within a wide domain of variation, leading to a great variety of film structures. Strictly speaking, the possible structures of CVD layers are so numerous that it is almost impossible to discuss all of them in detail; nevertheless, they can be classified into several categories and sub-categories, so that their essential features can be captured. The first level of classification consists in dividing materials into amorphous materials and crystalline materials. Amorphous materials (AM) consist of metastable solid phases, generally highly non-organized, exhibiting only short-range atomic order. Their composition may differ significantly from the stoichiometric composition of the corresponding thermodynamically stable phases. Moderate temperatures are required in order to obtain this kind of material using the CVD technique, and PECVD, which allows near room temperature deposition, is often preferred to prepare AM on substrates which are very temperaturesensitive. One of the main features of deposited CVD AM is that, despite their disordered structure at the atomic scale, they may exhibit some structural organization, exemplified by columnar microstructures which have been observed in numerous cases (Thornton 1977). Crystalline materials (CM) can be divided into epitaxially grown monocrystalline materials and polycrystalline materials. The term epitaxy is used to indicate that the growing film structure is controlled by the (mono)crystalline substrate structure. When the mismatch between the respective lattice parameters of the substrate and the desired film is not too large, and also when particular growing conditions are satisfied (see further), the substrate acts as a seed crystal, imposing its lattice parameters to the growing material. When the film and the substrate are of the same nature and the film grows exactly in the same crystalline structure as the substrate, the growth process is called homoepitaxy. When the film and the substrate are made of different materials (lattice mismatch is highly probable in this case), the film may still grow with a crystalline structure imposed by the substrate lattice, but mechanical strains appear inside the film, the intensity of which depends mainly on the degree of mismatch. If the initial mismatch is non negligible, structural defects such as dislocations will nec-

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269

[ZONE 1] lZONE 2l [ZONE 3] '~

~' .,#'

.~_ ..

~'~

0.3

0.5

SUBSTRATE TEMPERATURE (TITm)

Figure 10.6. (a) Processing/morphology map for CVD of TiCN (Cheng et al. 1987). (b) Structure zone model (Movchan and Demchishin 1969).

essarily be created in order to relax the strains. This second epitaxial growth process is referred to as heteroepitaxy. A large variety of materials such as Si, GaAs, GaN, SiC, etc. are epitaxially grown using the CVD technique, mainly for applications in the field of electronics. Polycrystalline materials (PCM) are easily obtained using standard CVD conditions. They have generally highly anisotropic microstructures (depending on the operating conditions but also on the film thickness), generating anisotropic physical or chemical properties. The various surface morphologies of CVD PCM have been classified in many categories such as: plate-like, lenticular, star-shaped, pyramidal, spherulitic, cauliflower, dendritic, etc. Processing/morphology maps can be drawn to provide a general idea of the influence of operating parameters on surface morphology. As an example, Figure 10.6a shows such a map for CVD TiCN (Cheng et al. 1987). Another approach used to schematize the microstructure evolution is to draw a structure zone model (Movchan and Demchishin 1969, Thornton 1977). Three zones have been distinguished by Movchan and Demchishin (1969) for the coating microstructures (Fig. 10.6b) which appear to be a function of T~ Tm, where T is the substrate temperature and Tm the PCM melting point. The first coating structure is observed at T/Tm < 0.3. At such moderate temperatures, surface diffusion of the adsorbates is low: tapered crystals with domed tops are observed, separated by voided boundaries. The second structure corresponds to 0.3 < T~ Tm < 0.5. Surface diffusion is favored in this case, columnar faceted crystals with dense boundaries are observed (Thornton 1977), the size of which tends to increase with coating thickness. The last structure (0.5 < T~ Tm < 1) is controlled by bulk diffusion: the deposit is composed of equiaxed grains and exhibits a bright surface. In order to better account for the substrate roughness, a fourth zone (not presented in Fig. 10.6b), zone T, has been added by Thornton (1977) between zone 1 and zone 2, which consists of fibrous grains with dense boundaries. The above zone models were established for PVD, however, the general trends remain the same for CVD. The values of T~ Tm separating the various zones may vary significantly from one material to another, and should only be considered as approximate values in the best cases. For any zone, it is observed that the grain size increases with T~ Tin.

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F. Teyssandier and A. Dollet

It should be kept in mind that although they seem to be very attractive, the above approaches remain empirical since they are completely specific to the reactor, and also to the material for processing/morphology maps. 10.3.3 The driving f o r c e f o r crystal growth f r o m the v a p o r p h a s e

The gas-solid thermodynamic equilibrium assumption, discussed above, is obviously never satisfied in practice in the CVD process since it implies that there is no net mass exchange between the two phases, that is, no condensation or crystallization of the vapor, nor solid evaporation. The equilibrium condition is: # g -- ~s

(10.1)

w h e r e # g and/Zs are the chemical potentials of the gas phase and solid phase, respectively.

Nevertheless, when #g ~: /Zs, the system tends to evolve towards its minimum energy, i.e. thermodynamic equilibrium. The condition #g > /Zs is then required in order for deposition to be observed. The chemical potential excess in the vapor phase, A/z, is the chemical driving force for crystal growth from the vapor. For quantification purposes, the relative supersaturation a defined as: cr -- e Alz/RT

(10.2)

is generally used instead of A#. The absolute supersaturation: aa = a - 1, is also used in some cases. From the values of a or aa, the trends in the evolution of a solid-gas systems can be predicted: 9 0 < a < 1 (or aa < 0) dissolution of solid into gas phase 9 cr = 1 (or aa = 0) solid-gas equilibrium 9a > 1 (or O"a > 0) solid deposition from vapor phase a can be related to the partial pressure of the various species present in the gas phase by using the classical thermodynamic relations for the chemical potentials. However, rigorous analytical expressions of a are not practically easy to obtain for real CVD systems, as many chemical species and sometimes several complex solid phases have to be taken into account. Authors often illustrate the evolution of the microstructure of CVD films by the correlation diagram proposed by Blocher (1974) in the PVD case. However, the microstructure of CVD films depends on supersaturation in a more complex manner, because of the chemical reactions. Figure 10.7 shows the experimentally observed evolution of the microstructure of CVD films as a function of temperature (which influences the surface mobility of adsorbed species) and flux of active species towards the surface. The microstructural order of the deposit is mainly the result of a competition between these two factors. Several attempts have been made to evaluate the supersaturation in CVD processes, in order to correlate supersaturation with film microstructure. The validity of the calculations carried out for that purpose is often questionable, either because equilibrium

271

Chemical Vapor Deposition

Effect of increased surfacetemperature

l

epitaxialgrowth whiskers,dendrites.... polycrystals fine grain polycI3"stals amorphousdeposits

1

Effectof increased flux of reactivespecies

Figure 10.7. Effect of temperature and flux of reactive species on CVD film microstructure.

approximations have been used (Langlais et al. 1982), or because simple but meaningless definitions of supersaturation have been chosen (Goto et al. 1992). Vandenbulcke and Vuillard (1977) have proposed a more rigorous approach by evaluating supersaturation from the concentration of species at the gas-solid interface, taking into account surface reactions. They have plotted the observed microstructure of boron deposits as a function of the supersaturation and flow rates. It has been found that formation of amorphous deposits is favored at low temperatures and high supersaturation. Later, Vandenbulcke et al. (1991) have also proposed a schematic plot of the various microstructures of carbon (graphite and diamond) prepared by CVD as a function of the calculated supersaturation. However, in this study like in most of the others, only very simplified chemical schemes have been considered.

10.3.4 Surface mechanisms and kinetics Surface kinetics are completely ignored in the thermodynamic approach which predicts only the final state of a solid film assuming its thermodynamic equilibrium with the saturated vapor. However, the formation of structural defects in solid films and metastable solid phases (which is evidence of the non-equilibrium state of a CVD system) can only be explained by kinetic limitations induced by the heterogeneous processes. In this section, a very brief survey of elementary surface processes will be first given, then the methods for obtaining the related kinetic parameters will be presented and illustrated on several examples dedicated tO the CVD process. 10.3.4.1 Rate of elementary surface processes The most important heterogeneous phenomena occurring at the film surface during deposition are (i) adsorption (physisorption or chemisorption), (ii) surface diffusion, (iii) reaction on the surface, and (iv) desorption (Fig. 10.2). The amount of species adsorbed at steady-state depends on the nature of both the surface and the adsorbates, on the pressure and on the temperature. Surface processes are thermally activated but it is worth to note that the energy barrier for a given heterogeneous reaction is almost always significantly lower than that of its homogeneous counterpart (Laidler 1987). As will be seen in Section 10.3.6.1, a solid surface is not smooth at the atomic level: various kinds of surface site exist at which adsorbed species may be attached more or less strongly (Laidler 1987, Tsong 1990). At low temperature, most of the adsorbates are not able to overcome the energy barriers of the various surface processes. They cannot move over distances large

272

F. Teyssandier and A. Dollet

enough to find the energetically most favorable surface sites for incorporation, a condition which is required for the film to grow as a crystalline material. Nevertheless, some local minima of potential energy exist at the immediate vicinity of the adsorption sites, in which they can be trapped; in this case, only an amorphous solid phase can be grown. Among the most important factors determining the kinetics of the above-mentioned heterogeneous processes are the binding energies of surface atoms. The rate constants of the elementary surface processes are generally assumed to obey an Arrhenius equation. However, the measurement or the evaluation of Arrhenius parameters of surface processes is not a simple task since they depend on the surface coverage by adsorbates (Zhdanov 1991), and also, when occurring, on adsorbate-induced surface reconstruction (Zhdanov 1991, Somorjai and Van Hove 1995). 10.3.4.2 Estimation o f the kinetics of surface processes in CVD Kinetic parameters or binding energies of surface atoms can be obtained experimentally by means of various techniques such as temperature-programmed desorption (TPD), isothermal desorption (ID), field-ion microscopy (FIM), scanning tunneling microscopy (STM), laser-induced thermal desorption (LITD) (Zhdanov 1991, Tsong 1990). Among the most efficient techniques, FIM and STM are used to obtain direct information on surface diffusion, whereas kinetic information on desorption and surface reactions are generally obtained from TPD and ID analyses. Molecular beam reactive scattering (MBRS) is also an efficient method for providing information on the surface kinetics in CVD according to Gates (1996). It should be pointed out that such experiments are usually expensive and difficult to perform. Moreover, since chemical vapor deposition is generally carried out at relatively high temperatures, there are so many possible reaction pathways that a complete experimental determination of the kinetics of all possible surface processes in CVD appears rather unrealistic. Because of the complexity of the heterogeneous processes, the kinetic equations which are most often included in CVD reactive transport models are empirical overall kinetic equations. Some reaction mechanisms are assumed a priori and the related kinetic parameters are subsequently adjusted until an acceptable agreement between theoretical and experimental results (mainly the deposition rate profiles) is obtained (Roenigk and Jensen 1987). The simplest procedure for describing the surface kinetics in CVD is probably the introduction of a reactive sticking coefficient (s), i.e., the probability of surface reaction. During thermal accommodation (Herlin et al. 1991), the increase of the internal energy may be either transformed into kinetic energy favoring desorption or into rotovibrational energy inducing breaking of bonds. Sticking coefficients can be obtained from specific measurements such as molecular beam scattering at very low pressure (Buss et al. 1988, Fisher et al. 1992). They can also be deduced from indirect experimental ways, for instance, by examining the conformality of films deposited inside micrometric trenches. A model simulating deposition in such geometries, i.e., Monte-Carlo (Dew et al. 1992) or analytical (Cale and Raupp 1990, IslamRaja et al. 1991) models, provides the s value (Fig. 10.8). It is possible to find the value of s for a particular chemical species if its incident flux at the surface can be monitored independently (Matsuda et al. 1990, Fisher et al. 1992). Langmuir kinetic equations, taking into account the surface coverage by adsorbates, are also often chosen a priori to describe the surface kinetics of the CVD process (Roenigk and Jensen 1987).

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273

Figure 10.8. Simulationof the deposition profile in an infinite trench for various values of the sticking coefficient (1, 0.1, 0.01) (IslamRajaet al. 1991).

Several theoretical methods, that allow calculating the kinetics of the various surface processes (Molecular Dynamics (MD) and Monte-Carlo (MC) simulations, transitionstate theories (TST), electronic structure calculations) can be used to analyze the kinetics of the surface processes involved in CVD (Koseki and Ishitani 1992, de Sainte Claire et al. 1996); however, a severe drawback of these sophisticated methods lies in their complexity. A paper by Wang and Pollard (1995) has been recently dedicated to the presentation of a simpler theoretical approach for modeling surface reaction kinetics in CVD, which combines simple statistical mechanics, TST, and bond dissociation energies. The method requires four steps: (i) identification of the nature of the various sites and the corresponding bonding configurations of the adsorbates, (ii) evaluation of the thermodynamic properties of the reacting species, (iii) description of the nature of the elementary processes, and (iv) assessment of the structure and properties of the transition states. The method has been applied to tungsten deposition from WF6 and Sill4, and Fig. 10.9 shows some typical adsorbates on the (100)W surface. Because of its relative simplicity, this approach can only provide rough estimations of the kinetic parameters of the various surface reactions but calculations are not time consuming and, as a consequence, a very large number of elementary steps can be investigated. In conclusion of this section, it must be recalled that surface chemistry plays a key role in determining the structure of films grown by CVD. This subject has been further discussed, for instance, in a recent review article by Gates (1996). Unfortunately, it should be also emphasized that in spite of the importance of the elementary surface processes in CVD, reliable kinetic data remain generally difficult to obtain.

10.3.5 Heterogeneous nucleation Since it is the initial stage of crystal growth from a vapor phase, heterogeneous nucleation is one of the most dominating factors in the microstructure formation. In this section, a brief insight into the fundamentals of heterogeneous nucleation will be first given and some of its implications in the field of CVD will be illustrated. Nucleation consists in the formation of crystalline embryos (clusters or nuclei) at the substrate surface, due to localized density fluctuations in the supersaturated vapor. Some gaseous species are first adsorbed on the substrate surface, which either move along this surface or are temporarily trapped in potential wells. During their movement on the sur-

274

F. Teyssandier and A. Dollet

ooeO H F Si

W

Figure 10.9. Schematic diagram of the (100)W surface with typical adsorbates on the surface. A tungsten atom is present at the apex of each pyramid (Wang and Pollard 1995).

face the adatoms may interact with each other to build-up clusters, but the adsorbates which have not been incorporated in nuclei of critical size before the end of their life time re-evaporate. If a sufficient number of stable nuclei can be formed, the moving adatoms increasingly attach to the stable nuclei, which become bigger, and a redistribution of their sizes is observed. When they become big enough, they come into contact with each other, leading to their size increase at the expense of their number. 10.3.5.1 Nucleation theories Two distinct theories are generally used for analyzing the nucleation process. The first one, mainly based on thermodynamic considerations, is referred to as classical nucleation theory. Fundamental data such as nucleation barrier height, critical size of nuclei, nucleation frequency, equilibrium shape of nuclei, can be obtained using this theory (see Chakraverty 1973, Bennema 1973, Mutaftschiev 1993, among other articles). When the supersaturation is high, the classical nucleation theory becomes difficult to use and a second theory, the atomistic nucleation theory, must be applied. Nucleation is then regarded as a kinetic polymerization process: instead of a critical size and a critical energy, a critical number of atoms and cluster dissociation energies are defined (Rhodin and Walton 1963). Convenient analytical expressions of the nucleation rate, critical size and cluster density are given by several authors, such as Lewis and Campbell (1967), in the case of atomistic theory. Most often in CVD, the critical nucleus size is of the order of only a few atoms and the atomistic nucleation theory is used instead of the classical theory (Liu and Dandy 1996). An example of the use of the non-classical (atomistic) theory in the field of CVD can be found in a paper by Liu and Dandy (1996) who have studied the influence of substrate materials, surface diffusion and adsorption state on the initial stage of diamond nucleation on carbide-forming substrates. The theoretical results obtained have

Chemical Vapor Deposition

275

been found to agree qualitatively well with previous experimental observations, in spite of the hypotheses formulated (not recalled here). The major findings of this study are that: (i) surface diffusion of adatoms is the most determining factor influencing the nucleation kinetics of diamond (it depends on the substrate and temperature); (ii) when increasing the substrate temperature, chemisorption is favored at the expense of physisorption and surface diffusivity is also increased. As a result, the mean residence time of adsorbed species on the surface is increased and nucleation is enhanced. 10.3.5.2 Two- and three-dimensional nucleation When an embryo does not perfectly wet the substrate, the value of the free energy of interface formation is high and a significant supersaturation is required to overcome the nucleation barrier. Three-dimensional stable clusters are formed in this case, the orientation of which with respect to the substrate does not depend on supersaturation. On the other hand, when the substrate is perfectly wetted, two-dimensional nucleation may take place, provided that the gas phase is supersaturated with respect to the 2D solid phase. In this case, however, the orientation of the 2D phase may depend on the ledge free energy and on the supersaturation (Mutaftschiev 1993). A temperature increase favors the transition from 3D nucleation to 2D nucleation, since the enhancement of their mobility allows the adatoms to surmount local energy barriers, such as barriers for jumping down steps (Levi and Kotrla 1997).

10.3.5.3 Epitaxial nucleation and growth At equilibrium, an adsorbate layer undergoes two kinds of opposing interactions: the firstones are the interactions between the adatoms, and the second-ones are the interactions between the adatoms and the substrate atoms. If the substrate-adsorbate interactions are strong as compared to the adsorbateadsorbate interactions, and if the distance between neighbor-substrate atoms is comparable to the distance between atoms in the normal bulk structure of the growing film, the nucleation of two-dimensional clusters on the substrate and layer-by-layer growth may be favored. In this case, the substrate imposes its structure on the two-dimensional growing material, even if the distance between the atoms of the growing crystal planes (hkl) parallel to the substrate surface is slightly different from the distance between the atoms of the corresponding isolated three-dimensional bulk material. Because the influence of the substrate on the film structure decreases with increasing number of new monolayers, the chemical potential of the 2D phase, which is minimum for the first monolayer, gradually increases for the upper monolayers. For layer-by-layer growth to be observed, the gas phase may be undersaturated with respect to the infinite three-dimensional solid phase, but must always remain supersaturated with respect to the growing two-dimensional phase. This process, known as layer-by-layer or Frank-Van der Merwe mechanism (Frank and Van der Merwe 1949a,b), can only be observed below the so-called roughening temperature (RT). If the substrate-film lattice mismatch is more significant, the substrate may still impose its structure on the first monolayer (the gas phase is then supersaturated with respect to the 2D film structure), but the internal strains in this layer become large. However, as mentioned above, the influence of the substrate on the upper ad-layers decreases and, up

276

F. Teyssandier and A. Dollet

Frank Van der Merwe (FVM)

Stranski-Krastanov ( S K )

Volmer-Weber (VW)

Figure 10.10. Nucleation/growth modes in epitaxy.

to a critical film thickness, the substrate can no longer impose an abnormal structure to the film and layer-by-layer growth mode stops. If the system is supersaturated with respect to the 3D solid phase, three-dimensional nucleation follows and the film continues to grow with its normal bulk structure. This particular process is known as the Stranski-Krastanov mechanism (Stranski and Krastanov 1938). If the substrate-film lattice mismatch is large and the substrate-film interactions are not strong compared to the adsorbate-adsorbate interactions, three-dimensional nucleation occurs on the bare substrate surface, at supersaturation with respect to the 3D film structure. This simple three-dimensional nucleation process is called the Volmer-Weber mechanism (Volmer and Weber 1926). The above mechanisms are schematically summarized in Fig. 10.10.

10.3.5.4 Multi-component nucleation and chemical reactions Up to now, the discussion has been restricted to the case of single-component nuclei formed from a single-component vapor without taking into account the kinetics of the heterogeneous chemical reactions, which are of primary importance in CVD. Although most of the fundamental features of the nucleation process have already been exposed, formation of multi-component nuclei with chemical reactions presents some specific aspects which deserve particular discussion. A general analysis of the problem has been presented by Kukushkin and Osipov (1995) who have classified the elementary processes which are likely to occur when heterogeneous nucleation from a multi-component vapor takes place. They have taken into account the possible formation of solid solutions as well as eutectics and chemical compounds and have proposed the following processes: (i) direct formation of multi-component nuclei, (ii) formation of nuclei after a chemical reaction within an adsorbed layer, (iii) chemical reaction within the nuclei, (iv) eutectic separation in a solid phase, (v) evaporation of nuclei, and (vi) further growth of nuclei. A particular sequence of some of these processes can be formulated to describe nucleation from a multi-component vapor under conditions of interest. 10.3.5.5 Nucleation enhancement Increasing the substrate temperature generally enhances nucleation. This is mainly the result of the increased surface mobility of adatoms which have thus a better chance to be incorporated in existing embryos and subsequently form critical nuclei. The use of scratched substrates as well as ion bombardment are known to be efficient solutions for nucleation enhancement in CVD. For example, a bias pretreatment can be

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used for the growth of heteroepitaxial diamond films on monocrystalline silicon substrates by plasma CVD (Strckel et al. 1998). A dc bias voltage is applied to the substrate for a few minutes prior to deposition that results in a bombardment of the substrate by ions from the plasma. This special procedure can increase the nucleation density up to about 1012 cm -2, whereas the diamond nucleation density is only about 104 cm -2 on untreated substrates; however, only a very small fraction of the nuclei are epitaxially oriented. Surface active species (surfactants) can be intentionally incorporated during epitaxial growth in order to change the growth mode, and particularly, to promote the layer-by-layer mode (Markov 1997). Predeposition of one monolayer of such species (which are not incorporated into the film) may suppress the undesirable islanding observed in some cases and may also reduce the defects resulting from strain accommodation. 10.3.5. 6 Selective vapor deposition Selective deposition (Carlsson 1990, Gladfelter 1993) is one of the most interesting features of the CVD process. The deposition selectivity has most interesting potential for advanced submicron devices on patterned substrates, but other fields are emerging such as growth of artificial 2-D and 3-D materials. Selectivity usually refers to the process ability to control the deposition mechanisms in such a way that the film growth occurs only on one surface in the presence of others (area-selective deposition). However, several categories of selective growth systems exist, and CVD allows depositing simultaneously and selectively on different materials:

- different phases: phase-selective deposition different microstructures: microstructure-selective deposition - different chemical composition: chemical-composition-selective deposition -

Selective growth is based on differences in the initial interfacial reactions between the vapor and the various surfaces of the substrate and particularly on the differences in incubation times required for the nucleation. The selective growth period is thus up to the longest incubation time for nucleation. As seen above, epitaxial growth is carried out at supersaturation below that required for 3D heterogeneous nucleation, i.e., at near equilibrium conditions. Thermodynamic analysis is for that reason used to predict trends in selectivity (Cho et al. 1987, Madar and Bernard 1990, Engqvist et al. 1992). A new understanding of selectivity has been proposed for the area-selective deposition of epitaxial silicon in the presence of SiO2 blanket areas (Fitch 1994). At low pressures, hydrogen atoms passivate and block adsorption of silicon species, whereas, at high pressures, nuclei etch away faster than they grow. 10.3.5. 7 Theoretical and experimental studies o f nucleation in CVD For a long time, only analytical approaches based on empirical or semi-empirical relations have been used for the theoretical analysis of nucleation, but owing to the increasing power of computers, numerical simulations become more and more the rule for that purpose. Several numerical models have already been proposed for studying the nucleation process in CVD (Zachariah et al. 1996, Mahalingham et al. 1997), which provide key information on the reaction kinetics and coalescence mechanisms.

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However, it must be kept in mind that some basic experimental information is required in order to check the accuracy of the predictions. Many experimental studies on nucleation kinetics have been conducted in CVD. Atomic Force Microscopy (AFM) and Scanning Electron Microscopy (SEM) have often been used to determine the cluster density and size as a function of time and CVD process parameters. Since AFM and SEM measurements are usually performed ex-situ, several deposition experiments of variable duration are necessary in order to obtain the time evolution of the cluster characteristics. Nevertheless, some experimental analyses of nucleation kinetics have also been carried out in-situ using techniques such as ellipsometry (Drevillon 1993), or light-scattering (Sheldon et al. 1992). In this case, however, a schematic nucleation model must be used in order to interpret the experimental measurements.

10.3.6 Crystal growth Heterogeneous nucleation is the first stage, let us say the birth of a crystalline solid. The following stage, the ripening of this solid, is crystal growth. Like for nucleation, we have thought that it could be helpful to remind very briefly some of the most basic ideas of crystal growth before discussing some particular examples related to CVD. Far more complete and rigorous analyses of the fundamentals of crystal growth will be found in many excellent books and review articles (Hartman 1973a, Hurle 1993, Villain and Pimpinelli 1995, Levi and Kotrla 1997).

10.3.6.1 Mechanisms of crystal growth At an atomic scale, the surface of a growing crystal can be represented by terraces, ledges (or steps) and kinks as first suggested by Kossel (1927). The species to be incorporated are referred to as growth units. The most favorable sites (with the lowest energy) for incorporating the growth units are the kink sites (so-called half-crystal sites). Accordingly, the growth units first adsorb on a terrace, then diffuse until they either reach a ledge and then finally a kink site where their incorporation takes place, or desorb (TLK model). Nucleation should normally provide the source of repeatable steps on flat terraces, assuming that supersaturation is high enough for the nucleation barrier to be overcome; however, surprisingly, crystal growth can be observed even at very low supersaturation. The source of repeatable steps has remained unidentified until Frank (1949) proposed a spiral growth model in which the steps are continuously generated at the emergence of a screw dislocation.

10.3.6.2 Classification of crystal faces The various faces of a given crystal have different density of steps and kinks. Some faces present theoretically no steps (like terraces) and have a flat appearance: they are referred to as singular faces. Other faces, which have a rough aspect at an atomic scale present a more or less important density of steps (and kinks) and are referred to as non-singular faces. A more precise classification has been proposed from the concept of periodic bond chain (PBC) (Hartman and Perdok 1955). A given (hkl) crystal face is referred to as a flat face (F), stepped face (S), kinked face (K). F faces correspond to singular faces, whereas

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S and K faces correspond to non-singular faces. It must be mentioned that F faces have a fiat aspect only below the so-called roughening transition.

10.3.6.3 Crystalline morphology Some experimental observations have suggested that a crystalline form {hkl } has a greater morphological importance when the spacing between the (hkl) planes, dhkl, is large. This result, known as the Bravais-Friedel law, has been further extended by Donnay-Harker (Hartman 1973a) and Donnay-Donnay (Hartman 1973a). The concept of attachment energy Eatt, the energy (per molecule) released when a new slice is attached to a crystal face, is most often used to obtain the crystalline morphologies (habits) of large crystals. In the Hartman-Perdok theory (Hartman and Perdok 1955), it is proposed that (i) the higher Eatt for a given (hkl) face, the less the morphological importance of this face in the growth form, (ii) the growth rate of F faces, R• is proportional to Eatt. This law has been found to be satisfying in many cases (Hartman and Bennema 1980), although the relation: R• (x Ea2tt, should be preferred for needle-like crystals (Van der Sluis and Kroon 1990). Using the above relations, it is possible to estimate the growth rate of the various slow-growing faces, and then to construct the crystal habit (see below). 10.3.6.4 Morphology and normal growth rate of crystal faces Since incorporation of growth units is energetically more difficult on terraces than at kink or step sites, F-faces should theoretically be the slowest-growing crystal faces. This is experimentally verified unless the reactivity of non-singular faces is lowered due to strong surface reconstruction and/or adsorption of impurities (elements which are not incorporated in the solid structure). Crystal habit is controlled by the slowest-growing faces, that is immediately obvious from an examination of Fig. 10.11. Starting with a morphology in which all possible crystal faces exist, and keeping the growth rate of these faces constant, it can be inferred that: (i) the fastest growing faces disappear quickly, (ii) two adjacent faces grow independently when the angle 0 between them is <90 deg, whereas the persistence of one of the two faces (or both) depends on their relative rate of growth and on cos 0 when 0 > 90 deg (Alexandru 1969). Stable growth of CVD layers. Aspar et al. (1993) and Rodriguez-Clemente et al. (1993) have studied the surface morphology of polycrystalline A1N films prepared by CVD. They have first used the Donnay-Harker (Hartman 1973a) approach in order to draw a theoretical equilibrium morphology assuming that the growth rate of a given (hkl) face is inversely proportional to its lattice spacing dhkl. From the PBC of A1N structure deduced from literature analysis, they have concluded that the slow-growing F faces are the non-polar face { 100 } and the polar faces { 101 }, { 10T }, {001 }, {001 }. In this case, the two approaches Donnay-Harker (Hartman 1973a) and Hartman-Perdok (1955) have given similar results, but the real morphologies observed did not match exactly the theoretical-one, and varied with the operating parameters. To account for these discrepancies, Rodriguez-Clemente et al. (1993) have discussed the reactivity of A1N crystal faces using a broken bond model. They have reached the conclusion that the negative polar faces { 101 } and {001 } should be more reactive than their opposite positive polar faces, and that they should be inhibited by hydrogen adsorption. Some other authors (Dollet et

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R1

~

l

RFI

'~RF2

FX~'l IRr3 F lll a)

b)

Figure 10.11. Left: time evolution of crystalline morphology from the relative growth rate of possible faces. (a) F faces with comparable growth rates (all faces survive), (b) S face between two F faces (the fast-growing S face disappear). Right: illustration of the competitive growth process (2D). 2 identical cubic crystals are considered. The growth rates perpendicular to the substrate surface are R1 et R2 for crystal 1 and 2 respectively. Each line represents a time step.

(oo-l)~ {lOl} 33 ~

a)

I subs rate b)

Figure 10.12. Surface morphology of (112) oriented AIN, prepared by CVD (Dollet et al. 1997). (a) SEM picture (left). (b) Theoretical morphology, using the SHAPE program (Dowty 1991) (right).

al. 1997) have also studied the growth of polycrystalline A1N films by CVD, and have observed surface morphologies essentially composed of pyramidal crystals in which the { 101 } and {001 } forms dominate (Fig. 10.12), in agreement with the theoretical findings of Rodriguez-Clemente et al. (1993). Owing to reactor modeling results, it has also been suggested that the surface coverage by ammonia, or perhaps the deposition rate, may play a key role in the preferred orientation. Unstable growth. During crystal growth from the vapor, some fluctuations (thermal, chemical, etc.) may disturb the moving solid-gas interface. If these fluctuations cannot be smoothened, they may become amplified and morphological instabilities will result. Dendrites, skeletal crystals (Nanev 1994), etc. are good examples of crystal morphologies

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which form under unstable growth conditions. It has been shown in previous atomistic simulations (Xiao et al. 1991) that a growing crystal may exhibit compact faceted or open dendritic morphologies depending on various factors such as supersaturation, mean free path or temperature. The apparition of protrusions has been observed to result from volume diffusion whereas the anisotropic surface kinetics has been found to exert a stabilizing effect. Although atomistic simulation seems to be the most powerful method to tackle the problem of morphological instability, it should be mentioned that continuum equations with moving boundary conditions have also been solved for the same purpose in CVD (Palmer and Gordon 1989) but this approach has generally been restricted to the growth of non-faceted crystals or amorphous materials.

10.3.6.5 Relation between preferred orientation and morphology of polycrystalline films Polycrystalline films are composed of a large number of individual crystals which are generally not oriented in the same direction. A mechanism known as evolutionary selection or competitive growth has been proposed by van der Drift (1967) to explain the preferential orientation of deposited polycrystalline layers. The nuclei are assumed to be randomly oriented on the substrate surface and during their growth, some crystals come into contact and a geometrical selection operates. Only the crystals having their fastest direction of growth nearly perpendicular to the substrate surface will finally emerge at the film surface. The crystals which are less favorably oriented are gradually buried by the others, and, as a result, the grain size and orientation sharpness increase with increasing film thickness. This is illustrated in Fig. 10.11, where it can be seen that the crystal on the fight will be soon buried by the one on the left. An evolutionary selection process has been proposed by Wild et al. (1990), to explain the growth of polycrystalline diamond films by CVD, which agrees well with the observations from X-ray texture analyses, showing that the initial orientation of the tiny crystals is random, and that the sharpness of the preferred (110) orientation increases with increasing film thickness. A 2D geometrical model has been formulated by the authors in order to simulate the growth of the diamond films.

10.3.6.6 Theoretical studies of crystal growth from the vaporphase Evolution of crystalline morphology, grain size and texture, formation of defects inside the film (voids, vacancies, etc.) during film growth results from a large number of complex phenomena at the atomic scale which are difficult to apprehend without numerical models. Beside continuous geometrical models (Wild et al. 1990, Tang et al. 1990, Barrat et al. 1996), atomistic models such as ballistic aggregation models (Henderson et al. 1974, Dew et al. 1992), molecular dynamics (MD) and Monte Carlo simulation models (Dawnkaski et al. 1996, Smith 1997) become more and more used for understanding the exact mechanisms of solid phase formation in CVD. 10.4. C O N C L U S I O N S

Modern manufacturing methods are placing more emphasis on high performance techniques. An accurate control of these techniques requires an improved knowledge of the physico-chemical phenomena that are involved in the growth of materials. We have tried

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to summarize in this review the current status of mechanistic research in the field of materials grown from a vapor phase in the presence of a purely thermally activated process. In particular, it has been emphasized that numerical modeling supported by in situ diagnostics is a key tool for the understanding and improvement of the CVD process. However, things become more problematic for the heterogeneous phenomena since they are even more complex than their homogeneous counterparts. Calculations assuming thermodynamic equilibrium can be run in order to observe, at least qualitatively, the influence of the operating parameters on the composition of the solid-gas system. Nevertheless, since CVD remains in essence a non-equilibrium process, whenever possible, a kinetic analysis of both the homogeneous and the heterogeneous processes should be preferred. It is worth to note that an increasing amount of work is dedicated to the numerical simulation of nucleation and crystal growth in CVD. Several interesting simplified approaches have been proposed for that purpose. However, sophisticated atomistic simulations are increasingly used and already provide invaluable information though their performances are still limited by the computing power available. Unfortunately, in spite of the progress accomplished in this field, evaluating the kinetics of the elementary surface processes will undoubtedly remain, for a rather long time, the most difficult task. The understanding and full-scale modeling of the whole process, including the formation of film microstructure, appears to be the most important challenge of the next years for research in chemical vapor deposition.

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Wang, Y. E and Pollard, R. (1995) J. Electrochem. Soc., 142, 1712. Wang, Y. B., Bisch, C. and Teyssandier, E (1995) Chem. Eng. Sci., 50, 1605. Westmoreland, E R., Howard, J. B., Longwell, J.E and Dean, A. M. (1986) AIChE J., 32, 1971. Wild, C. H., Herres, N. and Koidl, E (1990) J. Appl. Phys., 68, 973. Wu, J. J., Nguyen, H. V., Flagan, R. C., Okuyama, K. and Kousaka, Y. (1988) AIChE J., 34, 1249. Xiao, R. E, Iwan, J., Alexander, D. and Rosenberger, E (1991) J. Cryst. Growth, 109, 43. Zachariah, M. R., Carrier, M. J. and Tsang, W. (1996) in Proceedings of the XIII International Conference on CVD, Los-Angeles USA, eds. Besmann T. M., Allendorf, M. D., Robinson McD. and Ulrich R. K. (The Electrochemical Society, Pennington), p. 1. Zhdanov, V. E (1991) Surf. Sci. Rep., 12, 183.

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Chapter 11 Combustion Synthesis 11.1. 11.2. 11.3. 11.4. 11.5. 11.6. 11.7. 11.8. 11.9.

Introduction Thermodynamic Considerations Kinetic Considerations Field-Activated Combustion Synthesis The "Azide" Process SHS Reactions in Synthesizing Ti3SiC2 Controlled Reactions in the Ti-B Binary System Auto-Ignition Synthesis of Nanocrystalline Oxides Non-Equilibrium Effects 11.10. Concluding Remarks Acknowledgments References

289 291 294 295 297 299 301 304 307 307 308 308

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Chapter 11 Combustion Synthesis S.B. BHADURI and S. BHADURI

11.1. I N T R O D U C T I O N

The technique of self-propagating high-temperature synthesis (SHS) is being used all over the world, not only in synthesizing many advanced materials but also in various industrial applications. Other terminologies used to describe the process include combustion synthesis (CS), gasless combustion, etc. The current interest in the topic can be credited to the pioneering work of Merzhanov and his group in Russia, starting in the early seventies (Merzhanov and Borovinskaya 1975). However, contrary to the popular belief, the process itself has a much longer history. The earliest mention of a combustion phenomenon in materials synthesis can be found in 1825. Berzelius (1825) reported that Zr metal powders burn to form the corresponding oxide. The thermodynamics of the process was explained by Goldschmidt (1898). In fact, Goldschmidt can be rightly called the originator of the aluminothermic process, a subgroup of SHS reactions. During the next 100 years, there were some sporadic industrial discoveries utilizing the combustion phenomena (Alexander 1941, Krapf 1964). Despite those discoveries, there was no widespread scientific interest in the process until vigorous research was initiated by Merzhanov and his colleagues. In the USA, it is generally believed that the work in this field started with support from the Advanced Research Project Agency (ARPA) in the mid- 1980s. Long before that, however, Walton and'Poulos (1959) started utilizing the aluminothermic processes for synthesizing cermets. This activity went on for several decades. As the SHS related activities took firm root, excellent reviews started appearing in the literature. Frankhouser et al. (1986) published an extensive monograph on the subject. This monograph was interesting in that it had an industrial bias and was a balanced review of the work done in the Soviet Union and elsewhere in the West. Soon after that, two of Munir's reviews appeared in 1988 and 1989 (Munir 1988, Munir and Anselmi-Tamburini 1989). Munir's reviews together describe the important aspects of the then Soviet work, which by then was already a decade and a half old. Yi and Moore (1990), as well as Merzhanov (1990), reviewed somewhat similar topics. The paper by Koizumi is an important contribution in reviewing the Japanese work (Koizumi 1990). A short paper by Hlavacek (1991) points out the historic aspects of the process and mentions the earliest work by Berzelius, Goldschmidt and others. Other reviews also appeared in the literature (Subrahmanyam and Vijaykumar 1992) including the most recent ones by Moore and Feng (1995a, b). However, in spite 289

290

S.B. Bhaduri and S. Bhaduri

of almost ten reviews in the field within the last decade, the contents of the reviews have not progressed much further since the publication of the review by Munir and AnselmiTamburini (1989). Since some of the information included in this review was actually quoted from the earlier Russian references, it is believed that the data need confirmation in view of the more recent reports. Furthermore, there have been some interesting developments in the field, which need proper attention and dissemination. Finally, in spite of a great deal of applications in industry, worldwide skepticism towards the process has prevailed. Hence, there is a need for reviewing some of the recent and interesting developments in the field. The present attempt will not be comprehensive. Instead, it will focus on the synthesis of newer materials and developments of new techniques with a view towards ultimate application. Unless large-scale applications of SHS are realized, skepticism will remain. We would like to point out that the industrial applications of SHS are by no means insignificant in this country. But elsewhere in the world, particularly in Russia, other countries of the CIS, Europe (notably Spain), and Asia (specifically Japan, China, Korea, Taiwan and India), industrial applications have been substantial. Combustion synthesis, or Self-propagating High temperature Synthesis (SHS), relies on the ability of highly exothermic reactions to be self-sustaining. In a typical combustion synthesis reaction, the mixed reactant powders are pressed into a pellet of a certain green density and subsequently ignited either locally at one point (propagating mode) or by heating the whole pellet (simultaneous combustion mode). The reaction can be ignited in a number of ways including laser irradiation, radiant flux, resistance heating coil, spark, chemical oven, etc. The exothermic reaction is initiated at the ignition temperature T/g and generates heat, which is manifested in a maximum or combustion temperature, Tc. The maximum temperature obtained can be controlled by adding diluents, the nature and amount of which can be determined depending on the maximum temperature desired. The products of combustion synthesis are usually extremely porous, typically 50% of the theoretical density. Techniques such as hot pressing, hot isostatic pressing, and shock waves have been employed to densify the products. Compared with conventional processing methods, combustion synthesis has the following advantages: (1) the high temperatures generated can volatilize low-boiling-point impurities and therefore produce high-purity products, (2) the simple exothermic nature of the process avoids the need for expensive processing facilities and equipment, (3) the short processing times result in low operating and processing costs, (4) the high thermal gradients and rapid cooling rates can result in the formation of non-equilibrium phases, (5) inorganic materials can be synthesized and consolidated into a final product in one step by utilizing the chemical energy of the reactants. Because of these advantages of the process, it has been possible to produce novel and improved materials with improved mechanical, electrical, optical, and chemical properties. It has been estimated (Yi and Moore 1990) that over 500 compounds have been synthesized by the combustion synthesis process. These include resistive heating elements, abrasives, cutting tools, and polishing powders, shape memory alloys, high-temperature intermetallics, electrodes for electrolysis of corrosive media, thin films and coatings, functionally-graded materials, and composite materials.

Combustion Synthesis

291

This review will start off with a discussion of the thermodynamic aspects of the process. This will be followed by some specific examples of materials development such as nitrides, silicides, metal matrix composites, and oxides.

11.2. T H E R M O D Y N A M I C

CONSIDERATIONS

Once ignited, high temperatures can be achieved in a very short time owing to the highly exothermic nature of the reactions. The enthalpy of the synthesis reaction for the propagating case, AHreaction, is equal to AHf,298, the standard enthalpy of formation of the product. For an adiabatic reaction with the reaction temperature denoted by Tad, the heating of the product from 298 K to Tact takes place in response to the input of A Hreaction. Depending upon whether Tad is lower than, equal to, or higher than the product melting temperature Tm, the following three cases arise: AHf,298 --

f T~d

Cp(product) dT

if Tad
(11.1)

,/298

Tad AHf,298 --

Cp(product) d T +

vAHm

if Tad = Tm

(11.2)

d298

rap

[ Tad

Cp(produc0d T + AHm +

AHf,298 -3 298

,l Trap

Cp(liquid)dT

if Tad > Tm

(11.3)

where, v is the fraction of the liquid phase in the mixture of liquid and solid phases, since both of them are expected to exist at the melting point. For many compounds, the thermodynamic data required to calculate Tad are available in the literature. However, agreement between the experimentally obtained data and theoretically calculated values is lacking. This is because the reactions are not truly adiabatic. At best, the calculations predict the upper limit of Tad. Nevertheless, such calculations provide important guidelines for maintaining the exothermic state of the reactions. Based on the observations of a number of reactions (Munir and AnselmiTamburini 1989, Moore and Feng 1995a, b), a linear correlation has been found between AHf,298/Cp,298 and Tad. It has been suggested in the literature that the reaction will not be self-sustained if AHf, z98/Cp,298 < 2000~ (or Tad values are less than 1800~ (Munir and Anselmi-Tamburini 1989). Table 11.1 shows the results of Tad calculations for various silicides. AH/,298 for the materials under consideration was obtained from (Barin et al. 1977, Chart 1973). Cp (in U/mole K) is written as A + B T 4-C T 2, where the coefficients A, B and C are taken from (Barin et al. 1977) and are listed in Table 11.1. If substantial melting takes place, (as in the case of TisSi3 (Rupp et al. 1992)), the calculations may require the value for Cp (liquid), which is taken as 33.74.n J/mole K, where n is the number of atoms in a molecule and is equal to 8 for TisSi3. Table 11.2 compares the as-predicted values of Tad for various silicides with those reported (Subrahmanyam and Vijaykumar 1992), which are based on the Russian data. Even though the Russian values are available for only some materials,

292

S.B. Bhaduri and S. Bhaduri

Table 11.1. Thermodynamic data and adiabatic temperatures, Tad, for various silicides.

Compound

A

B

C

A Hf

A Hm

(kJ/mol)

(kJ/mol)

M.P. (K)

Tad A H f / Cp

(K)

ZrSi

44.9

8.74

2.95

153.92

58.73

2368

3.84

2909

ZrSi2

62.8

15.27

-2.78

158.50

44.75

1790

2.47

2228

188.0 47.84

30.58 11.36

-14.9 -5.41

572.42 128.96

59.59 40.18

2483 1843

3.17 2.86

2739 2349

TiSi2

70.01

17.47

8.99

133.54

43.44

1773

2.05

1808

Ti5Si 3 CrSi

46.75 51.71

10.70 8.69

-4.8 -8.36

138.40 54.60

178.88 35.12

2408 1730

2.90 1.21

2528 1215

Zr5Si 3 TiSi

CrSi2 Cr5Si 3

65.23 197.5

22.38 49.00

-7.74 -25.5

79.62

42.39

1730

1.26

1252

221.73

33.02

1920

1.21

1233

ReSi

52.29

9.57

-3.74

52.42

50.16

2153

1.03

997

ReSi2

67.39

10.98

-6.07

89.86

73.22

2253

1.41

1420

Re5Si 3 MnSi

189.7 49.00

44.93 12.69

-13.9 -6.36

156.42 78.29

46.89 29.41

2233 1548

0.83 1.71

901 1553

Mn5Si 3

200.2

53.87

-19.5

199.68

21.71

1573

1.03

1100

12.57

-2.78

137.28

69.17

2203

2.14

1863

449.28

73.51

2753

2.49

2222

NbSi2 Nb5Si 3 TaSi2 Ta5Si 3

62.82 188.1 72.84 178.6

30.61

-14.9

7.67

-9.01

118.39

89.03

2473

1.82

1671

38.90

-8.86

332.80

72.10

2773

1.85

1772

MoSi2

67.44

11.90

-6.53

130.96

74.83

2303

2.06

1872

Mo5Si3 W5Si 2

182.3 67.43

34.84 10.98

-11.9 -6.06

307.84 92.19

56.65 85.16

2463 2433

1.72 1.44

1672 1471

W5Si 3

178.6

38.98

-8.82

133.79

63.53

2593

0.74

945

11.61

-9.36

124.8

51.75

1953

2.00

1792

459.26

49.08

2283

2.26

1703

VS2 V5Si 3

71.05 187.3

118.1

-17.2

A, B, and C in J mol-1 K - l , 10 -3 J mol-1 K - l , 105 J mol-1 K - l , respectively.

the agreement is quite sound for most of the cases provided, while there are substantial differences in the predicted values of Tad in the case of TiSi and CrSi2. Tables 11.3 and 11.4 show similarly the calculated Tad values for various carbides as well as borides. Again, there are agreements as well as discrepancies between our data and the Russian predictions. Notable differences exist in the carbides of ZrC and HfC and the borides CrB2 and MoB. Similar comments can also be made in the case of various borides. At this point several comments can be made. First, since these predictions serve as important guidelines, the new values should be carefully considered before initiating any reaction. Second, combustion of stoichiometric TiB composition has a peculiar result. Even though the predictions favor reactions of higher exothermicity than that of the TiB2 reaction, in practice such reactions are quite sluggish. This is due to the formation of a significant amount of the liquid phase, which can be attributed to the large quantity of unreacted Ti remaining within the system. This occurrence can be exploited in fabricat-

Combustion Synthesis

293

Table 11.2. Comparison of predicted Tad values of selected silicides.

Material

Tad (K) Present work

Tad (K) Russian work

TiSi TiSi2 Ti5Si 3 ZrSi ZrSi 2 Zr5Si 3 VSi2 V5Si 3 NbSi2 Nb5Si 3 TaSi2 Ta5Si3 CrSi CrSi2 Cr5Si3 MoSi2 Mo5Si 3 WSi 2 WsSi 3

2349 1808 2528 2909 2228 2739 1792 1703 1863 2222 1671 1772 1215 1252 1233 1872 1672 1471 945

2000 2500

2800

1900 1800

1800 1900 1500

ing m e t a l m a t r i x c o m p o s i t e s . Third, it is s e e n that, the f o r m a t i o n o f several c o m p o u n d s is quite s l u g g i s h as w e l l ( c o r r e s p o n d i n g to Tact t e m p e r a t u r e s in the r a n g e o f 1000 K). T h e r e are m a n y strategies that are f o l l o w e d in initiating r e a c t i o n s in such cases. C o m m o n m e t h ods to ignite such r e a c t i o n s are the " E x p l o s i o n M o d e " , the " C h e m i c a l O v e n " , and u s i n g

Table 11.3. Calculated adiabatic temperatures for selected carbides.

Material TiC ZrC HfC VC NbC TaC Mo2C WC

Tad (K) Present work

Tad (K) Russian work

3290 3758 4512 2228 2821 2901 1017 1163

3210 3400 3900 2400 2800 2700 1000 1000

294

S.B. Bhaduri and S. Bhaduri Table 11.4. Calculated adiabatic temperatures for selected borides. Tad (K) Present work

Tad (K) Russian work

TiB

3388

3350

TiB2

3198

3190

ZrB 2

3322

3310

HfB 2

3662

3520

VB 2

2546

2670

VB

2490

NbB 2

2715

2400

TaB 2

2728

2700

CrB 2

1475

2470

CrB

1646

Material

MoB 2

1480

MoB

2659

1800

WB

1617

1700

FeB

1605

Fe2B

1172

CoB

1939

Co 2B

1721

NiB

2026

Ni2B

1106

ternary reactions with several different phases taking part in the reaction. A recent advancement in this field is the use of Field Assisted Combustion Synthesis (FACS), which will be described later. Fourth, Tad calculations in the case of nitrides have not been discussed since it is felt that the "Azide" process has better capabilities, which again will be discussed later. Finally, the thermodynamics of the "Auto Ignition" or "Combustion Synthesis in the Auto Ignition" mode will also be described.

11.3. K I N E T I C C O N S I D E R A T I O N S

In order to understand the kinetics of the process, a cylindrical compact is placed vertically in a reaction chamber and ignited from the top. A self-supporting reaction wave is initiated under suitable conditions and its steady speed of combustion is measured. Of interest is the dependence of the reaction upon the variables of the experiment (such as specimen diameter, specimen composition, size of the initial pressed particles, and the extent of dilution with an inert component such as the reaction product itself). These are well documented in Munir (1988), Munir and Anselmi-Tamburini (1989), and Moore and Feng (1995a, b). Azatyan et al. (1979) have derived an expression for the rate of propagation of a planar combustion front based upon an elementary model of the process. In this model, the

Combustion Synthesis

295

calculations are performed assuming a uniform radial distribution (a planar front). With the additional assumptions of there existing a single heat source and that the only heat loss from the body is by conduction and radiation, the generalized form of Fourier's one dimensional heat equation takes the following form:

OT _ kO2T

Cpp Ot

ot

a T4

O~ 4- Qpdp(T, rl) - 2-(Tr - To) - 2e--(r

O(T, o) -

Orl

-

Td)

(11.4)

(11.5)

Ot

where Cp is the heat capacity, p, the density, T, the temperature, To, the ambient temperature, 77, the fraction reacted, t, the time, k, the thermal conductivity, x, the axial distance, Q, the heat of reaction, or, the heat transfer coefficient, r, the sample radius, E, the emissivity coefficient, a, the Stefan-Boltzman constant, and 4~, the kinetic function. Neglecting the transient effects from the ignition surfaces and the heat losses, the above equation can be rewritten as:

OT kO2T Cpp 0---;~ 4- Qp~(T, 7) _

(11.6)

The solution to the above equation can be written as

u2 - a n Q p

E Koexp

RTad

(11.7)

where an is a constant depending upon the order of the reaction, u is the velocity, E is the activation energy and K0 is a constant. The apparent activation energy of the reaction can be found from a graph of ln(u / Tad) vs. 1/ Tad. 11.4. F I E L D - A C T I V A T E D C O M B U S T I O N

SYNTHESIS

During the discussion on thermodynamic calculations, it was mentioned that the reactions are frequently quite sluggish and do not propagate in a self-sustaining manner. Several alternate methods for propagating such reactions were also suggested. Unfortunately, these alternative methods lead to other undesirable phases. A specific example of this phenomenon is the combustion of MoSi2 in the Mo-Si binary system. As noted earlier, the Tad value for MoSi2 formation is at the threshold of self- propagation. In order to enhance the propagating reaction, the explosion mode can be tried, which in turn leads to the formation MosSi3, a compound that generally does not have the high temperature oxidation resistance of MoSi2. To alleviate this problem, Munir and his co-workers ignited the combustion synthesis reactions in the presence of electric fields (Gedevanishvili and Munir 1994, Feng and Munir 1994, Shon and Munir 1995, Munir et al. 1995, Feng and Munir 1995a, b, Gedevanishvili and Munir 1995, Xue and Munir 1996, Munir 1996). This process is now referred to as field-activated combustion synthesis (FACS). The idea here

296

S.B. Bhaduri and S. Bhaduri Table 11.5 Materials synthesized using FACS. Materials

Reference

SiC MoSi2-SiC MoSi2-Nb, MoSi2-ZrO 2 B4C-TiB 2 W5Si 3, WSi2

Feng and Munir (1994, 1995a, b) Gedevanishvili and Munir (1994) Shon and Munir (1995) Xue and Munir (1996) Gedevanishvili and Munir (1995)

is to create localized Joule heating in the reaction zone. Indeed, this occurred when monitoring the voltage, current, and the resistance of the sample, which essentially remained constant during propagation. A second discovery was the apparent dependence of wave propagation on the field itself, as clearly demonstrated by the propagation ceasing whenever the field was shut off. Thirdly, the propagation velocity was directly related to the magnitude of the field. A fourth observation noted a threshold value of the field, below which combustion did not take place, while at relatively high values of the field, combustion took place spontaneously. Finally, the actual mechanisms for such reactions could be studied by turning off the field and stopping the propagating reaction at any time. Such reactions are easy to model by adding Joule heating term of m E 2 to the right hand side of Eq. (11.6). Feng and Munir (1995a) modeled such a reaction using a twodimensional version of the modified Eq. (11.6) and the results are shown in Figure 11.1. The graphs show profiles of current densities near the top, center and bottom of the sample. The profiles show that a planar front is moving through the sample. The modeled experimental configuration is shown in Figure 11.2. The system consists of spring-loaded graphite electrodes between which the powder compact is placed. The compact is then ignited from the side. Table 11.5 shows some of the materials produced by the FACS process. Gedevanishvili and Munir (1994) studied the SHS reactions in the MoSi2-SiC system. In such a case, the formation of both constituents is sluggish. However, beyond the threshold voltage, a desirable composition of MoSi2-SiC composite was produced. The threshold field level decreased as the percentage of Si and C in the system was increased. Feng and Munir (1994, 1995a and b) used the process to react Si and C to form SiC. The threshold voltage was 9.5 V and simultaneous combustion occurred for voltages greater than 30 V. Shon and Munir (1995) reported FACS processing of MoSiz-ZrO2 and MoSiz-Nb composites. The secondary phases are added to enhance the mechanical properties with the application of field, without the formation of undesirable phases. In the absence of the field, however, the reactions do not self-sustain and undesirable phases form. The same result occurs in the case of B4C-TiB2 composites. While the formation of B4C is sluggish, the reaction to form TiB2 is highly energetic. Under normal circumstances, the thermodynamic equations predict the propagating reactions only in restrictive compositions. However, with FACS, propagating reactions can take place easily (Xue and Munir 1996). FACS is also useful in producing WsSi3 with a Tad of 954 K. Since a propagating reaction is not normally expected here, a high voltage is necessary.

Combustion Synthesis

297

top

~--.-4~..$s

I000 |

500

.

:"

?.Us "~Ss "less

1000

~"

500

.-. 1000 E 500

0

0.0

0.5

1.0

1.5

2.0

x (cm) Figure 11.1. Currentdensity profiles from side ignition (Feng and Munir 1995a).

11.5. THE "AZIDE" PROCESS Nitrides are typically formed by igniting a reaction between a metal and nitrogen gas. There are two competitive processes that take place during such reactions: the kinetics of reaction and the permeation of the gas. Typically, permeation of the gas is the rate-controlling step. The conversion efficiency can be increased by making use of two strategies: increasing permeation by increasing the porosity of the compact or by increasing the gas pressure. Frequently, the formation of a molten phase adversely affects the permeation of the gas. Inclusion of diluents helps to reduce Taa values and without allow-

S.B. Bhaduri and S. Bhaduri

298 Graphite

Laser beam Igniter!

ON2N

()

,a~.... ~:,.;~,J, PrO luct Rea tants

power (

O

~

?

Wave Fro ~ ' K~actants 'Product

Figure 11.2. Schematic representations of field activated combustion synthesis (Feng and Munir 1995a).

ing a molten metal to form. Despite these strategies, conversion to the final nitride may not be significant. For example, if the initial starting materials have low melting temperatures, then controlling their melting is extremely difficult. A case in point is the synthesis of A1N by SHS, during which substantial melting of A1 is a problem. Solid-state nitrogen sources can solve this problem. One such source of nitrogen is sodium azide, NAN3. The material is cheap, readily available, and is unstable. Holt (1984a, b) pioneered the "Azide" process for making nitrides by the SHS process. Originally, the process was developed for making nuclear fuels. Typical reactions with the respective Tad values are written as: 3U + NaN3 = 3UN + Na 3UO2 + 6Ca + NaN3 = 3UN + 6CaO + Na

Tad = 3669 K Tad -- 2824 K

During the process, the nitrides form and metallic sodium vaporizes. In the case of the second reaction, CaO is extracted from the product and is expected to offer much better control. Similar ideas can be used in making important ceramic materials as well, which are described by the following equations: A1 + NaN3 = A1N + N2 + Na 3Si + 4NAN3 = Si3N4 + 4N2 + 4Na Ti + NaN3 = TiN + N2 + Na

Tad -- 3010 K Tad = 2537 K Tad = 3022 K

Table 11.6 shows the Tad values calculated in the case of the "Azide" process vs. the conventional gaseous N2 reactions. In some cases the calculated Tad values for the "Azide" process are lower (e.g., in the cases of TiN and Si3N4) than the conventional situation, while in other cases the Tad values by the "Azide" process are higher. Nevertheless, generating N2 in-situ is an important advantage, so much so that excess pressure of N2 is not required (Holt 1984a, b, Amosov et al. 1992). In spite of these potential advantages, the growth in this field has not been significant. Recently, Amosov et al. (1992) used this process for synthesizing several nitrides. They reported that NH4C1, when incorporated into the powder mixture, sometimes catalyzes the reaction. Lee and Chung (1997) explained the catalytic effect of NH4C1 on the basis

299

Combustion Synthesis

Table 11.6 Comparisonof calculated Tad values of nitride formation. Material TiN ZrN Si3N4 A1N UN

Tad (K) (Azide process)

Tad (K) (Conventional)

3022 2864 2537 3010 3669, 2824

4900 4900 4300 2900 3000

of vapor phase decomposition of NH4C1, formation of SiClx, followed by the nitridation of SiClx to form Si3N4. The maximum temperature of combustion was 2123 K, which is much lower than the Tad value calculated. Since vapor phase reactions took place, the product contained loose powders, whiskers, and platelets. Lee et al. (1995) also used the "Azide" process to synthesize A1N powder. As predicted by the thermodynamic calculations, Tad for such reactions are higher than those for the formation of Si3N4. In order to counteract the fragmentation of the processed pellet, a casket of Ti and C powder was built around the A1 and NaN3 mix. The reaction between A1 and NaN3 was actually ignited by the combustion reaction between Ti and C. The conversion was reasonably high. Similar to the Si3N4 reaction, gas phase reactions also took part in A1N formation. 11.6. SHS REACTIONS IN SYNTHESIZING T i 3 S i C 2

Considerable attention has recently been focused on Ti3SiC2 due to its mechanical behavior (Goto and Hirai 1987, Pampuch et al. 1989, Lis et al. 1993, Racault et al. 1994, Tong et al. 1995, Arunajatesam and Carim 1995, Radhakrishnan et al. 1996, Barsoum and E1-Raghby 1996). A complete understanding of the mechanical properties can be obtained only if bulk samples of Ti3SiC2 are synthesized. Such samples have been obtained by sintering or hot-pressing Ti3SiC2 powders obtained through combustion synthesis. Combustion synthesis yields powders that contain up to 10-30 vol% TiC impurities. Consequently, these powders must be pulverized and hot pressed to obtain bulk samples for testing. As opposed to this two-stage process, we have hot pressed Ti, Si, and C powders at 1400~ and 40 MPa for 90 min to obtain bulk samples (Radhakrishnan et al. 1998). X-ray diffraction (XRD) pattems indicate the presence of Ti3SiC2 with about 10 vol% of TiC. The hardness of Ti3SiC2 at room temperature is 4.1 GPa and decreases to 1 GPa beyond 1000~ This suggests significant yielding of the compound beyond 1000~ If this is the case, it should be possible to fabricate complex shapes by extrusion, forging, injection molding, etc. A fracture toughness of 6.15 4-0.4 MPa m 1/2 was obtained from chevron-notched specimens. The fracture surface revealed a layered profile and cleavage planes (Fig. 11.3). Despite possessing superb mechanical properties, Ti3SiC2 lacks sufficient resistance to oxidation. A composite with SiC was investigated to improve the resistance to oxida-

300

S.B. Bhaduri and S. Bhaduri

lO~m Figure 11.3. SEM micrograph of combustion synthesized Ti3SiC2.

Figure 11.4. SEM Micrograph of fracture surface of displacement reaction processed Ti3SiC2.

tion. The solid state displacement reaction between TiC and Si was utilized to fabricate Ti3SiC2/SiC composites. The composite microstructure revealed a uniform distribution of SiC in a TizSiC2 matrix (Fig. 11.4). Two morphologies of SiC were identified: blocky and needle-shaped. The composite, when subjected to oxidation in air at 1000~ registered a weight gain of the order of a few milligrams after 10 hours. A very thin oxide layer was formed. There-

301

Combustion Synthesis i: ~: ; ~ ~ ~

:.

9

Figure

:~ : ~ : ~ :!~i : ~

,

~ ,

:

~

:~

. . . . .

:,

~

; ~ ,,

I

~

~

:::,i~;:i~!~ I~

:i,',,,;,,,~

,~- ~,

~

~

11.5. A Vickers indentation on Ti3SiC2. The bar is 5 0 / z m ,

fore, our objective to improve the oxidation resistance was successful (Radhakrishnan et al. 1998).

The effect of SiC on the mechanical properties was also investigated. A hardness of 4.74 GPa was obtained. The increase in hardness can be attributed to the addition of SiC. Cracks from Vickers indentation, when examined using SEM, provided evidence of crack bridging and crack deflection (Fig. 11.5). The bend strength of the composite was 405 MPa. The fracture surface showed fracture around SiC particles as well as a layered Ti3SiC2 matrix.

11.7. C O N T R O L L E D

REACTIONS

IN THE Ti-B

BINARY

SYSTEM

Another example of the utility of thermodynamic predictions in controlling vigorous combustion reactions is in the Ti-B binary system. It is generally believed that the formation of single-phase TiB2 is highly exothermic (Frankhouser et al. 1986, Sheppard 1986, Munir 1988, Munir and Anselmi-Tamburini 1989, Yi and Moore 1990). A low melting ternary element such as Fe or Cu may be incorporated to control the reaction between Ti and B (Li et al. 1992). Addition of extra elements, however, gives rise to other undesirable phases, rendering the interpretation of mechanisms and the evaluation of properties difficult. Therefore, control of the reaction can be sought by simply shifting to a composition away from the TiB2 stoichiometry (Peng et al. 1996, Peng and Bhaduri 1997). In the following paragraphs, the results of this work are presented. First, a thermodynamic study was carried out to investigate how the exothermic nature changed. Pertinent thermodynamic properties of TiB and TiB2 are shown in Table 11.7. Figure 11.6 shows the variation of Tad with respect to the free Ti acting as a diluent. Phase changes affect the temperature of the system and consequently the reaction rate. The reaction (1 + x)Ti + 2B --+ TiB2 + xTi, for instance, absorbs the latent heat of fusion of 100.4 kJ/mol, while maintaining Tad close to the melting temperature. This trend con-

302

S.B. Bhaduri and S. Bhaduri

Table 11.7 Thermochemical properties of the compounds in the Ti-B system.

Tad

Melting point Tm (~

B

C

(oC)

Product

Ti

B

Comments

56.379

25.867

- 17.464

2919

2913

1670

2092

Tad = Tm

54.066

-0.033

-21.631

3115

2200

1670

2092

Tad > Tm

AH 0 Product

Heat capacity

kJ/mole

A

TiB 2

323.8

TiB

160.2

3100 ~El

2800

-

25~

-

22~

-

1900

-

\+

C

A

O v "O m

+ +

I.--

+ ~+ _+_1.

(1 + x ) T i +

_../X._2.

(1 + x ) ' l i + 2 B = T i B 2

16~

- --O--

13~

.......

B=TiB+xTi

~+

+xTi

~

3. (2 + x)'l'i + 3 . 5 B = 1 . 5 T i B 2 + 0 . 5 T i B + x r i

+

4. (2 + x)'ri + 2 . 5 B = 0 . 5 " f i B 2 + 1.5"fiB + x'fi

0

Figure 11.6.

I .......

0.2

I .......

0.4

I .......

0.6

I .......

0.8

Free Ti (mole)

I .......

1

I .......

1.2

!,,,

1.4

Thermodynamic predictions of how Tad varies with departure from stoichiometry.

tinues up to an addition of 0.65 mole of Ti, as depicted in Fig. 11.6. Tad in the reaction (1 + x)Ti + B ~ TiB + xTi drops drastically with the addition of Ti. Second, the combustion behavior was closely monitored in the chosen compositions. Elemental titanium ( - 3 2 5 mesh, >99.4 pct pure and - 1 0 0 mesh, >99.5 pct pure) and boron ( - 3 2 5 mesh, >99 pct pure) powders were used as starting materials. The three compositions chosen for reaction contained 18, 15, and 12% B, the rest being Ti. The compositions investigated were either stoichiometric or off-stoichiometric composition of TiB. Six grams of each composition were measured and then mechanically pressed at pressures of 135,340, and 500 MPa in a uniaxial press with a die of 13 mm in diameter to form cylindrical pellets, giving two groups of pellets with an identical density of 50-70% of theoretical. Figure 11.7 shows a series of photographs of the combustion front of the above compositions ignited under the same conditions. The white zone is the reaction product, while the black zone is the reactant mixture. The specimens of proper compositions in the Ti-B system retained their original form of cylindrical shape and dimension in cross-section

Combustion Synthesis

{a }

s

(

s

~s

303

!~

Figure 11.7. Propagation of the controlled combustion wave fronts in compositions" (a) Ti-18 wt.% B (b) Ti-15 wt.% B, and (c) Ti-12 wt.% B.

Table 11.8 Reaction rate versus composition and size of Ti powder.

Compositions

Density (% of theoretical)

Ti-12 wt.% B Ti-15 wt.% B Ti- 19 wt.% B

50-60 50-60 50-60

Velocity (mm/s) degassed at 550~ for 10 hrs -325 mesh Ti -100 mesh Ti 13 10 2.5

2 1.4 1.2

while shrunk in the longitudinal direction by about 10%. It is also noted that the combustion front was a planar surface propagating through the compact. The steady front is attributed to an optimum combination of the reaction parameters (such as composition, green density, and particle size). For compositions with Ti content of 88 wt.% the combustion of reactant mixture did not take place, whereas combustion reactions in compositions with Ti of 82 and 85 wt.% were completed. The effects of green density and composition on reaction rate are summarized in Table 11.8 based on visual observations. For compositions of 15-18 wt.% B with green density less than 70% of theoretical, the combustion of the reactants was smooth. Compositions with higher Ti content (e.g., 88 wt.%) propagated half way down the sample and then were self-quenched. When the green density of the compact reached up to about 70% of theoretical, none of the compacts of the three compositions were ignited. It is believed that the higher green density of the compact resulted in higher thermal conductivity, resulting in more heat loss from the combustion front to its surroundings.

304

S.B. Bhaduri and S. Bhaduri 14.

J.

- o - - fine Ti

12-"

.-e--- c o a r s e Ti

A

E E 0 tU

10-"

8

="

6

="

4

= =

2-

-

~

O-

9

80

82

9

9

9

9

9

84 wt%

Figure 11.8.

41

- . . . m _ . . O . ~ . ~ . " 9

|

.

86

.

|

.

L

,

88

.

.

.

90

Ti

Effect of titanium content with different particle size on propagation rate of combustion reaction.

The rate of combustion wave propagation was measured normal to the direction of the combustion based on the photograph of the reaction. The photographs were taken with a video camera at a film speed of about 30 frames per second. Both the Ti particle size and its composition have a large effect on the reaction rate, as shown in Table 11.8. The rate is seen to be proportional to particle size and the content of Ti, as depicted in Fig. 11.8. This implies that combustion in the Ti-B system may be optimized by selecting the proper size of powders and compositions for controlling the reaction. Under otherwise identical conditions, the velocity of the propagation through the fine powder mixture is considerably faster than that of the coarser particle mixture, which is quite consistent with conventional wisdom. Most importantly, the combustion velocity ranges are much lower than a typical highly exothermic reaction. Yet, the propagation of the combustion front was still smooth, indicating a controlled combustion. Figure 11.9 contains XRD patterns of SHS-produced samples, showing three phases of TiB, TiB2 and Ti. This figure clearly evidences the effectiveness of the control of reactions attempted in this work. In order to appreciate this, close attention must be paid to the TiB and Ti peaks located at the 20 angles between 38 ~ and 40 ~ From the top to the bottom patterns, the ratio of the peak intensity of TiB over Ti is reduced. Since the bottom pattern is for the TiB stoichiometric composition, free Ti retained in the stoichiometric composition seems to be less. It is also observed that the formation of TiB2 is very insignificant in off-stoichiometric compositions as compared to the stoichiometric composition.

11.8. A U T O - I G N I T I O N S Y N T H E S I S O F N A N O C R Y S T A L L I N E O X I D E S

In the auto-ignition process, atomistic mixing is achieved, resulting in homogeneous complex chemical products. The process is similar to the "Pechini" process (Pechini 1963) or "Glycine-Nitrate" process (Pederson et al. 1992) or has been termed as "Combustion Synthesis" (Aksay et al. 1991). While the term "Combustion Synthesis" tends to get

Combustion Synthesis 9T i B 2 o TiB

305

T rl

eTi

o

(c) 77- 12v~%B

o II m

b)_ Ti- 15 Ht%B

(9

=_

_

_

_

(~) Ti-18Ht%B .

25

.

.

.

I

30

.

.

.

.

.

.

35

.

.

.

.

40

.

.

.

.

45

.

.

.

.

.

50

"

;

- - "

55

-

-

-

i

60

-

-

-

-

65

Figure 11.9. X-ray diffraction patterns of the CS-synthesized products of compositions. (a) Ti-18 wt.% B, (b) Ti-15 wt.% B, (c) Ti-12 wt.% B.

confused with SHS, the "Glycine-Nitrate" and the "Pechini" processes are very special cases of such auto-ignition processes. Yet another variation is called the "Spray Pyrolysis" process (Kingsley and Patil 1988). The auto-ignition process had its origin in rocket propellant technology and utilizes oxidizers (e.g., metal salts) and fuels (typically organic compounds). When properly controlled, high temperatures can be generated by the exothermic properties of reduction-oxidation (redox) reactions between decomposition products of the oxidizer and the fuel, which drives the reaction. The process is controlled by such parameters as heating rate, stoichiometry, choice of fuel, etc. The burning mixtures generate a large amount of gas in a very short period of time leading to very fine particulates of friable agglomerates. The following are important features of the auto-ignition process: 9 It is a versatile process leading to the synthesis of solid solution composites as well as complex compound oxide phases in homogeneous form. 9 It can synthesize ultra-fine powders. 9 It uses cheap raw materials. 9 Rapid reaction leads to nucleation of crystallites without much growth. 9 It is a scaleable and a high production rate process. 9 The products are of high purity and are loosely agglomerated. 9 The products are easily sintered. While several ceramics have been reportedly synthesized by the auto-ignition technique, the effort in this area of research is not well coordinated. A great deal of work is being carried out by chemists with little appreciation for the materials themselves, while materials scientists have been performing this type of work without much understanding of fuels and their combustion behavior. To the best of our knowledge, there has been

306

S.B. B h a d u r i a n d S. B h a d u r i

no review on this subject. In the following, some examples of materials synthesized are presented. Examples of synthesis of simple single-phase oxides include A1203 (Aksay et al. 1991) and CeO2 (Verma et al. 1990). Examples of solid solutions include yttria doped zirconia and ceria doped zirconia (Aksay et al. 1991, Dhas and Patil 1994). However, the most important use of this technique, so far, is in synthesizing complex oxides. Because of the atomistic level mixing of raw materials, the final products are very homogeneous. Examples of complex oxides produced by this process include YAG (Aksay et al. 1991), lanthanum aluminate (Aksay et al. 1991), mullite (Hong et al. 1994), cordierite (Gopichandran and Patil 1993), 1-2-3 superconductors (Rambabu 1990), lanthanum chromite (Chick et al. 1990), barium titanate (Zhong and Gallagher 1995), and various ferrites (Ravindranathan and Patil 1987). It has been realized that heating and cooling rates are fast in such reactions. Therefore, they can be utilized in synthesizing nanocrystalline materials. A group at Alfred University has made use of this process for synthesizing nanocrystalline YzO3-doped ZrO2, and pure ZrO2 (Venkatachari et al. 1995, Huang et al. 1995). We have used the auto-ignition process to directly synthesize nanocrystalline oralumina (Bhaduri et al. 1996a). It is believed that the high temperatures generated for a short time crystallized the c~-phase. Since the cool down period was fast, the crystallites did not have the time to grow. Both XRD and transmission electron microscopy (TEM) studies confirmed the crystallite sizes to be less than 100 nm. This study was extended to the AlzO3-ZrO2 binary system to synthesize nanocomposites (Bhaduri et al. 1994, 1995, 1997b). It is well known that grain growth is considerably reduced whenever ZrO2 is added as a second phase to A1203. Nanocomposite powders containing A1203-14 wt.%ZrO2 were synthesized using the auto-ignition process. Both XRD and TEM studies showed that A1203 was in the c~ form and that ZrO2 has the tetragonal structure. The average crystallite size for A1203 was approximately 40 nm and that of ZrO2 was 13 nm. Examples of solid solution were CaO and CeO2 doped nanocrystalline ZrO2 powders (Zhou et al. 1996, 1997). The compositions contained 2 to 10 mol% of CaO and 4 to 20 mol% of CeO2 in ZrO2. The powders were characterized by XRD, Raman spectroscopy, and TEM. The crystallite sizes were about 5 nm in most of the compositions. The results indicated that the as-synthesized powders were in the tetragonal phase and not in the cubic phase. There is also evidence of extension of solid solubility because of the non-equilibrium nature of the process. Nanocrystalline spinel (or spinel-alumina nanocomposite) powders were also synthesized (Bhaduri et al. 1997a). Al-nitrate and Mg-nitrate were used as the starting materials and urea was used as the fuel. The crystallite sizes and lattice parameters were calculated from the XRD results. Lattice parameters of non-stoichiometric spinel solid solutions satisfied the Vegard's law. Spinel grains were about 13-20 nm range and or-alumina 30-40 nm, as determined by XRD and TEM.

Combustion Synthesis

307

11.9. N O N - E Q U I L I B R I U M E F F E C T S

As mentioned in the Introduction, the rapid heating and cooling effects involved in combustion synthesis can produce non-equilibrium phases. The non-equilibrium phases produced by this technique are neither as extensive nor as large in number as those obtained by rapid solidification or mechanical alloying techniques. Another important difference between combustion synthesis and other non-equilibrium processing techniques is that while combustion synthesis is used more for ceramic systems, rapid solidification and mechanical alloying are used more for metallic systems. Therefore, a direct comparison of the non-equilibrium effects produced by these different techniques is not fair. However, by combustion synthesis also supersaturated solid solutions and glass formation have been reported as also the formation of intermetallic phases away from the equilibrium stoichiometry. Kunrath (1998) reported that supersaturated solid solutions could be obtained in the TiC-Cr3C2 system. While the room temperature solid solubility of Cr3C2 in TiC under equilibrium conditions is very small (<5%), a single-phase solid solution phase was obtained in a TIC-50% Cr3C2 powder mixture produced by the combustion synthesis reaction. Glass formation was reported in the A1203-B203 system (Yi 1998). It has also been reported that combustion synthesis of this system under microgravity conditions has an important effect on the glass formation characteristics. When a combustion synthesis reaction occurs in a Ni-50 at% Ti powder mixture it is observed that a second phase is usually present along with NiTi. That this is due to the non-equilibrium conditions prevalent in the combustion synthesis processing is confirmed by annealing the product at high temperatures when only the NiTi phase formed (Moore and Yi 1992). It has also been reported (Merzhanov 1981) that during the synthesis of niobium and tantalum monoborides, the corresponding diboride forms as an intermediate product. This diboride will then combine with more metal and forms the monoboride. It has also been mentioned (Aleksandrov et al. 1984) that formation of NiA1 is preceded by the formation of two unidentified phases.

11.10. C O N C L U D I N G R E M A R K S

This chapter reviewed some recent developments in the Combustion Synthesis research. The results described are some of our own while other results were quoted from the literature. Nevertheless, the topics presented are all very exciting and have potential industrial uses. It is our opinion that the research in this area is going on unabated all over the world. This statement is corroborated by the record number of papers that were presented at the 3rd International Conference on the topic held in October 1997. The next conference is scheduled to be held in Moscow, Russia, where the number of papers is expected to be even larger.

308

S.B. B h a d u r i a n d S. B h a d u r i

ACKNOWLEDGEMENTS

It is a p l e a s u r e to a c k n o w l e d g e the s u p p o r t o f N S F g r a n t D M R s u p p o r t e d b y an N S F E P S C o R f e l l o w s h i p .

9 3 1 5 0 5 7 . SB w a s

REFERENCES

Aksay, I. A., Hahn, C., Maupin, G. D., Martin, C. B., Kurosky, R. E and Stangle, G. C. (1991) US Patent No. 3, 330, 697. Alexander, P. P. (1941) US Patent No. 2, 378, 368. Aleksandrov, V. V., Korchagin, M. A., Tolochko, B. P. and Sheromov, M. A. (1984) Combust. Explos. Shock Waves USSR, 19, 430. Amosov, A. P., Bichurov, G. V., Bolshova, N. E, Erin, V. M., Makavenko, A. G. and Markov, Y. M. (1992) Int. J. SHS., 1,239. Arunajatesam, S. and Carim, A. H. (1995) J. Amer. Ceram. Soc., 78, 667. Azatyan, T. S., Maltsev, V. M., Merzhanov, A. G. and Selzenev, V. A. (1979) Comb. Expl. Shock Waves, 15, 35. Barin, I., Knacke, O. and Kubaschewski, O. (1977) Thermal Properties of Inorganic Substances, Springer Verlag, Berlin. Barsoum, M. W. and E1-Raghby, T. (1996) J. Amer. Ceram. Soc., 79, 1953. Barsoum, M. W., Brodkin, D. and E1-Raghby, T. (1997) Scripta Met., 36, 535. Berzelius, J. J. (1825) Prog. Ann., 4, 126. Bhaduri, S. B., Radhakrishnan, R. and Linch, C. (1994) Ceram. Eng. & Sci. Proc., 13, 694. Bhaduri, S. B., Bhaduri, S., Watt, G. L., and Radhakrishnan, R. (1995) in Advances in P/M and Particulate Materials, eds. Philips, M. and Porter, J. (APMI, Princeton), p. 117. Bhaduri, S., Bhaduri, S. B. and Zhou, E. (1996a) Nanostructured Mater., 7, 487. Bhaduri, S., Bhaduri, S. B., Lewis, I. R. and Griffiths, P. R. (1996b) in Advanced Synthesis and Processing of Composites and Advanced Ceramics H, eds. Logan, K. V., Munir, Z. A. and Spriggs, R. M. (The ACerS, Westerville), p. 81. Bhaduri, S., Bhaduri, S. B. and Radhakrishnan, R. (1997a) Nanostructured Mater., 8, 755. Bhaduri, S., Zhou, E. and Bhaduri, S. B. (1997b) Ceram. Eng. & Sci. Proc., 18, 645. Chart, T. G. (1973) High Temperature--High Pressures, 5, 1241. Chick, L. A., Pederson, L. R., Maupin, G. D., Bates, J. L., Thomas, L. E. and Exharos, G. J. (1990) Mater. Lett., 10,6. Dhas, N. A. and Patil, K. C. (1994) Int. J. SHS., 3, 311. Feng, A. and Munir, Z. A. (1994) J. Appl. Phys., 76, 1927. Feng, A. and Munir, Z. A. (1995a) Metall. Mater Trans., 26 B, 581. Feng, A. and Munir, Z. A. (1995b) Metall. Mater Trans., 26 B, 587. Frankhouser, W. L, Brendly, K. W., Kieszek, M. C. and Sullivan, S. T. (1986) Gasless Combustion Synthesis of Refractory Compounds, Noyes Publications, NJ Gedevanishvili, S. and Munir, Z. A. (1994) Scripta Met. Mater., 31, 741. Gedevanishvili, S. and Munir, Z. A. (1995) J. Mater. Res., 10, 2642. Goldschmidt, H. (1898)Z. Elektrochem., 4, 494. Gopichandran, R. and Patil, K. C. (1993) Br. Ceram Trans., 92, 239. Goto, T. and Hirai, T. (1987) Mater. Res. Bull. 22, 1195. Hlavacek, V. (1991) Bull. Amer. Ceram. Soc., 70, 240. Holt, J. B. (1984a) US Patent No. 4, 446, 242. Holt, J. B. (1984b) US Patent No. 4, 459, 363. Hong, C. S., Ravindranathan, P., Agrawal, D. K., and Roy, R. (1994) J. Mater. Sci. Lett., 13, 1072. Huang, D., Venkatachari, K. R., and Stangle, G. C. (1995) J. Mater Res., 10, 762. Kingsley, J. J. and Patil, K. C. (1988) Mater. Lett., 6, 427. Koizumi, M. (1990) in Proc. 1st US-Japan Workshop on Combustion Synthesis, eds. Kaieda, Y. and Holt, J. B. (National Research Institute for Metals, Tokyo), p. 101. Krapf, S. (1964) US Patent No. 3,143,413. Kunrath, A. (1998) Ph.D. Thesis, Colorado School of Mines, Golden, CO, USA. Lee, W-C and Chung S-L (1997) J. Mater Res., 12, 805.

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Lee, W-C, Tu, C-L, Weng C-Y and Chung, S-L (1995) J. Mater. Res., 10, 774. Li, H. P., Bhaduri, S. B., and Sekhar, J. A. (1992) Metall. Trans., 23A, 251. Lis, J., Pampuch, R., Pierkarczyk, J. and Stobierski. L. (1993) J. Mater. Synthesis and Processing, 1, 93. Merzhanov, A. G. (1981) Arch. Combust., 1, 23. Merzhanov, A. G. (1990) in Combustion & Plasma Synthesis of High Temperature Materials, eds. Munir, Z. A. and Holt, J. B. (VCH Publishers), p. 1. Merzhanov, A. G. and Borovinskaya, I. P. (1975) Comb. Sci. and Tech. A., 10, 195. Moore, J. J. and Feng, H. J. (1995a) Prog. Mater. Sci., 39, 243. Moore, J. J. and Feng, H. J. (1995b) Prog. Mater. Sci., 39, 275. Moore, J. J. and Yi, H. C. (1992) Mater Res. Soc. Symp. Proc., 246, 331. Munir, Z.A. (1988) Bull Amer. Ceram. Soc., 67, 342. Munir, Z. A. (1996) Metall. Mater. Trans., 27A, 2642 Munir, Z. A. and Anselmi-Tamburini, U. (1989) Mater. Sci. Eng. Reports, R 3, 277. Munir, Z. A., Lai, W. and Ewald (1995) US Patent No. 5,380, 409 Pampuch, R., Lis, J., Stobierski, L. and Tymkiewicz, M. (1989) J. Eur. Ceram. Soc., 5, 93. Pechini, M. P. (1963) US Patent No. 3, 330, 697. Pederson, L. R., Chick, L. A. and Exarhos, G. J. (1992) US Patent No. 5, 114, 702. Peng, Z. and Bhaduri, S. B. (1997) Ceram. Eng. and Sci. Proc., 18, 637. Peng, Z., Watt, G. L. and Bhaduri, S. B. (1996) in Advanced Synthesis and Processing of Composites and Advanced Ceramics, eds. Logan, K. V., Munir, Z. A., and Spriggs, R. M. (The ACerS, Westerville), p. 11. Racault, C., Langlais, E and Naslain, R. (1994) J. Mater Sci., 29, 3384. Radhakrishnan, R., Brimhall, J. L., Henager Jr., C. H. and Bhaduri, S. B. (1996) Scripta Mater., 34, 1809. Radhakrishnan, R., Bhaduri, S. B. and Henager Jr., C. H. (1998) to be submitted to J. Mater Res. Rambabu, D. (1990) J. J. Appl. Phys., 29, 507. Ravindranathan, E and Patil, K. C. (1987) Amen. Ceram. Soc. BulL, 64, 688. Rupp, B., Holt, J. B. and Wong, J. (1992) CALPHAD, 16, 377. Sheppard, L. M. (1986) Adv. Mater Proc., 2, 25. Shon, I. J. and Munir, Z. A. (1995) Mater Sci. & Eng., 202, 256. Subrahmanyam, J. and Vijaykumar, M. (1992) J. Mater Sci., 27, 6249. Tong, X., Okano, T., Iseki, T. and Yano, T. (1995) J. Mater Sci., 30, 3087. Venkatachari, K. R., Huang, D., Ostrander, S. P., Schulze, W. A. and Stangle, G. C. (1995) J. Mater Res., 10, 748. Verma, H. K., Mukundan, P., Warrier, K. G. K. and Damodaran, A. D. (1990) J. Mater. Sci. Lett., 9, 377. Walton, J. D. and Poulos, N. E. (1959) J. Amen. Ceram. Soc., 42, 40. Xue, H. and Munir, Z. A. (1996) Metall. Mater Trans., 27A, 475. Yi, H. C. (1998) CCACS Workshop, Golden, CO, USA, September 1998. Yi, H. C. and Moore, J. J. (1990) J. Mater. Sci., 25, 1169. Zhong, Z. and Gallagher, E K. (1995) J. Mater. Res., 10, 945. Zhou, E., Bhaduri, S., Bhaduri, S. B., Lewis, I. R. and Griffiths, P. R. (1996) in Processing and Properties of Nanocrystalline Materials, eds. Suryanarayana, C., Singh, J. and Froes, E H. (TMS, Warrendale), p. 123. Zhou, E., Bhaduri, S. and Bhaduri, S. B. (1997) Ceram. Eng. & Sci. Proc., 8, 653.

This Page Intentionally Left Blank

Chapter 12 Nanostructured Materials 12.1. Introduction 12.2. Classification 12.3. Preparation 12.3.1 Inert Gas Condensation 12.3.2 Rapid Solidification 12.3.3 Electrodeposition 12.3.4 Chemical Reactions 12.3.5 Mechanical Attrition 12.3.6 Devitrification 12.4. Structure 12.4.1 Microstructure 12.4.2 Atomic Structure of the Crystal Lattice 12.4.3 Atomic Structure of the Grain Boundaries 12.4.4 Triple Junctions and Higher-Order Grain Junctions 12.5. Stability 12.5.1 Grain Growth in Nanocrystalline Materials 12.5.2 Grain Growth at Ambient Temperature 12.5.3 Examples of Grain Growth Inhibition 12.5.4 Isothermal Grain Growth Kinetics 12.6. Particulate Consolidation 12.7. Properties 12.7.1 Diffusion and Sinterability 12.7.2 Mechanical Properties 12.7.2.1 Elastic Properties 12.7.2.2 Hardness and Strength 12.7.2.3 Ductility and Toughness 12.7.2.4 Superplastic Behavior 12.7.2.5 Deformation Mechanisms in Nanoscale Materials 12.7.3 Magnetic Properties 12.7.3.1 Fundamental Properties 12.7.3.2 Soft Magnetic Materials 12.7.3.3 Hard Magnetic Materials 12.7.3.4 Other Ferromagnetic Nanocrystalline Materials 12.7.4 Chemical Properties 12.7.4.1 Corrosion Behavior 12.7.4.2 Catalytic Properties

313 313 314 314 316 317 318 318 319 319 320 321 321 322 323 323 324 324 325 326 327 327 328 328 329 329 331 331 333 333 333 333 335 335 335 335

12.8. Applications--Present and Potential 12.8.1 Structural Applications 12.8.1.1 Cutting Tools 12.8.1.2 Nanocomposites 12.8.1.3 Superplastic Materials 12.8.1.4 Coatings 12.8.2 Magnetic Applications 12.8.3 Catalysts and Hydrogen Storage Materials 12.8.4 Functional NanostructuresmElectronic Applications 12.9. Concluding Remarks References

337 337 337 338 339 339 340 340 341 341 342

Chapter 12 Nanostructured Materials C. SURYANARAYANA and C. C. KOCH

12.1. I N T R O D U C T I O N

The field of nanostructured (or nanocrystalline, or nanophase) materials as a major identifiable activity in materials science results to a large degree from the work of Gleiter (1981, 1989) and coworkers in the 1980's who synthesized ultrafine-grained materials by the in-situ consolidation of nanoscale atomic clusters. The ultra-small size (<100 nm) of the grains in these nanocrystalline materials can result in dramatically improved-or different--properties from those of conventional grain-size (>1 #m) polycrystalline or single crystal materials of the same chemical composition. This is the stimulus for the present high interest in these materials. The nanocrystalline materials pioneered by Gleiter were preceded by studies of nanoparticles by researchers such as Uyeda (1991). At present the very broad field of nanostructured materials includes (i) nanoparticles, (ii) nanocrystalline materials, and (iii) nanodevices. The potential applications for the various kinds of nanoscale materials include dispersions and coatings, high surface area materials, functional nanostructures (e.g., optoelectronic devices, biosensors, nanomachines) and bulk nanostructured materials for structural or magnetic applications. This chapter will focus on the preparation, structure, properties, and possible applications of bulk nanostructured materials.

12.2. C L A S S I F I C A T I O N

Nanostructured materials are single-phase or multi-phase polycrystals, the crystal size of which is of the order of a few (typically 1 to 100) nanometers in at least one dimension. Thus, depending on the dimensions in which the length scale is nanometers, they can be classified into (a) nanoparticles, (b) layered or lamellar structures, (c) filamentary structures, and (d) bulk nanostructured materials. Nanoparticles are essentially atom clusters and can be considered zero-dimensional (0-D) in nature (Siegel 1994). A layered or lamellar structure is a one-dimensional (l-D) nanostructure in which the magnitudes of length and width are much greater than the thickness that is only a few nanometers in size. One can also visualize a two-dimensional (2-D) nanostructure that can be termed filamentary and in this the length is substantially larger than width or diameter, which are of nanometer dimensions. The most common of the nanostructures, however, is basically equiaxed (all the three dimensions are of nanometer size) and are termed nanostructure crystal313

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Table 12.1. Classificationof nanocrystalline materials Dimensionality

Designation

Typical method(s) of synthesis

Zero-dimensional (0-D) One-dimensional (I-D)

Clusters Layered (lamellar)

Two-dimensional (2-D) Three-dimensional (3-D)

Filamentary Crystallites (equiaxerd)

Sol-gel method Vapor deposition Electrodeposition Chemical vapor deposition Gas condensation Mechanical alloying

lites (three-dimensional (3-D) nanostructures). The nanostructured materials may contain crystalline, quasicrystalline, or amorphous phases and can be metals, ceramics, polymers or composites. If the grains are made up of crystals, the material is called nanocrystalline. On the other hand, if they are made up of quasicrystalline or amorphous (glassy) phases, they are termed nanoquasicrystals or nanoglasses, respectively (Suryanarayana 1995). Table 12.1 shows this classification and Fig 12.1 illustrates the four types of nanostructures schematically. Amongst the above, maximum attention has been paid to the synthesis, consolidation, and characterization of the 3D-nanostructured crystallites followed by the 1D-layered nanostructures. While the former are expected to find applications based on their high strength, improved formability, and a good combination of soft magnetic properties, the latter are targeted for electronic applications. Relatively few investigations have been carried out on the 2D-filamentary nanostructures and it is only recently that zero-dimensional clusters are being investigated to tailor the optical properties.

12.3. P R E P A R A T I O N

In principle, any method capable of producing very fine grain-sized polycrystalline materials can be utilized to produce nanocrystalline materials. The grain size, morphology, and texture can be varied by suitably modifying/controlling the process variables in these methods. If a phase transformation is involved, e.g. liquid-to-solid or vapor-to-solid, then steps need to be taken to increase the nucleation rate and decrease the growth rate during formation of the product phase. We shall now describe the different methods of preparing nanostructured materials starting from the vapor, liquid, or solid phases. Table 12.2 summarizes the commonly used methods to synthesize nanocrystalline materials. The full details of the different methods can be found in Gleiter (1989), Siegel (1991), Suryanarayana (1995), and Koch and Suryanarayana (1999). Therefore, only the most commonly used methods will be described here.

12.3.1 Inert gas condensation

Vapor condensation has been known to produce very fine-grained or amorphous alloys depending on the substrate temperature and other operating conditions. Thus, this tech-

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Figure 12.1. Schematic representation of the four types of nanocrystalline materials (after Siegel 1994).

Table 12.2 Methods to synthesize nanocrystalline materials Starting phase

Technique

Nature of product

Vapor

Inert gas condensation Physical vapor deposition--Evaporation and sputtering Plasma processing Chemical vapor condensation Chemical reactions Rapid solidification Electrodeposition Chemical reactions Mechanical attrition Devitrification Spark erosion Sliding wear

3-D 1-D 3-D 3-D, 2-D 3-D 3-D l-D, 3-D 3-D 3-D 3-D 3-D 3-D

Liquid

Solid

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nique was originally used to synthesize small quantities of nanostructured pure metals. A number of variants have also been subsequently developed. This method consists of evaporating a metal (by resistive heating, radio-frequency heating, sputtering, electron beam, laser/plasma heating, or ion sputtering) inside a chamber that was evacuated to a very high vacuum of about 10 -7 torr and then backfilled with a low pressure inert gas, typically a few hundred Pascals of helium. The evaporated atoms collide with the gas atoms inside the chamber, lose their kinetic energy, and condense in the form of small, discrete crystals of loose powder. Convection currents, generated due to the heating of the inert gas by the evaporation source and cooled by the liquid nitrogen-filled collection device (cold finger), carry the condensed fine powders to the collector device, from where they can be stripped off by moving an annular teflon ring down the length of the tube into a compaction device. Compaction is carried out in a twostage piston-and-anvil device initially at low pressures in the upper chamber to produce a loosely compacted pellet, which is then transferred in the vacuum system to a highpressure unit where final compaction takes place. The scraping and compaction processes also are carried out under ultra-high vacuum conditions to maintain cleanliness of the particle surfaces (and subsequent interfaces) and also to minimize the amount of any trapped gases (Fig. 12.2) (Siegel 1990). This method produces equiaxed (3-D) crystallites. The crystal size of the powder is typically a few nanometers and the size distribution is narrow. The crystal size is dependent upon the inert gas pressure, the evaporation rate, and the gas composition. Extremely fine particles can be produced by decreasing either the gas pressure in the chamber or the evaporation rate and by using light (such as He) rather than heavy inert gases (such as Xe). Nanocrystalline alloys can be synthesized by evaporating the different metals from more than one evaporation source. Oxides, nitrides, carbides, etc. of the metals can be synthesized by filling the chamber with oxygen or nitrogen gases or by maintaining a carbonaceous atmosphere. High densities of as-compacted samples have been measured with values of about 75 to 90% of bulk density for metal samples. Sputtering (Terauchi et al. 1995), Electron Beam Vapor Deposition (Bickerdike et al. 1984-85), Plasma Processing (Uda 1992, Chou and Phillips 1992), and Chemical Vapor Condensation (Chang et al. 1994) techniques have also been used to synthesize nanostructured materials.

12.3.2 Rapid solidification

If the cooling rate (or the undercooling) during rapid solidification of metallic melts is increased to high values, then it is possible to increase the nucleation rate (and decrease the growth rate) of the solid phase and consequently the grain size of the resultant product will be in the nanometer range. But, this prediction has not been realized in practice, except in some special situations. A single-phase HfNi5 nanocrystalline alloy (Lu et al. 1995) has been produced directly by rapid solidification. More frequently, nanocomposites have been produced by this method during rapid solidification of metallic melts or during devitrification of metallic glasses produced by rapid solidification (see below).

317

Nanostructured Materials LIQUID NITROGEN

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FUNNEL

7

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~ ~

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\\

- PISTON

LOW PRESSURE COMPACTION UNIT

HIGH PRESSURE COMPACTION UNIT

\\\\N!

Figure 12.2. Schematicrepresentation of a gas condensation chamberfor the synthesis of nanocrystalline materials (after Siegel 1990).

12.3.3 Electrodeposition Electrodeposition of multilayered metals can be achieved using either two separate electrolytes or much more conveniently from one electrolyte by appropriate control of agitation and the electrical conditions (particularly voltage) (Lashmore and Dariel 1988). Also, 3-D nanostructure crystallites can be prepared using this method by utilizing the interference of one ion with the deposition of the other. These processes (a) can be applied to the synthesis of pure metals, alloys, and composites, (b) have high production rates, (c) have few size and shape limitations, and (d) require a low initial capital investment. Erb (1995) and his collaborators have extensively used this process to study the synthesis and properties of 3-D nanocrystalline materials.

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12.3.4 Chemical reactions

WC-Co nanocrystalline composites have been produced on a commercial basis using the "spray conversion processing method" (Kear and McCandlish 1993, Kear and Strutt, 1995) starting from aqueous solution precursors such as ammonium metatungstate (NH4)6(H2W12040)4H20) and COC12, (Co(CH3COO)2), or cobalt nitrate (Co(NO3)2). The solution mixture is aerosolized and rapidly spray dried to give extremely fine mixtures of tungsten and cobalt complex compounds. This precursor powder is reduced with hydrogen and reacted with carbon monoxide in a fluidized-bed reactor to yield nanophase cobalt/tungsten carbide powder. The tungsten particles are 20-40 nm in size. A typical powder particle consists of a hollow, porous 75 # m sphere containing hundreds of millions of WC grains in a cobalt matrix. To prevent grain growth of tungsten, additions of inhibitors such as VC and Cr3C2 are made as binders during the sintering steps. Recently, vanadium is being introduced into the starting solution itself to achieve a more uniform distribution in the powder mixture. The process parameters are being further optimized and since the process is fully integrated, on-line control is also being planned.

12.3.5 Mechanical attrition

Processing/synthesis of nanocrystalline materials starting from the solid state has been a popular method because of the possibility of producing large quantities of material. Even amongst the many possible ways, mechanical attrition has attracted the maximum attention. Mechanical attrition produces nanostructured materials by the structural decomposition of coarser-grained structure as a result of severe plastic deformation. Mechanical attrition consists of repeated welding, fracturing, and rewelding of powder particles in a dry high-energy ball mill until the composition of the resultant powder corresponds to the percentages of the respective constituents in the initial charge. These processes have produced nanocrystalline structures in pure metals, intermetallic compounds, and immiscible alloy systems. It has been shown that nanometer-sized grains can be obtained in almost any material after sufficient milling time. The grain sizes were found to decrease with milling time down to a minimum value that appeared to scale inversely with the melting temperature. Koch (1993, 1997) has recently summarized the characteristics and properties of the nanocrystalline materials produced by mechanical attrition. Contamination of the powder from the milling tools is of concern in this method. (See the Chapter on "Mechanical Alloying" by C. Suryanarayana in this book). In recent years, the process of severe plastic deformation of bulk solids (by the EqualChannel-Angular technique) has been shown to produce ultra fine-grained structures. Even though the grain size is strictly not in the nanometer range (it is usually about 0.3-0.5 #m), there has been considerable amount of research work on the structure and properties of materials produced by this method (Valiev 1994). This has been essentially due to the possibility of producing bulk high-purity materials possessing sub-micron grain sizes.

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12.3.6 Devitrification

Many processes such as rapid solidification, mechanical attrition, electrodeposition, and vapor deposition can produce amorphous (glassy) alloys. Controlled crystallization of these amorphous alloys (by increasing the rate of nucleation and decreasing the rate of growth) leads to the synthesis of nanostructured materials. In fact, the most common method to produce nanocrystalline magnetic materials has been to obtain an amorphous phase by rapidly solidifying the melt of appropriate composition and then crystallizing the glassy phase at a relatively low temperature. These materialsmreferred to as FINEMET-were first investigated by Yoshizawa et al. (1988) and this method has now become an established practice to study the structure and properties of nanocrystalline magnetic materials (Lu 1996). This simple devitrification method has been commonly employed to study the magnetic properties of nanocrystalline materials because it can produce (a) porosity-free samples, (b) samples with different grain sizes by controlling the crystallization parameters, and (c) large quantities of material. Furthermore, since no artificial consolidation process is involved (like in inert gas condensation, mechanical alloying, or plasma processing, where fine powders are involved), the interfaces are clean and the product is dense. Spark Erosion (Kodama et al. 1994) and Sliding Wear (Ganapathi et al. 1991) have also been used to synthesize nanocrystalline materials. Nanofibers have also been produced by several methods including filling of carbon or boron carbide nanotubes with metal oxides and performing subsequent reactions to prepare nanorods, template synthesis in the cylindrical pores of a membrane, decomposition of selected hydrocarbons, and by locking in the mesophasic structures formed by diblock copolymers (Liu 1997). Apart from the methods described above, other methods have also been used to synthesize nanostructured materials. These include co-precipitation, laser ablation, hydrothermal pyrolysis, thermophoretic forced flux system, autoignition, and sonochemical synthesis. Additionally, variations of the above-mentioned methods and some laboratory improvisations are also being adopted to achieve certain morphology, grain size, or nature of the nanocrystalline particles. Several companies all over the world have started producing nanocrystalline powders on a commercial scale. According to a recent estimate (Rittner and Abraham 1998), in the United States alone, over 50 small and large companies are active at various levels in the development and production of nanocrystalline materials. More than a dozen of these businesses are involved in the manufacture of nanostructured materials on an industrial scale.

12.4. S T R U C T U R E

A schematic representation of a hard-sphere model of an equiaxed nanocrystalline metal is shown in Fig. 12.3 (Gleiter 1989). Two types of atoms can be distinguished: crystal atoms with nearest neighbor configuration corresponding to the lattice and the boundary atoms with a variety of interatomic spacings, differing from boundary to boundary. A nanocrystalline metal contains typically a high number of interfaces (~6 x 1025 m -3 for

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Figure 12.3. Schematic representation of an equiaxed nanocrystalline metal distinguishing between the atoms associated with the individual grains (filled circles) and those constituting the grain boundary network (open circles) (after Gleiter 1989).

a 10 nm-grain size) with random orientation relationships, and consequently, a substantial fraction of the atoms lies in the interfaces. Assuming that grains have the shape of spheres or cubes, the volume fraction of atoms associated with the boundaries can be calculated as

C = 3A/d

(12.1)

where A is the average grain boundary thickness and d is the average grain diameter. Thus, the volume fraction of atoms in the grain boundaries can be as much as 50% for 5-nm grains and decrease to about 30% for 10-nm grains and 3% for 100-nm grains. Thus, nanocrystalline metals can be considered to consist of two structural componentsmthe numerous, small crystallites with long-range order and different crystallographic orientations constituting the "crystalline" component and a network of intercrystalline regions, the structure of which differs from region to region; this will be referred to as the "interfacial" component. The interatomic spacings in the interfacial component have a wide distribution and further the average atomic density is considerably less than the crystal density depending on the type of chemical bonding between the atoms.

12.4.1 Microstructure

In nanocrystalline single-phase alloys and pure metals, the most important structural parameter is the grain size. The properties of materials are mostly dependent on the grain size and therefore, an accurate determination of the grain size is important. Both direct (elec-

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tron microscopic) and indirect (scattering) techniques have been employed to determine the grain sizes. Transmission electron microscopy (TEM) techniques (especially, the high-resolution TEM studies) are ideal to directly determine the grain sizes of nanocrystalline materials using the dark-field technique. The width of the Bragg reflection in an X-ray (large-angle) diffraction pattern can provide grain (or crystal) size information after the appropriate corrections (for instrumental and strain effects) are incorporated. The TEM techniques can clearly indicate whether there is a distribution of grain sizes and it is also possible to obtain a grain size histogram by measuring the grain sizes and counting the number of grains. On the other hand, the X-ray diffraction technique gives only the average crystal size.

12.4.2 Atomic structure of the crystal lattice The structure of the grains (crystallites) in nanocrystalline materials has been normally accepted to be the same as in a coarse-grained material. Consequently, there have not been many investigations into this aspect. High-resolution TEM experiments have indicated that nanocrystalline materials consist of small crystallites of different crystallographic orientations separated by grain boundaries. Even though not frequently reported, the grains contained different crystalline defects such as dislocations, twin boundaries, multiple twins, and stacking faults. Sui and Lu (1994) and Liu et al. (1993) used X-ray diffraction and Mrssbauer spectroscopy techniques and observed that even though both Ni3P and Fe2B (obtained by crystallization of the Ni-P and Fe-Mo-Si-B amorphous alloys) have the tetragonal crystal structures in both coarse-grained and nanocrystalline states, the "a" parameter is always larger and the "c" parameter smaller in the nanocrystalline state. Further, these values changed with a change in grain size. The relative deviations of the parameters reached as much as Aa = 0.22% and Ac = - 0 . 2 4 % for the Fe2B lattice with a grain size of 25 nm and Aa = 0.21% and Ac - --0.13% for the Ni3P lattice with a grain size of 6 nm. The unit cell volume increased in both the cases with respect to the equilibrium values for coarse-grained materials. Results different from the above were also reported. While Cr has a bcc structure under equilibrium conditions, it had, in the nanocrystalline state, the A15 structure under high-purity conditions, and the bcc structure in presence of oxygen impurities. Weismfiller (1996) has reported that the lattice parameters of nanocrystalline metals are not significantly different from those of coarse-grained materials.

12.4.3 Atomic structure of the grain boundaries The structure of the grain boundaries has received a lot of attention and has been discussed extensively in the literature, especially to decide whether it is different in the nanocrystalline and coarse-grained materials of the same composition. The conclusions differ and some believe that the structure is fundamentally different in both the types of materials while others believe that it is the same. The present status of the structure of grain bound-

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aries can be found in some recent reviews (Siegel 1992, Birringer 1994, Weismtiller et al. 1995, Suryanarayana 1995, Weismtiller 1996). Gleiter (1989), Birringer (1989), and Ouyang et al. (1993) raised the possibility that the structure of the grain boundaries in nanocrystalline materials might have unique characteristics owing to the constraints imposed on the grain boundary atoms by their synthesis via cluster consolidation. Based on a recent computer simulation study, Phillpot et al. (1995) suggested that the structure of grain boundaries in nanostructured materials is rather dissimilar from the structures found in glasses (which have short-range order) and that derived from bicrystal studies. Based on these studies, it has been suggested that the grain boundaries in nanocrystalline materials are fundamentally different from those in conventional coarse-grained materials. Siegel and Thomas (1992) also simulated the structure of grain boundaries in nanocrystalline materials and concluded that the grain boundaries in nanocrystalline materials and conventional polycrystalline materials have similar structures. Based on their high-resolution TEM studies, Horita et al. (1996) also concluded that there was no evidence for a gas-like or widely diffuse structure in the grain boundaries of sub-micron grain size materials. The details of the atomic structure of the grain boundaries in a given nanocrystalline material may vary as a function of the time-temperature history of the sample (crystal size and external pressures). From the above it is apparent that the question of whether the structure of the grain boundaries in nanocrystals is different from the coarse-grained crystals is not answered clearly. Direct evidence for an accurate structure of the grain boundary is difficult to obtain in view of the relaxation at surfaces; additional results from FIM and STM may shed new light. Additionally, it would be useful to compare the structure of the grain boundaries in the same material using the same characterization technique, but produced by different methods to resolve some of the conflicting data available in the literature. Ping et al. (1995) studied the microstructure of nanocrystalline materials obtained by two different techniques--inert gas condensation and devitrification of glassy phases. They reported that much less defect structure was detected inside the crystals or interfaces in a devitrified sample than in an inert-gas-condensed and in-situ consolidated samples. They have also reported that while both ordered and disordered regions and nanovoids were present in the grain boundaries of inert gas condensed palladium, the samples obtained by devitrification of Ti-Ni-Si and Fe-Mo-Si-B amorphous alloys have only ordered structures. 12.4. 4 Triple junctions and higher-order grain junctions

In recent years, it has been repeatedly pointed out that when the grain size of the material is very small, triple junctions and higher-order grain junctions become an important component of the microstructure. Thus the total intercrystalline region consisting of grain boundaries (interface planes shared by two grains), triple junctions, (intersection lines where three grains meet), and four grain junction points is more important than just the grain boundaries. The grain boundary area is inversely proportional to the grain size while the triple line length is proportional to the inverse square of the grain size. Thus, the contribution of triple junctions and higher-order grain junctions becomes significant at

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Materials

323

very small grain sizes. It has been suggested that the contribution of triple junctions and higher-order junctions is not very significant at grain sizes > 10 nm (most materials produced to-date are in this size range) and therefore they may not exert much influence on the properties of nanocrystalline materials.

12.5. S T A B I L I T Y

Knowledge of the thermal stability of nanocrystalline materials is important for both technological and scientific reasons. From a technological point of view, the thermal stability is important for consolidation of nanocrystalline powders without coarsening the microstructure. That is, many methods for synthesis of nanocrystalline materials result in powder products. These powders then require consolidation by techniques such as hot forging to obtain useful bulk material. The goal of such powder consolidation is to attain essentially 100% theoretical density while preventing or minimizing grain growth of the nanometer-sized grains. Another processing method that requires information on grain growth is the devitrification of amorphous precursors (Lu 1996). In order to design an appropriate heating schedule to provide the desired nanometer grain size, knowledge of the nucleation rate and grain growth rate are both needed. The basic scientific question regarding grain growth of nanocrystalline materials, as for all phenomena for nanocrystalline materials, is whether the behavior involves "new physics" or is simply the expected grain-size-dependent behavior extrapolated to nanometer grain sizes.

1 2 . 5 . 1 G r a i n g r o w t h in n a n o c r y s t a l l i n e m a t e r i a l s

Grain growth occurs in polycrystalline materials to decrease the system energy by decreasing the total grain boundary energy. Under ideal conditions grain growth kinetics should obey parabolic kinetics (Atkinson 1988, Vandermeer and Hu 1994), namely D 2 -

D 2 -- kt

(12.2)

where D is the length scale, e.g., mean grain diameter, after annealing for time t, k is a temperature-dependent rate parameter and Do is the length scale or the grain size at t = 0. However, most experimental data for isothermal grain growth do not obey this ideal behavior and experimenters often analyze grain growth data with the expression:

D1/n

1/n

- Do

-- k t

(12.3)

where n is an empirical constant typically <0.5 (the grain growth exponent is also often defined as n I = 1/n and is then typically >2). By applying an Arrhenius-type equation of the form

0exp{

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to the rate parameter k one gets access to the activation energy Q, which is frequently used to deduce the grain growth mechanism. Another method of obtaining Q for nanocrystalline materials is the so-called Kissinger method in which differential scanning calorimetry (DSC) is used to analyze isochronal grain growth at different heating rates. We shall describe the isothermal grain growth kinetics data by the use of n and Q. However, first we will review several aspects of thermal stability which have been observed in nanocrystalline materials.

12.5.2 Grain growth at ambient temperature Significant grain growth, that is doubling of the grain size in 24h, has been observed in a number of pure, relatively low melting elements such as Sn, Pb, A1, and Mg (Birringer 1989). Because of the large amount of grain boundary enthalpy stored in nanocrystalline materials a high driving force for grain growth is expected. GUnther et al. (1992) and Kumpmann et al. (1993) studied grain growth in pure Cu, Ag, Pd and Ag alloys and found grain growth occurring at much lower temperatures than those observed for recrystallization of the elements after heavy cold deformation. In fact grain growth in Cu and Pd was observed even at room temperature. In some of these cases grain growth was observed to be abnormal. Gertsman and Birringer (1994) observed abnormal grain growth in nanocrystalline Cu after holding the samples for times of more than a month at room temperature. The abnormally coarse grains exhibited a range of sizes, most were <1/zm but some were >2 #m. The nanometer grains surrounding the coarse grains were 10-50 nm. The general explanation given for abnormal grain growth is an inhomogeneous structure for the as-processed samples such that grain growth inhibitors such as pores, impurities, or even grain boundary structure are not evenly distributed and rapid grain growth can occur where such inhibitors are absent due to the large driving force and high grain-boundary mobility.

12.5.3 Examples of grain growth inhibition Significant stabilization of nanometer grain structures has been observed in many multicomponent materials. It has been observed in several systems that on heating a metastable nanocrystalline solid solution alloy, limited grain growth occurs while solute atoms also segregate to the grain boundaries. The grain boundary segregation then may lower the specific grain boundary energy sufficiently to reduce the driving force and limit further grain growth. This effect was illustrated clearly in Pdl-xZrx alloys (Krill et al. 1995) where x = 0.1 is an equilibrium solid solution and x = 0.2 is a metastable supersaturated solid solution. The 10 at.% Zr nanocrystalline alloy showed significant grain growth with an onset temperature of 325~ while the 20 at.% Zr nanocrystalline alloy had little or no grain growth at temperatures up to 500~ While it is not possible to separate possible kinetic effects, such as solute drag, from this thermodynamic stabilization, the much slower grain growth in the 20 at.% Zr supersaturated alloy strongly suggests that the thermodynamic stabilization is dominant here. Solute segregation to nanocrystalline grain boundaries being responsible for grain size stabilization in supersaturated metastable nanocrystalline

Nanostructured Materials

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Figure 12.4. Time exponent, n, for isothermal grain growth of various nanocrystalline materials as a function of the normalized annealing temperature (after Malow and Koch 1996).

solid solutions has also been observed in Nb-Cu (Abe et al. 1992) and Fe-Cu (Eckert et

al. 1993) alloys. A number of factors can influence grain boundary mobility in nanocrystalline alloys. These include solute drag, pore drag, second phase (Zener) drag, and chemical ordering.

12.5.4 Isothermal grain growth kinetics The isothermal grain growth kinetics will be considered using Eqs. (12.3) and (12.4) to define the grain growth exponent n (_<0.5) and the activation energy Q, which also can be obtained by the Kissinger Analysis. Isothermal grain growth data from the literature for nanocrystalline materials has been analyzed by fitting to Eq. (12.3) and is presented in Fig. 12.4 where n is plotted against the reduced annealing temperature T~ Tm (Malow and Koch 1996). Tm is the melting or liquidus temperature of the given material. The lines connecting data for a given material/reference are drawn only to aid the reader. While there is a good deal of variation from n < 0.1 to n ~ 0.5 the general trend is for n to increase toward the ideal 0.5 value with increasing temperature, T~ Tin. These results follow the trend observed in coarse-grained materials (Hu and Rath 1970). In fact, the ideal value of n = 0.5 has only been observed in high purity elements such as Cd, Fe, and Sn at high values of T~ Tm (Vandermeer and Hu 1994).

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C. Suryanarayana a n d C. C. Koch

Often the activation energy for grain growth is compared to the activation energy for lattice (Q1) or grain boundary (Qgb) diffusion in order to help determine the underlying mechanisms involved. Activation energy data for grain growth in nanocrystalline materials has been summarized by Malow and Koch (1996). The data were compared to diffusion data, with the rough estimate for activation energy for grain boundary diffusion (Qgb = 0.5 Q1), where data does not exist. Most of the results on Q for nanocrystalline materials compare better with Qgb than with QI. This is consistent with most data for grain growth in conventional grain size materials. There are several exceptions to this general trend (GUnther et al. 1992, Kumpmann et al. 1993, Gertsman and Birringer 1994). The activation energy for grain growth at T > 775 K in nanocrystalline Fe was found to be 248 kJ/mol which is comparable to that for grain growth in zone refined pure Fe with conventional grain sizes (> 10/zm) (Vamdermeer and Hu 1994) and with Q1 for Fe. But at T < 775 K, the activation energy was 125 kJ/mol which is lower than both Qgb and Q1 for Fe. This suggests that grain growth in nanocrystalline Fe at T > 775 K is similar to that in conventional grain size pure Fe, but grain growth in nanocrystalline Fe at T < 775 K is due to another mechanism (Malow and Koch 1997). In conclusion it can be said that in order to learn about the mechanisms that govern the thermal stability/grain growth in nanocrystalline materials it is important to keep factors in mind that might interfere with grain growth. Specifically it would be desirable to avoid porosity, contamination and multiple phases (or precipitation) in future investigations. Furthermore it would be of advantage to compare the resultant activation energies to diffusion data for nanocrystalline materials, if available, as it has been done by H6fler and Averback (1990).

12.6. PARTICULATE CONSOLIDATION Widespread application of nanocrystalline materials requires production of the powder in tonnage quantities and also efficient methods of consolidating the powders into bulk shapes. All the consolidation methods generally used in powder metallurgy processes can also be used for nanocrystalline materials. However, because of the small size of the powder particles (typically a few microns, even though the grain size is only a few nanometers), some special precautions need to be taken to minimize their activity and the high level of interparticle friction. Conventional consolidation of powders to full density through processes such as hot extrusion and hot isostatic pressing normally requires use of high pressures and high temperatures for extended periods of time. Unfortunately, however, this results in significant coarsening of the nanometer-sized grains and consequently the benefits of nanostructure processing are lost. Therefore, novel and innovative methods of consolidating nanocrystalline powders are required. Consolidation of nanocrystalline powders has been achieved by electro-discharge compaction (Okazaki 1993), plasma-activated sintering (Groza 1993, 1994), shock (explosive) consolidation (Korth and Williamson 1995, Suryanarayana et al. 1997), hotisostatic pressing (Suryanarayana et al. 1997, Haji-Mahmood and Chumbley 1996, He and Ma 1996a), Ceracon processing (Suryanarayana et al. 1997), hydrostatic extrusion (Liang et al. 1996), strained powder-rolling (Liang et al. 1996), and sinter-forging (He

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and Ma 1996b). Optimization of the consolidation parameters is important because retention of the nanostructures requires use of low consolidation temperatures while achieving full (theoretical) density requires use of high temperatures. However, it should be noted that because of the increased diffusivity in nanocrystalline materials, sintering (densification) takes place at a temperature much lower than in coarse-grained materials. This is likely to reduce the grain growth. (See also the chapter on "Powder Consolidation" by J.R. Groza in this book). Full consolidation of nanocrystalline materials has also been achieved by modifying/improving some of the other existing processes. Although it is typically difficult to obtain full density while retaining the nano-scale microstructure, the high driving force and enhanced kinetics due to large curvature effects facilitate full densification of nanocrystalline materials at temperatures consistently below those of coarse-grained materials of the same composition by a few hundred degrees. Retention of fine grain sizes and elimination of sintering aids (especially for ceramic materials) are the specific advantages of the lower processing temperatures. 12.7. P R O P E R T I E S

Because of the very fine grain sizes, nanocrystalline materials exhibit a variety of properties that are different and often considerably improved in comparison with those of conventional coarse-grained polycrystalline materials. These include increased strength~ardness, enhanced diffusivity, reduced density, higher electrical resistivity, increased specific heat, higher thermal expansion coefficient, lower thermal conductivity, and superior soft magnetic properties. 12.7.1 Diffusion and sinterability

Since nanocrystalline materials contain a very large fraction of atoms at the grain boundaries, the numerous interfaces provide a high density of short-circuit diffusion paths. Consequently, they are expected to exhibit an enhanced diffusivity in comparison to single crystals or conventional coarse-grained polycrystalline materials with the same chemical composition (Horvath 1989, H6fler et al. 1993). This enhanced diffusivity can have a significant effect on mechanical properties such as creep and superplasticity, ability to efficiently dope nanocrystalline materials with impurities at relatively low temperatures, and synthesis of alloy phases in immiscible metals and at temperatures much lower than those usually required in other systems. The measured diffusivities in nanocrystalline copper are about 14 to 20 orders of magnitude higher than lattice diffusion and about 2 to 4 orders of magnitude larger than grain boundary diffusion. For example, the measured diffusivity at room temperature is 2.6 x 10 -20 m2/s for 8 nm-grain-sized copper samples compared to 4.8 x 10 -24 for grain boundary diffusion and 4 x 10 -40 for lattice diffusion (Schumacher et al. 1989). The increased diffusivity (and consequently the reactivity) leads to increased solid solubility limits, formation of intermetallic phases (at temperatures much lower than those required for coarse-grained materials and sometimes new phases) and increased sinterability of nanocrystalline powders.

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Solid solubility limits are usually enhanced when the material is in the nanocrystalline state. In extreme cases, solid solutions can also form in alloy systems that exhibit miscibility gaps both in the liquid and solid states. Typical examples of this phenomenon are the Ag-Fe, Ti-Mg, and Cu-Fe systems (Suryanarayana 1995). While the room temperature solid solubility of Mg in Ti under equilibrium conditions is <0.2 at.%, it could be increased to 38 at.% when titanium grains are in the nanometer range. Similarly, the solid solubility of Bi in Cu is < 10 -4 at.% under equilibrium and this has been increased to 4 at.% in the nanocrystalline state. Further, the solid solubility of Hg in nanocrystalline Cu has been increased to 70 at.%, against an equilibrium value of < 1 at.%. The enhanced solid solubility limits in these systems have been explained on the basis of elastic strains at the interfaces; demonstrated to be the reason both theoretically and experimentally for various alloys. The high diffusivity of nanocrystalline materials leads to alloying by diffusion along the grain boundaries resulting in the formation of stable and metastable phases at relatively low temperatures. For example, formation of the Pd3Bi intermetallic has been found to occur at 120~ a temperature much lower than normally observed. Similarly, nanocrystalline mixtures of copper and erbium crystallites formed the equilibrium CuEr compound by compaction. Two new polymorphs of erbium oxide were also detected in the nanocrystalline state. Another important consequence of the increased diffusivity is that sintering of nanocrystalline powders can occur at temperatures much lower than those required for sintering coarse-grained polycrystalline powders. Significant improvements in the sinterability have been obtained in nanocrystalline TiO2 (rutile) in comparison to conventionally synthesized coarse-grained rutile (Mayo et al. 1990). For example, TiO2 with a grain size of 12 nm could be sintered at ambient pressures at temperatures 400 to 600~ lower than that required for ball-milled 1.3 #m powder, and without the need for any compacting or sintering aid, such as polyvinyl alcohol, which is usually required. Similarly, the HIP temperature for full consolidation could be reduced by about 400~ for nanocrystalline titanium aluminides (Suryanarayana et al. 1997).

12. 7.2 Mechanical properties 12.7.2.1 Elastic properties

The early measurements of the elastic constants on nanocrystalline materials prepared by the inert gas condensation method gave values, for example for Young's Modulus, E, which were significantly lower than values for conventional grain size materials (Suryanarayana 1995). While various reasons were given for the lower values of E, it was suggested (Krstic et al. 1993) that the presence of extrinsic defects, e.g., pores and cracks, was responsible for the low values of E in nanocrystalline materials compacted from powders. This conclusion was based on the observation that nanocrystalline NiP produced by electroplating with negligible porosity levels had an E value comparable to fully dense conventional grain size Ni (Wong et al. 1993).

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12.7.2.2 Hardness and strength

The hardness and strength of conventional grain size materials (d > 1 #m) can be expressed by the empirical Hall-Petch equation ~o = ai + k d - 1 / 2 where ~0 = yield stress, cri = friction stress opposing dislocation motion, k = constant, and d - grain diameter. Similar results are obtained for hardness with H0 = Hi + kd -1/2. Most of the data on nanocrystalline materials have been hardness data while some tensile test data are becoming available as well. Several recent reviews have summarized the mechanical behavior of nanocrystalline materials (Morris 1998, Siegel and Fougere 1994, Siegel 1997, Weertman and Averback 1996). It is clear that as grain size is reduced through the nanoscale regime (< 100 nm) hardness typically increases with decreasing grain size and can be 2 to 7 times harder for pure nanocrystalline metals (10 nm grain size) compared to large grained (> 1 #m) metals. Indeed, if the Hall-Petch plot is extrapolated from large grain to nanograin regions very high yield stresses--approaching the theoretical limit for shear of atomic planesmmight be reached. However, such high predicted yield stresses or hardness values are not observed. The experimental results of hardness measurements show different behavior for dependence on grain size at the smallest nanocrystalline grains (<20 nm) including a positive slope ("normal" Hall-Petch behavior), essentially no dependence (~zero slope), and in some cases a negative slope. Most data that exhibit the negative Hall-Petch effect at the smallest grain sizes have resulted from nanocrystalline samples that have been annealed to increase their grain size. It is suggested (Siegel and Fougere 1994) that thermally treating nanophase samples in the as-produced condition may result in such changes in structure as densification, stress relief, phase transformations, or grain-boundary structure changes, all of which may lead to the observed negative Hall-Petch behavior.

12.7.2.3 Ductility and toughness

It is well known that grain size has a strong effect on ductility and toughness of conventional grain size (> 1 #m) materials. For example, the ductile/brittle transition temperature in mild steel can be lowered about 40~ by reducing the grain size by a factor of 5. The results of ductility measurements on nanocrystalline metals are mixed and are sensitive to flaws and porosity, surface finish, and method of testing (e.g., tension vs. compression). While conventional grain size pure copper has values of % elongation of 60 (annealed) to 14 (cold drawn) nanocrystalline copper gives <4% total elongation values (Eastman et al. 1997). Miniaturized disk bend tests, which also provide tensile stress components, exhibited brittle behavior for 27 nm grain size Cu but ductile behavior for nanocrystalline Cu with grain sizes >40 nm (Gertsman et al. 1994). Compression tests on nanocrystalline Cu, however, have shown much larger ductilities with % elongations of >20% (Suryanarayanan et al. 1996) and 12-18% (Eastman et al. 1997) observed. The yield stresses were also about twice those observed in tension. While it is likely that flaws and porosity present in many nanocrystalline samples seriously affect the results of mechanical tests and may be partly responsible for the asymmetry of results in compression vs. tension, the nature of the deformation process may also be important. In any event, the ductility of nanocrystalline materials that exhibit ductile behavior with conventional grain sizes

C. Suryanarayana and C. C. Koch

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typically show reduced ductilitymsometimes brittle behaviormat the smallest nanometer grain sizes as illustrated in Fig. 12.5. This is presumably due to the inability of usual dislocation generation and motion to occur at these nanometer grain sizes. An intriguing suggestion based on observations of ductile behavior of brittle nanocrystalline ceramics at low temperatures (Karch et al. 1987) is that brittle ceramics or intermetallics (Bohn et al. 1991) might exhibit ductility with nanometer-sized grains. Karch et al (1987) observed apparent plastic behavior in compression in nanocrystalline CaF2 at 80~ and in nanocrystalline TiO2 at 180~ These observations were attributed to enhanced diffusional creep providing the plasticity at these temperatures where conventional grain size materials would fail in the elastic regime. It was assumed that diffusional creep was responsible for the plasticity and the observations were rationalized with boundary diffusion dominating the behavior such that the strain (creep) rate is: de dt

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at moderate temperatures find creep rates comparable to or lower than corresponding coarse-grain rates. It is suggested that the high fraction of low-energy grain boundaries together with the limitation on dislocation activity by the small grain sizes cause the low creep rates observed. The initial interesting results on nanocrystalline CaF2 and nc TiO2 have not been reproduced per se and it is likely these results were due to extensive porosity in the samples. However, compressive creep tests on 99% dense nanocrystalline TiO2 at 810~ (0.57 Tin) showed total strains of up to 0.6. Thus, there have not been reports for brittle ceramics or intermetallics of the dramatic increases in ductility predicted by Eq. (12.5) at temperatures significantly <0.5 Tm (Morris 1998).

12.7.2.4 Superplastic behavior Superplasticity is the ability of some polycrystalline materials to exhibit very large tensile deformations without necking or fracture. Typically elongations of 100% to > 1000% are considered the defining features of this phenomenon. As grain size is decreased it is found that the temperature at which superplasticity occurs is lowered, and the strain rate for its occurrence is increased. As discussed in Section 12.7.2.3 it was suggested from Eq. (12.5) that creep rates could be enhanced by many orders of magnitude and superplastic behavior might be observed in nanocrystalline materials at temperatures much lower than 0.5 Tm. However, actual creep experiments have not borne out this prediction, but showed creep rates comparable or lower than in coarse-grained samples of the same material. This is presumably why little enhancement in ductility or superplastic behavior has been observed for nanocrystalline materials at temperatures <0.5 Tm. While there is no evidence for "low"-temperature superplasticity in tension in nanocrystalline materials there are several reports of large ductilities exhibited by nanocrystalline materials under compressive loads. These include studies of ceramics (Hahn et al. 1991) and intermetallics (Jain and Christman 1994). 12.7.2.5 Deformation mechanisms in nanoscale materials While there is still limited data on the mechanical behaviormespecially tensile properties--of nanocrystalline materials some generalizations may be made regarding the deformation mechanisms. It is likely that for the larger end of the nanoscale grain sizes, about 50 to 100 nm, dislocation activity dominates for test temperatures <0.5 Tm. As grain size decreases, dislocation activity apparently decreases. The essential lack of dislocations is presumably the result of the image forces that act on dislocations near surfaces or interfaces. The creation of new dislocations is also made difficult as the grain size reaches the lower end of the nanoscale (_< 10 nm). Stresses needed to activate dislocation sources, such as the Frank-Read source, are inversely proportional to the distance between dislocation pinning points. Since nanoscale grains will limit the distance between such pinning points the stresses to activate dislocation sources can reach the theoretical shear stress of a dislocation-free crystal at the smallest grain sizes (,~2 nm). Thus at the smallest grain sizes we may have new phenomena controlling deformation behavior. It has been suggested (Siegel 1997) that such phenomena may involve grain boundary sliding and/or grain rotation accompanied by short-range diffusion assisted healing events.

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Eastman et al. (1997) have carried out mechanical tests using compression, tension and disk-bending (which has tensile components) on nanocrystalline Cu, nanocrystalline A1-A13Zr alloys, and nanocrystalline NiA1. From these studies they concluded that materials such as Cu or A1-A13Zr that are ductile in coarse-grained form, due to dislocation creation and motion, become less ductile in nanocrystalline form as dislocation-based deformation becomes difficult. On the other hand, brittle coarse-grained materials in which dislocation-based deformation is already minimal, such as NiA1, exhibit slightly enchanced ductility in nanocrystalline form. This enhanced ductility is presumably due to the deformation mechanism for nanocrystalline materials, which is not yet well defined but may involve grain rotation and grain boundary sliding. Several examples of deformation by shear banding have been reported for nanocrystalline materials. Carsley et al. (1997) have studied nanocrystalline Fe-10% Cu alloys with grain sizes ranging from 45 to 1680 nm. In all cases deformation in compression proceeded by intense localized shear banding. The stress-strain curves exhibited essentially elastic-perfectly plastic behavior. That is, no measurable strain hardening was observed. Shear banding is also the deformation mode observed in amorphous metallic alloys and amorphous polymers. The deformation shear banding in nanocrystalline Fe-10% Cu was compared to these observations for metallic glasses, amorphous polymers, and coarsegrained polycrystalline metals after significant plasticity and work hardening have taken place. While this suggests a close similarity between deformation in nanocrystalline materials and amorphous materials, not all tensile data on nanocrystalline materials exhibit a lack of strain hardening and the Fe-10% Cu samples of Carsley et al. (1997) showed shear bands even in their larger grained specimens (i.e., about 1000 nm). The existing data on mechanical properties provides the following tentative conclusions, although much work remains to be done to clarify the detailed deformation mechanisms. (i) Except for very small grain nanocrystalline samples (<5 nm) the elastic properties are essentially identical to those of coarse-grained material. (ii) High hardnesses and yield strength values are observed for nanocrystalline materials. (iii) Ductile coarse-grained materials are less ductile, perhaps brittle as nanocrystalline materials. (iv) Brittle coarse-grained materials may exhibit slight ductility when nanocrystalline. Not enough evidence to verify this. (v) Superplasticity has not been observed at low temperatures (<0.4 to 0.5 Tm) for nanocrystalline materials, although it has been observed at somewhat lower temperatures and higher strain rates as grain size decreases. (vi) Shear banding and perfectly plastic behavior have been seen in some nanocrystalline materials analogous to plastic deformation in amorphous materials. (vii) Mechanical twinning may be an alternate deformation mechanism in nanocrystalline materials.

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12.7.3 Magnetic properties 12.7.3.1 Fundamental properties Changes in the magnetic properties of nanoscale ferromagnetic materials can be attributed to their small volume which means they are typically single domain, and to the large fraction of atoms associated with the grain boundaries/interfaces. If the single domain nanoscale ferromagnetic particles (e.g., Fe, Co, or Ni) are put in a non-magnetic matrix such that their magnetic interaction is negligible, a superparamagnetic material is obtained. If the grain sizes are small enough, the structural distortions associated with the surfaces or interfaces can lower the Curie temperature, Tc, and reduce the magnitude of the saturation magnetization Ms. For example, the Ms of 6 nm Fe was reported to be 130 emu/g compared to 220 emu/g for normal polycrystalline Fe. The Curie temperature of 10 nm Gd was decreased about 10 K from that of coarse-grained Gd, and the transition was broader. The significant decreases in Ms (Birringer et al. 1986) for nanocrystalline Fe and nanocrystalline Fe, Ni, and Co (Gong et al. 1991) reported for gas condensed nanocrystalline powders have not been observed for fully dense electrodeposited nanocrystalline Ni (Aus et al. 1992) where Ms decreased less than 5% down to grain sizes of 10 rim. Similarly the Curie temperature was found to be unchanged (Erb et al. 1997). These results were consistent with calculations to determine the effect of the structural disorder associated with various grain boundaries on local magnetic moments. 12.7.3.2 Soft magnetic materials Attractive soft magnetic properties were observed for nanocrystalline Fe-base alloys obtained by crystallization of an amorphous alloy ribbon produced by rapid solidification. In particular, the crystallization of amorphous Fe-Si-B alloys containing Nb and Cu results in the formation of a nanoscale bcc structure and the bcc/amorphous alloys exhibit good soft magnetic properties of 1.2 to 1.4 T for saturation induction (Bs) and 10 • 104 for effective permeability (/~c) at 1 kHz (Yoshizawa et al. 1988). Although good soft magnetic properties are obtained for these "FINEMET" alloys, e.g., Fe73.5Si13.sB9Nb3Cu1, which have 15 nm bcc particles in an amorphous matrix, the relatively low Fe concentration leads to the limitation of Bs to less than 1.4 T. It is desired to have soft magnetic materials with Bs > 1.5 T and/Zc above 105 at 1 kHz for various kinds of power transformers. Higher Fe content amorphous alloys of type Fe-M-B (M = Zr, Hf, or Nb) were subsequently studied and mixed phase nanocrystalline/amorphous structures were found to exhibit excellent soft magnetic properties (Makino et al. 1995). These alloys (the "Nanoperm" family) developed by the Alps Electric Co. of Japan are typified by a composition such as Fe86ZrTB6Cu1 (Makino et al. 1991). The addition of Cu aids in increased nucleation while Zr acts to inhibit growth (as well as allowing for amorphization). The core losses for this material are compared to those for conventional coarse grained Fe-3.5% Si and amorphous Fe-9% Si-13 %B in Fig. 12.6. 12.7.3.3 Hard magnetic materials The first attempts to produce nanoscale microstructures to enhance the magnetic properties of the Nb-Fe-B permanent magnetic materials used mechanical alloying of blended

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elemental powders followed by heat treatment (Schultz et al. 1987). Since the grain structure so obtained does not exhibit any crystallographic texture--and limits the energy product--special processing methods such as die-upsetting were needed to provide the crystallographic anisotropy (Schultz et al. 1989). While the coercivities of these nanocrystalline alloys are high, the remanent magnetization is decreased. Recent approaches to increase the magnetic induction have utilized exchange coupling in magnetically hard and soft phases. The Fe-rich compositions (e.g., Fe90NdTB3) result in a mixture of the hard Fez4Nd2B phase and soft ot-Fe phase. The nanoscale two-phase mixtures of a hard magnetic phase and a soft magnetic phase can exhibit values of remanent magnetization, Mr, significantly greater than the isotropic value of 0.5 Ms. This "remanence enhancement" is associated with exchange coupling between the hard and soft phases which forces the magnetization vector of the soft phase to be rotated to that of the hard phase (Smith et al. 1996). Two important requirements for alloys to exhibit remanence enhancement are a nanocrystalline grain size and a degree of coherence across interphase boundaries sufficient to enable adjacent phases to be exchange coupled. The significant feature of the exchange coupling is that it allows crystallographically isotropic materials to exhibit remanence values approaching those achieved after full alignment. Such two-phase nanoscale ferromagnetic alloys have been prepared by non-equilibrium methods such as melt spinning, mechanical alloying, and sputter deposition. Besides the highly reduced remanence the material cost is reduced by reduction in the content of the expensive hard rare-earth containing magnetic phase.

Nanostructured Materials

335

While the very low losses of the nanocrystalline soft magnetic materials (FINEMET or Nanoperm) are dependent on the nanometer grain size for their properties, the hard magnetic nanocrystalline alloys with remanence enhancement provide flexibility in processing, especially with powder materials. These remanence enhanced nanocrystalline hard magnetic alloys may find many applications as permanent magnet components. 12.7.3.4 Other ferromagnetic nanocrystalline materials The magnetocaloric effect is another magnetic property that can be enhanced in nanocrystalline microstructures. Shull and co-workers (1993b) have described new magnetic nanocomposite refrigerants that have 4 times the magnetocaloric effects of the best low temperature magnetic refrigerant. Small magnetic particles in a non-magnetic (or weakly magnetic) matrix exhibit alignment of the magnetic spins in a magnetic field. This increase in magnetic order lowers the magnetic entropy of the spin system. If this process is performed adiabatically the reduction in spin entropy is offset by an increase in lattice entropy and the specimen temperature will rise. This temperature rise A T is reversible--the specimen cools down on removal of the magnetic field--and is known as the magnetocaloric effect. The entropy change at a given (low) temperature for a system of magnetic spins is enhanced when the isolated spins are clustered. Shull et al. (1993a) have shown that the nanocomposite Gd3Gas-xFexOle gives superior magnetocaloric effects which increase with x up to x = 2.5 and can be extended to higher temperatures than existing paramagnetic refrigerants. 12.7.4 Chemical properties 12.7.4.1 Corrosion behavior The corrosion behavior of nanocrystalline materials has received very little attention. Thorpe et al. (1988) studied the corrosion behavior of nanocrystalline Fe32Ni36Cr14P12B6 alloy obtained by crystallization of the melt-spun amorphous ribbon, while Rofagha et al. (1991, 1993) studied nanocrystalline Ni-P alloys produced by the electrodeposition technique. The average dissolution rate of nanocrystalline Ni was found to be higher than that for coarse-grained material. However, while coarse-grained Ni suffers excessive intergranular corrosion in acidic media, nanocrystalline Ni exhibited more uniform corrosion morphology (Rofagha et al. 1991). 12.7.4.2 Catalytic properties A majority of the processing methods produce the nanocrystalline materials in the form of powder, and therefore the total surface area available can be accurately tailored by controlling the particle and grain sizes through the processing parameters. Beck and Siegel (1992) have shown that the chemical reactivity of nanocrystalline TiO2 which has been only lightly consolidated (to maximize the porosity and hence the surface area) is significantly higher than that in other commercially available TiO2 samples. It has been reported that the nanocrystalline sample has a much higher catalytic activity for S removal from H2S than any of the other samples both initially and even after exposure for 7h. The activity for nanocrystalline TiO2 is about 5 times greater than

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for the next most active materials (Fig. 12.7). Nanocrystalline samples not only have the increased activity, but the high activity is retained for a longer time than in commercial coarse-grained samples. This greatly enhanced activity was shown to result from a combination of unique and controllable features of the nanocrystalline materialmits high surface area combined with its rutile structure and its oxygen deficient composition. Annealing this material in an oxygen atmosphere decreased the number of anion vacancies and lowered the activity. There has been an increased activity in this area recently (Zaluski et al. 1996, Chen et al. 1997, Trudeau and Ying 1996). A new class of Pd-Fe alloys, showing continuous solid solubility at all compositions, was shown to have a high degree of H2 permeation through the membrane material, without experiencing poisoning by CO, HzS, and H20, which is

Nanostructured Materials

337

commonly encountered in membrane materials that are used for H2 separation. Hydrogen absorbing alloys (FeTi, LaNis, MgzNi, and other Mg-based materials) in the nanocrystalline form show much better hydrogen sorption properties than their polycrystalline counterparts. Further improvement was accomplished by modifying the metal powders with a small amount (<1 wt.%) of a catalyst, such as Pd, present in the form of small clusters or particles dispersed on the surface of the absorbing material (Zaluski et al. 1996). 12.8. A P P L I C A T I O N S m P R E S E N T

AND P O T E N T I A L

Nanocrystalline materials are relatively new and it is only during the past few years that researchers have started exploring the many potential benefits of these materials. A number of applications have been suggested based on the special attributes of these materials. Currently, bulk nanostructured soft magnetic iron-based alloys and WC-Co composites have found industrial uses (Rittner and Abraham 1998).

12.8.1 Structural applications Nanocrystalline materials have much higher hardness and strength than conventional coarse-grained materials. The enhanced diffusivity of these materials allows them to be deformed more easily, and at least in some cases, superplastically. The present and potential applications of nanocrystalline materials based on these properties are highlighted.

12.8.1.1 Cutting tools It has been shown that the hardness and strength of nanocrystalline materials is 4-5 times higher than that of coarse-grained materials of the same composition. McCandlish et al. (1994) have shown that the hardness of WC-Co composites with nanometered WC grains increases with decreasing grain size, reaching values as high as 2,200 VHN. Scratch tests suggested that the higher hardness in nanograined WC-Co is also accompanied by increased toughness. Thus, the abrasive wear resistance of nanostructured WC-Co is higher than the abrasive wear resistance of conventional WC-Co when the comparison is made at equal hardness (Fig. 12.8). Thus, nanocrystalline WC-Co cutting tools are expected to have more than double the lifetime than conventional coarse-grained composites. It has also been suggested that these materials could be used as fine drill bits for drilling holes in printed circuit boards and as rotary cutting tools (because of the fine grain size, sharp edges, and smooth surface finish). Ontario Hydro Technologies developed the Electrosleeve steam generator tube repair technology based on bulk nanostructured deposits (Cheung et al. 1996). This in-situ repair technology involves the application of an internal sleeve in the area of the tubesheet to effect a complete pressure boundary restoration of degraded tubing. The technology produces a nanostructured nickel microalloy sleeve (several hundred microns thick) having a continuous high strength metallurgical bond to the steam generator tube, superior strength and ductility, and resistance to wear, fatigue and corrosion. The Electrosleeve was found to significantly outperform other commercially available repair technologies.

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I 1800

'

I 2000

'

I 2200

' 2400

Hardness (kg/mm 2) Figure 12.8. Abrasion resistance versus hardness of conventional and nanocrystalline WC-Co materials (after McCandlish et al. 1994).

12.8.1.2 Nanocomposites Nanocomposites are those materials in which either the reinforcement or the matrix or both are on a nanometer scale. Further, nanocomposites may be such that the nanoreinforcement may be distributed either in the grain boundaries or inside the grains (Niihara 1991). A comprehensive review on ceramic nanocomposites has been published very recently (Bhaduri and Bhaduri 1998). It has been known that dispersing a second phase on a microscopic scale can considerably enhance the fracture toughness of ceramics. Since reducing the grain size to nanometer dimensions can provide increased strength and hardness, it is suggested that fabrication of micro/nanohybrids will lead to a class of supertough and superstrong ceramics. In these hybrids, a nanocomposite matrix is reinforced with sub micronsized particlesmwhiskers, platelets, and long fibers--and these hybrids show enhanced fracture toughness and strength up to very high temperatures, where the properties of coarse-grained composites normally degrade. Further, nanoreinforcements increase creep resistance by suppressing grain boundary sliding. Consequently, nanoreinforcements are now being used to boost the performance of a growing number of composites. Nanoparticles of SiC markedly improve the fracture toughness and hardness of MgO/SiC composites developed by Mitsubishi Mining & Cement Co., Ltd., Japan. Nanocrystalline fibers also are used to reinforce composites. Nextel 610, a polycrystalline A1203 fiber

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developed by 3M has a grain size of 100 nm and strength of 2.4 GPa (Das 1996). Microstructure of the mullite-based 3M fiber is a two-phase nanocomposite of 55% mullite and 45% A1203 by weight, with 20% 200-nm mullite needles randomly dispersed in a matrix of 100-nm A1203 particles. The fiber retains 85% of its tensile strength at 1200~ (Anon 1994). Such superceramics may find use in severe engineering applications such as high efficiency gas turbines, and in aerospace and automotive components. While high hardness and strength are usually associated with low ductility, it has been recently shown that a uniform dispersion of about 20 vol% of a nanocrystalline phase in an amorphous matrix increases the strength of the material significantly while maintaining an adequate level of ductility in the composite (Inoue et al. 1994). 12.8.1.3 Superplastic materials

It was mentioned earlier that ceramic materials such as TiO2 and CaF2 in the nanocrystalline state underwent plastic deformation at low temperatures (Karch et al. 1987); probably due to the porosity present in the samples. Whether such a phenomenon will be observed in other nanocrystalline materials is a matter of concern at the moment. However, it has been shown that materials in the nanocrystalline state show tendencies of superplastic deformation at temperatures lower and strain rates higher than conventional coarse-grained materials. Thus, the deformation behavior of SiC/SiC nanocomposites has been comparable with metals and Si3N4/SiC composites have been shown to deform superplastically (Niihara 1991). Since ceramics and conventionally brittle intermetallics can be rendered ductile at elevated temperatures (at least in some cases) by nanostructure processing, these materials can be formed to near-net shapes by means of deformation processing methods previously applicable only to produce ductile metal parts.

12.8.1.4 Coatings The high hardness and strength of nanocrystalline materials can be utilized to produce wear-resistant coatings. Thermal barrier coatings consist of a metallic bond coat, usually MCrA1Y, a thermally grown A1203 film, and a ceramic layer, usually yttria stabilized zirconia. These coatings reduce superalloy metal temperatures by as much as 300 F, and thereby provide durability to combustors, transition ducts, and turbine vanes. However, these coatings fail by spallation of the coating at or near the ceramic to metal bond line. With the gas turbine industry's demand for higher operating temperatures to provide enhanced fuel efficiency, thermal barrier coatings with reduced thermal conductivity provide the highest pay-off coating application. Because of the small grain size and increased phonon scattering at grain and layer boundaries, nanocrystalline coatings should adhere better and at the same time exhibit reduced thermal conductivity (Gell 1995). One can also deposit nanocrystalline coatings to improve the wear-resistance while retaining a tough interior for low-temperature applications. Nanocrystalline coatings deposited by laser plasma discharge increased the life of ZnS samples more than five times against abrasion/erosion/rain water corrosion/impact damage (Davanloo et al. 1993). A double layer coating has been found to be a suitable method to improve the mechanical properties and recording characteristics of amorphous-ultrafine thin films; controlling the polymer coating thickness is still a problem (Kunimoto et al. 1987).

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12.8.2 Magnetic applications In the area of permanent magnet materials, an exciting new development is that dealing with exchange coupled magnets. In these magnets there is an excess of ot-Fe (up to 40 vol.%) which couples to the hard magnetic phase, either NdzFe14B or SmzFe17N3. The phases have to be on the nanometer scale (typically _<30 nm) and the required microstructures are produced either by melt spinning or mechanical alloying followed by appropriate annealing treatments for devitrification. There is now a large market for strong rare-earth based permanent magnets in a wide variety of applications, and flexibility in manufacturing and lower material costs (less expensive rare-earth metals needed in remanence enhanced magnets) should encourage use of nanostructured materials. The exceptional soft magnetic properties of the "FINEMET" and "Nanoperm" nanocrystalline/amorphous Fe-base alloys promise applications in transformer cores for power frequencies, saturable reactors, high frequency transformers, and magnetic heads. Additional probable applications are for data communication interface components (pulse transformers), sensors (current transformers, magnetic direction sensors), common mode choke coils, and magnetic shielding. Magnetic recording materials to some extent already use nanostructured materials. Magnetic materials are used in both the information storage media and in the "write" and "read" heads. It is predicted that in the future essentially all media and heads will contain nanostructured materials (Berkowitz 1996). Thin film recording media (e.g., disks) typically have thicknesses of 10-100 nm and at very high bit densities, longitudinal media must be very thin (10-50 nm) while perpendicular recording media can be much thicker (50-250 nm) (Judy 1990). These films are mainly prepared by sputtering of Co-based alloys. In particulate media (e.g., tapes) the structure for high-density applications consists of an acicular iron core passivated by an oxide surface. The iron particles are about 200 nm long and about 20 nm in diameter (Sharrock 1990). For both particulate and film media the magnetic single domains must become even smaller to achieve higher densities with satisfactory signal-to-noise ratiosmbut not so small as to become paramagnetic.

12.8.3 Catalysts and hydrogen storage materials Finely-dispersed materials are traditionally used in catalysis because of the high specific surface area that is obtained when a sample is divided into small portions. In fact some of the first applications of nanoparticles were as catalysts (Ponec 1975). It is suggested that the most important synthetic materials discovered in the past 50 years are the metallic, highly dispersed heterogeneous catalysts such as Pt/A1203, NiSiO2, Pd/C, Rh/A1203 and many others. These catalysts are critical to the fuels industry, control of automobile emissions, and chemical synthesis. It has been shown in several studies that the catalytic activity and selectivity can be dependent on particle size. A strong correlation was observed between Pt particle sizes and catalytic activity. That is, very small Pt particles (3-6 nm) showed an unusual catalytic activity in the hydrogenation of pchloronitrobenzene. Commercial catalysts with about 20 nm Pt particles only promote the catalytic reduction of the NO2 group while the small (3-6 nm) Pt particles catalyze the unusual formation of dicyclohexylamine 8 in large quantity. However, the electronic

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structure of metal nanocrystallites supported on oxides suggest no change from bulk behavior even down to 1 nm for Pt, Ir, or Os. Furthermore it cannot be assured that the integrity of a nanocrystallitemits composition and positionmwill remain in tact when it is used for a catalyst. Metal hydrides offer the best compromise of safety, efficiency, and cost for storage of hydrogen for use as a fuel. Unfortunately their use is hindered by the sensitivity of their kinetic properties to surface oxidation. Hydrides of intermetallic compounds need activation treatments--often at elevated temperatures and pressures~after exposure to air. It was believed that a way to enhance their properties is to produce the materials in nanocrystalline form. The high density of grain boundaries/interfaces might increase diffusion. Nanocrystalline Fe40Ti60 and FesoTi50 exhibited, compared to coarse-grained alloys, increased solubility at low pressure, narrowing of the miscibility gap, lower plateau pressure, and disappearance of the transformation to y-FeTiH1.90 (Tessier et al. 1996). Nanocrystalline Mg2Ni inclusions in a two-phase Mg/Mg2Ni alloy showed improved hydrogen storage behavior in terms of the Mg2Ni catalyzing the decomposition of the molecular hydrogen, greatly increasing the hydrogen absorption/desorption kinetics (Wasz et al. 1995). A further advantage of the nanoscale microstructure is that the alloy powder does not comminute as the result of repeated hydrogen charging-discharging cycles. 12.8.4 F u n c t i o n a l n a n o s t r u c t u r e s m e l e c t r o n i c applications

"Functional" electronic and optical nanostructures are typically fashioned from singlecrystal semiconductor materials using a variety of possible pattern formation (e.g., electron beam lithography) and pattern transfer processes (e.g., reactive ion etching) which can create structures at the nanoscale. This is a large area of research and development of importance to the microelectronics industry for future directions such as quantum circuits and systems, massive memory, architecture and interconnected systems. Miniaturization of Si device size to the ultimate Si feature size of 3 to 6 nm is a goal. A discussion of the issues involved in this large sub area of nanotechnology is beyond the scope of this chapter. The reader is referred to reviews by Peyghambarian et al. (1996) and Dobisz et al. (1996). 12.9. C O N C L U D I N G

REMARKS

This chapter has described a subset of topics in the broad and rapidly growing research area of nanostructured materials. This review has concentrated on the synthesis, structure, and properties of what may be considered bulk, or the powder precursors to bulk materials. The wide variety of processing methods can lead to similar, but not identical nanocrystalline structures. The detailed structure of the grain boundaries in nanocrystalline materials remains controversial. Many of the unique properties reported for the early gas-condensed materials have been subsequently shown to result from sample porosity. These include low values of elastic moduli, ductility at low temperatures of nanostructured ceramics, increased coefficients of thermal expansion and specific heat, and large decreases in saturation magnetization. Some of the confirmed properties of nanostructured materials can be viewed as extrapolations of grain-size-sensitive behavior

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to the nanoscale, such as varistor breakdown voltage. Other properties, such as deformation behavior, appear to be the result of new phenomena that are not yet well understood. Applications of nanostructured materials are in their infancy, but the explosion of interest in this field is expected to provide an early maturation. REFERENCES

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Rofagha, R., Langer, R., E1-Sherik, A. M., Erb, U., Palumbo, G. and Aust, K. T. (1991) Scripta Metall. Mater., 25, 2867. Sanders, E G., Rittner, M., Kiedaisch, E., Weertman, J. R., Kung, H. and Lu, Y. C. (1997) NanoStructured Mater., 9, 433. Schultz, L., Wecker, J. and Hellstern, E. (1987) J. Appl. Phys., 61, 3583. Schultz, L., Schnitzke, K. and Wecker, J. (1989) J. Magn. Magn. Mater., 80, 115. Schumacher, S., Birringer, R., Strauss, R. and Gleiter, H. (1989) Acta Metall., 37, 2485. Sharrock, M. E (1990) MRS Bulletin, 15(3), 53. Shull, R. D., McMichael, R. D. and Ritter, J. J. (1993a) NanoStructured Mater., 2, 205. Shull, R. D., McMichael, R. D., Ritter, J. J. and Bennett, L. H. (1993b) MRS Symp. Proc., 286, 449. Siegel, R. W. (1990) MRS Bulletin, 15(10), 60. Siegel, R. W. (1991) in Processing of Metals and Alloys, vol. 15 of Materials Science & Technology - A Comprehensive Treatment, eds. Cahn, R. W. (VCH, Weinheim, Germany), p. 583. Siegel, R. W. (1992) in Materials Interfaces: Atomic Level Structure and Properties, eds. Wolf, D. and Yip, S. (Chapman & Hall, London), p. 431. Siegel, R. W. (1994) NanoStructured Mater., 4, 121. Siegel, R. W. (1997) Mater. Sci. Forum, 235-238, 851. Siegel, R. W. and Fougere, G. E. (1994) in Nanophase Materials: Synthesis, Properties, Applications, eds. Hadjipanayis, G. C. and Siegel, R. W. (Kluwer Acad. Publishers, Dordrecht, The Netherlands), p. 233. Siegel, R. W. and Thomas, G. J. (1992) Ultramicroscopy, 40, 376. Smith, E A. I., Ding, J., Street, R. and McCormick, E G. (1996) Scripta Mater., 34, 61. Sui, M. L. and Lu, K. (1994) Mater. Sci. Eng., A179-180, 541. Suryanarayana, C. (1995) Internat. Mater. Rev., 40, 41. Suryanarayana, C., Korth, G. E. and Froes, E H. (1997) Metall. Mater. Trans. A, 28A, 293. Suryanarayanan, R., Frey, C. A., Sastry, S. M. L., Waller, B. E., Bates, S. E. and Buhro, W. E. (1996) J. Mater. Res., 11,439. Terauchi, S., Koshizaki, N. and Umehara, H. (1995) NanoStructured Mater., 5, 71. Tessier, E, Zaluski, L., Zaluska, A., Str6m-Olsen, J. O. and Schulz, R. (1996) Mater. Sci. Forum, 225-227, 869. Thorpe, S. J., Ramaswami, B. and Aust, K. T. (1988) J. Electrochem. Soc., 135, 2162. Trudeau, M. L. and Ying, J. Y. (1996) Nanostructured Mater., 7, 245. Uda, M. (1992) NanoStructured Mater., 1, 101. Uyeda, R. (1991) Prog. Mater. Sci., 35, 1. Valiev, R. Z. (1994) in Nanophase Materials: Synthesis, Properties, Applications, eds. Hadjipanayis, G. C. and Siegel, R. W. (Kluwer Acad. Publishers, Dordrecht, The Netherlands), p. 275. Vandermeer, R. A. and Hu, H. (1994) Acta Metall. Mater., 42, 3071. Wasz, M. L., Schwarz, R. B., Srinivasan, S. and Kumar, M. P. S. (1995) Mater. Res. Soc. Symp. Proc., 393, 237. Weertman, J. R. and Averback, R. S. (1996) in Nanomaterials: Synthesis, Properties, and Applications, eds. Edlestein, A. S. and Cammarata, R. C. (Institute of Physics Publ., Bristol), p. 323. Weismfiller, J. (1996) in Synthesis and Processing of Nanocrystalline Powder, ed. Bourell, D. L. (TMS, Warrendale, PA), p. 3. Weismfiller, J., L6fler, J. and Kebler, M. (1995) NanoStructured Mater., 6, 105. Wong, L., Ostrander, D., Erb, U., Palumbo, G. and Aust, K. T. (1993) in Nanophases and Nanocrystalline Structures, eds. Shull, R. D. and Sanchez, J. M. (TMS, Warrendale, PA), p. 85. Yoshizawa, Y., Oguma, S. and Yamauchi, K. J. (1988) J. Appl. Phys., 64, 6044. Zaluski, L., Zaluska, A., Tessier, P., Str6m-Olsen, J. O. and Schulz, R. (1996) Mater. Sci. Forum, 225-22"7, 853.

Chapter 13 Powder Consolidation 13.1. Introduction 13.2. Metastability in Powder Consolidation 13.3. Consolidation of Metastable Powders 13.3.1 Sintering Mechanisms 13.3.2 Scaling Laws 13.3.3 Powder Contamination 13.4. Consolidation Methods 13.4.1 Conventional Sintering 13.4.1.1 Cold Compaction 13.4.1.2 Warm Compaction 13.4.1.3 Conventional Sintering of Nanopowders 13.4.1.4 Nano-Composite Densification 13.4.1.5 Grain-Size Control 13.4.2 Pressure Consolidation of Metastable Powders 13.4.3 Pressure-Assisted Consolidation Methods 13.4.4 Non-Conventional Sintering Methods 13.4.4.1 Microwave Sintering 13.4.4.2 Field-Assisted Sintering 13.4.4.3 Shockwave Consolidation 13.5. Concluding Remarks Acknowledgements References

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Chapter 13 Powder Consolidation JOANNA R. GROZA

13.1. I N T R O D U C T I O N

The unique properties of metastable materials may be fully exploited if powders are densifted into bulk parts that retain the initial metastable features. Numerous applications, particularly structural materials, require fully dense parts of certain size, geometry and final properties. In addition, dense specimens with metastable microstructures are important for magnetic, electrical or electronic functions such as permanent magnets, substrates in electronic circuitry, magnetic recording heads, capacitors, electrodes, and varistors. The key characteristic of metastable powder consolidation processes is to achieve densification with minimal microstructural coarsening and/or undesirable microstructural transformations. This requirement places significant restrictions on the consolidation process, particularly high temperature exposure which must be controlled more carefully than in the case of regular powders. Although lowest temperatures are desirable, full interparticle bonding and pore free structures must be achieved. These restrictions to the sintering process are dependent on the type of microstructural metastability that needs to be retained. For instance, retention of chemical homogeneity is less restrictive on the sintering process than retention of metastable phases. Consolidation of far-from equilibrium powder materials has been addressed quite extensively with the advent of rapid solidification (RS) techniques that have usually produced metastable materials as powders or flakes with sizes in the micron range. The main focus in the consolidation of RS powders has been to retain the metastable condition of the initial microstructures. The recently developed nanopowders have added a new step in metastability and its retention requirements. In this case, the particle size decreased dramatically, often down to nanometer ranges. While consolidation of micron size RS powders has been accommodated into the usual sintering patterns, consolidation of nanosize powders has presented significant additional challenges. An increased tendency to return to equilibrium results in inevitable loss of nanofeatures (e.g., by grain coarsening). This coarsening tendency and small specimen size produced by nanopowder consolidation severely limited the assessment of nanomaterial properties, particularly mechanical (Sanders et al. 1997a). In fact, numerous controversies on nanomaterial properties have originated from artifacts in nanopowder processing. 347

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Table 13.1. Types of metastabilities in far-from-equilibrium powder materials (adapted from Turnbull 1981). Nature of meta (or

Typical excess energy (R Tm)

"in")-stability

Material examples

Compositional Structural

Supersaturated/extended solid solution P' F Altemate crystal structure P, F, B Amorphous solid P, F, B

<0.5

Morphological

Compositional modulations F, P

<0.1

Interphase dispersions F, P

<0.1

Nanocrystalline materials F, P

<0.4

<1 <0.5

P = powders, F = film, B = bulk.

13.2. M E T A S T A B I L I T Y IN P O W D E R C O N S O L I D A T I O N

As shown in Table 13.1, metastable structures may be divided into three main categories (Turnbull 1981). Most of these metastable conditions have been associated with powder materials or, conversely, the powder materials display the whole range of metastabilities. Particularly, powder materials processed by rapid solidification, high-energy milling, precursor pyrolysis, and vapor condensation may display far-from-equilibrium structures. In many cases, a combination of such metastabilities may be found. For instance, powders in the nanosize range display inherent morphological metastability associated with topological (e.g., alternate crystal structures than at equilibrium) or compositional (extended solubility or amorphous phases) metastabilities. Topological metastability in nanopowders is the result of both undercooling and increase in pressure due to the extra surface tension in particles with large curvatures (Gibbs-Thompson effect). In materials with high-pressure polymorphism, synthesis of particles in the nanometer range often results in the formation of the high pressure phases such as cubic BaTiO3, monoclinic Y203 and tetragonal ZrO2 (Uchino et al. 1989, Skandan et al. 1992, Garvie 1978). Alternate crystal structures due to interfacial energy effects at large undercoolings were shown by Krauss and Birringer (1997) in nanocrystalline refractory metals (the A15 modification of fl-W and fl-Ta). It is noteworthy that such metastable phases are obtained only in powder form. When films of W and Ta have been synthesized with same grain size, only the equilibrium structures have formed. In contrast to high-pressure ceramic phases, the latter metal structures have a lower density than the equilibrium crystals. The consolidation process or densification or sintering is aimed at transforming the initial powders into bulk materials with a minimum amount of pores, or no pores. The driving force for densification is the tendency to reduce the free surface energy associated with individual powder particles. This driving force increases with the reduction of the initial particle size and becomes significant for powders in the nanometer range. Although there is a large driving force, sintering usually involves exposures at temperatures above room temperature for a certain time. These conditions may diminish or sometimes revert the initial metastable state of the powders. Since at least some departure from metastability will occur during sintering, it is important to define this metastability and the conditions

349

Powder Consolidation Table 13.2. Physical and structural metastability ranges.

Material

Critical grain size (nm)

BaTiO3

120

TiO2 Y2 03 ZrO2 Cr203

120 ~50 ~ 13 8-26 >80

Nano property

Equilibrium property

Cubic

Tetragonal

Uchino et al. (1989)

Drop in Tcurie* Anatase Monoclinic Tetragonal Superparamagnetic

Constant Tcurie Rutile Cubic Monoclinic Antiferromagnetic

Takeuchi et al. (1997) Hague and Mayo (1993) Skandan et al. (1978) Garvie (1978) Balachandran et al. (1995)

Reference

As the particle size decreases, the Curie temperature also decreases.

upon which it is lost. As already shown, compositional metastability (extended solubility) is the most compliant to subsequent sintering conditions or the consolidation conditions for these powders are the least restrictive (Flinn 1985). Consolidation parameters must be more controlled when initial powders display morphological and/or topological metastabilities. Morphological metastability is not a concern in RS powders in which grain sizes are on the order of microns. In contrast, for nanocrystalline powders with initial particle/grain size on the order of nanometers, morphological instability becomes critical. So far, the consolidation efforts in nanosize powders have been generally aimed at retaining this particular metastability (i.e., finest grain size), without regard for the other metastable conditions (e.g., compositional or structural). Topological metastability presents the most challenges for the consolidation step. This metastability is easily defined since qualitative sharp changes such as phase transitions occur. Some phase transitions from a metastable to an equilibrium phase and the critical grain size range at which they occur are listed in Table 13.2. It is noticeable that grain sizes associated with novel crystal structures are, generally, extremely small (e.g., less than 26 nm for Y203 and ZrO2). In other cases, the transformation path to equilibrium may be quite complex going through a range of metastable conditions. Table 13.2 also presents cases of physical property modifications upon recovery of the equilibrium property at a critical grain size. Densification commonly results in grain coarsening. With the exception of the above few cases shown in Table 13.2, a definite threshold value for the grain size to retain the new and unique metastable features has proven difficult to assess. For instance, there is not yet consensus on the strengthening mechanisms or superplasticity benefits in the nanoregime, but a continuous effort is to push the limit toward smallest grain sizes. In this context, an arbitrary limit for the maximum grain size in the nanomaterial consolidation will be considered < 100 nm throughout this chapter.

13.3. C O N S O L I D A T I O N OF M E T A S T A B L E P O W D E R S

The densification process for conventional powders is well known both theoretically and practically. The main steps are collapse of pores, plastic flow and diffusion to fill the void

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space and provide good interparticle bonding. However, the densification of metastable powders is more demanding due to restrictions related to grain growth and equilibrium phase development. Nanopowders present additional challenges. Powder agglomeration, high reactivity and inherent contamination, grain coarsening and ultimate loss of the nanofeatures, inability to fabricate large and dense parts are among the main problems. The most recent efforts have been very fruitful in overcoming some of these problems (e.g., agglomeration and grain size control). This has been accomplished by major improvements in the nanopowder synthesis methods and understanding of the densification process such as, for example, pore effects in nanosintering. For nanopowder consolidation, the critical issue is whether sintering of nanocrystalline materials involves the same mechanisms as conventional ones or some qualitatively new phenomena come into play. This chapter is written with this question in mind. More specifically, differences between the sintering of regular and nanosize powders will be highlighted. Alternatively, the possibility of new sintering mechanisms will be examined. Consolidation of metastable powders has been discussed in a number of reviews that address either RS or nanopowder densification (Flinn 1985, Grant 1985, Suryanarayana 1991, Mayo 1996, Weissmuller 1996, Lu 1996, Groza 1998). This chapter will first cover some thermodynamic and kinetic aspects of metastable powder densification, sintering mechanisms, activation energies and scaling laws. Due to the critical effect of surface contamination on small particles, the impurity role in sintering is separately dealt with. Next, conventional sintering, along with cold compaction, pore size and distribution and their effects on sintering and grain coarsening are addressed. Finally, methods for full densification of metastable powders and their ability to maintain the metastable features are presented. T h e r m o d y n a m i c a l l y , powders are unstable due to large surface area. This morphological metastability is highly dependent on particle size, with nanopowders being the most unstable. The surface area increases dramatically with the reduction in powder diameter. The specific surface area (per unit mass), S, is given as: S = 6/PthD

(13.1)

where Pth is the material theoretical density and D is the particle diameter. For comparison, the surface area of 5 #m copper powder is 0.134 m2/g while for 5 nm powder is 134 m2/g. The sintering process is driven by the tendency to reduce this surface area and therefore, fine particles will have an enhanced tendency for sintering. The driving force for sintering is proportional to ~,/r, where y is the surface energy, generally considered isotropic, and r the particle size. In the case of nanosize powders with a limited number of atoms, departures from an isotropic behavior (i.e., anisotropy) have been noticed. Direct TEM studies of y-A1203 (Bonevics and Marks 1992), ZrO2 (Rankin and Sheldon 1995, Carim 1994), and CeO2 (Guillou et al. 1997) showed that nanoparticles have a faceted appearance with anisotropic surface energies. Next, a different local atomic arrangement at the surface of a nanocrystal may result in a different surface energy value than in conventional powders. This may be the case of amorphous layers or non-stoichiometric oxides which have been observed on nanoparticle surfaces (Luo et al. 1995, Trudeau et al. 1995). As Cahn (1992) infers, the different nature of the surface oxide on nanopowders may be

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351

Table 13.3. Sintering onset temperatures for powders in the nanoregime.

Material Fe W ZrO2 TiO2 TiN Y203

Particle size (nm)

Sintering temperature (K)

40 2 (mm) 31 9 8-9 70 40 12-14 "~10 4

370 (0.21Tm) ~900 (0.5Tm) 1100 (0.3Tm) 900 (0.24Tin) 870-920 ('~0.3 Tin) 1370 (~0.5Tm) <950 (0.46Tm) 823 (0.4Tm) 1170* (0.37 Tm) ~600 (0.25Tm)

Reference Dominguez and Bitog (1995) Dominguez and Bigot (1995) Alymov et al. (1992) Alymov et al. (1992) Hahn (1993), Skandan (1995) Hahn (1993), Skandan (1995) Yan and Rhodes (1983) Hahn et al. (1990) Rabe and Wasche (1995) Hahn (1993)

* Sintering is completed at 1823 K. Nanocrystalline TiN is approximately 100% dense at 1823 K. Microcrystalline TiN is only 63% dense at 1823 K.

an indication of a different surface structure. Macroscopically, stabilization of particle surfaces, i.e., lower surface energy values when oxides are formed, results in a decrease of the driving force for sintering. This has been the case for any powder particles but becomes more critical for nanopowders that typically contain large amounts of adsorbates (Rabe and Wasche 1995, Andrievski 1994, Ying 1993). Kinetically, sintering of nanopowders is significantly enhanced. The dependence of the sintering kinetics on grain size may be illustrated using the equation for the densification rate (dL/L dt) developed by Johnson and co-workers (Hansen et al. 1992) for all stages of sintering:

dL L dt

FS2 (~DbFb

kT

--------Z-d +

DvFv) d3

(13.3)

where g is the surface energy, f2 is the atomic volume, 3 is the grain boundary width, Db and Dv are the grain boundary and bulk diffusivities, 1-'b and Fv are functions of density, kT has the usual meaning, and d is the grain size. From this equation, it is seen that a decrease in grain size by three orders of magnitude (e.g., from # m to nm) could enhance sintering rates by up to 12 orders of magnitude. Consequently, sintering of nanopowders may be accomplished at significantly lower temperatures and shorter durations than for conventional powders. This has been noticed for numerous real nanoparticle sintering. A few examples of such depressed sintering onset temperatures are given in Table 13.3. A more dramatic effect of the sintering temperature depression in nanopowders is seen in the liquid phase sintering systems, such as WC-Co. Due to the accelerated sintering at lower temperatures in the nanoregime, full densification may take place entirely during solid state sintering (Porat et al. 1996).

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13.3.1 S i n t e r i n g mechanisms

The densification process consists of solid particle bonding or neck formation followed by continuous closing of pores from a largely open porosity to essentially a pore-free body. For simplicity, solid state densification is considered to be accomplished through three stages: initial, intermediate and final. Multiple mechanisms are involved throughout these stages, viz., evaporation-condensation, surface diffusion, grain boundary diffusion, bulk diffusion, viscous flow and plastic deformation. Each transport process exhibits a particular dependence on the particle/grain size and defect density. The highest sensitivity on particle size is that of surface and grain boundary diffusion. Although simultaneous mechanisms operate, the common sintering models attribute a predominant mass transport path to a specific sintering stage. For instance, surface diffusion is considered the principal mechanism during the initial stage when the main event is the neck formation. So far, a similar partition of the sintering process is accepted in the existing sintering theories applied to nanopowders (Mayo 1996). The largest departure from conventional powder sintering is seen in the initial sintering stages. For nanoparticles with excessive surface area and highly curved surfaces, surface diffusion is expected to be rapid in early sintering stages. Evidence of the surface diffusion contribution to the neck formation in nanopowder sintering is given by many researchers such as in zirconia (Hahn et al. 1990, Thuenissen et al. 1993, Rankin and Sheldon 1995), alumina (Bonevics and Marks 1992) and TiO2 (Hahn et al. 1990). In these early sintering stages, lower activation energies than for conventional diffusion have been reported in 15 nm Y-TZP (Thuenissen et al. 1993), 13.5 nm A1203, 11.5 nm TiO2 (Vergnon et al. 1974) and 40 nm W (Trusov et al. 1989). For instance, the activation energy (134 kJ/mol) for sintering 40 nm W is lower than the value for volume diffusion (580 kJ/mol) and surface diffusion (~300 kJ/mol). Some of these low activation energy values cannot unambiguously be assigned to surface diffusion due to the lack of comparable surface diffusion literature data, particularly for ceramics. In addition, a more valid comparison should use nanomaterial diffusion data rather than diffusion data on regular materials. For conventional powders, it is known that surface diffusion does not lead to densification but to grain coarsening. As already shown, surface diffusion is the most sensitive to particle size. Therefore, enhanced surface diffusion with reduced low-temperature densification should be observed in nanoparticle sintering. Abnormal grain growth and pore coarsening without densification were noticed at low sintering temperatures in ultrafine silicon (Hayashi and Kihara 1987). However, this is more of an exception, since for nanoparticles, enhanced low-temperature densification has been typically observed (Table 13.2). Molecular dynamics simulations also indicated extremely fast sintering of nanoparticles (Zhu and Averback 1996, Averback et al. 1996). Surface diffusion cannot explain this behavior. Furthermore, the usual approach to use high heating rates to avoid deleterious surface diffusion does not seem to always apply to small powders. In some cases, slow heating rates induced surprisingly high densities (Yoo et al. 1987). In other cases, no effect of the heating rates on densification or grain growth has been observed (Chen and Mayo 1993). At this point, the contribution of surface diffusion to nanosintering is not clear. The ratio between bulk (and/or grain boundary) and surface diffusion contribution to overall densification is not known and may be different

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353

as compared to conventional materials. Either a major shift in the contribution of various diffusion mechanisms or a change in the sintering mechanisms is possible, at least in the initial densification stage. Some other mechanisms for neck formation in nanopowders have been suggested: grain boundary slip (Thuenissen et al. 1993), dislocation motion, grain rotation, viscous flow and even grain boundary melting (Trusov et al. 1989, Zhu and Averback 1995). Dislocation motion, viscous flow and grain rotation mechanisms have been identified in molecular dynamics (MD) studies that simulate nanoparticle sintering (Zhu and Averback 1995, 1996, Averback et al. 1996). Sintering of a copper nanoparticle pair takes place in picoseconds. Calculations for the observed shrinkage based on surface or grain-boundary diffusion cannot explain this rapid sintering behavior (Averback et al. 1996). To account for the shrinkage, the grain boundary diffusivity in Cu at 700 K has to be 10 -2 cm2/s, a value that is two orders of magnitude larger than liquid diffusivity in Cu. This rapid shrinkage simultaneous with neck formation is attributed to fast dislocation activity driven by the contact Hertzian stresses that exceed the ideal shear strength. After the neck forms by dislocation slip, the adjacent particles rotate to achieve a minimum grain boundary energy (Zhu and Averback 1996). TEM studies of zirconia, ceria and copper nanoparticles indicate that such a rotation process occurs during early sintering (Rankin and Sheldon 1995, Guillou et al. 1997 and Yeadon et al. 1997). For late sintering stages, activation energies are in better agreement with values determined for sintering of conventional powders (Vergnon et al. 1974, Messing and Kumagai 1994, Shaik and Milligan 1997). For instance, Messing and Kumagai (1994) found activation energies that correspond to volume diffusion for sintering ot-A1203 at 90% density. By late sintering stages, the initial nanopowders have already lost some of the more conspicuous initial nanofeatures and therefore, their densification behavior converges towards that of regular grain-sized materials. Only slight changes, such as the predominance of grain boundary diffusion instead of volume diffusion, have been noticed in the densification by power law creep of Fe-10 wt.% Cu during sinter forging (Shaik and Milligan 1997). This shift to grain boundary diffusion is not unreasonable since sintering temperatures are lower than in large grain powders.

13.3.2 Scaling laws As shown in Table 13.2, the sintering onset for nanoparticles occurs at temperatures lower (0.2-0.3Tm) than for conventional powders (0.5-0.8Tin). Full densification of nanopowders is completed at temperatures lower than that for conventional powders, as well. Nanosize Si3N4 was fully sintered at 700 K lower than the coarse grained analogs (Pechenik et al. 1993). TiC powder of 140-170 nm particle size sintered to 91% density at 1900 K as compared to 5 # m powder that achieved the same density at 3070 K (Panfilov et al. 1985). Nanometer size TiO2 (12-14 nm) completes its sintering at "-~1300 K as compared to more than 1670 K in 1.3/zm TiO2 (Hahn et al. 1990). To rationalize this decrease in the sintering temperature, different scaling laws have been applied. For particles under 100 nm, Alymov et al. (1992) developed an empirical relationship for the dependence of the sintering onset temperature, Ts, on the mean particle size, d:

354

J. R. G r o z a

In Ts = 1 / d

(13.3)

They calculated a ratio of sintering onset to melting temperature which is in good agreement with experimental data for sintering of nanometals. Herring (1950) developed the classical scaling rule for the effect of the particle size on sintering time. The time, t, to achieve the same sintering condition is related to the powder particle size, d, if the same sintering mechanisms operate, by the equation: tit2 = ( d l / d 2 ) n

(13.4)

where n is the exponent which has a value of 3 if the main sintering mechanism is volume diffusion and a value of 4 if the main sintering mechanism is grain boundary diffusion. Considering an Arrhenius expression for temperature, the sintering temperature dependence on the particle size may be calculated. Reasonable agreement of experimental and calculated temperature data was found in TiO2 and A1203 assuming certain diffusion mechanisms (Yan and Rhodes 1983, Messing and Kumagai 1994). For instance, in TiO2, the fitting value for the activation energy (139 kJ/mol) used is that of grain boundary diffusion (i.e., close to half of the lattice diffusion of 251 kJ/mol) (Yan and Rhodes 1983). When applying Herring's law to Andrievski's (1994) data on sintering pure nickel particles of different sizes to 60% density, the calculated activation energy value (239 kJ/mol) was closer to volume diffusion (284 kJ/mol) than to grain boundary diffusion (115 kJ/mol) in nickel. However, volume diffusion is less likely to be the principal sintering mechanism since temperatures are low and particles are fine. This points out to uncertainties in using the scaling law to determine the activation energy values and thus, deduce the sintering mechanisms. Conversely, the predictions of sintering temperatures as a function of particle size may not be accurate, giving primarily qualitative estimates. The main concern in using Herring's scaling law is related to a possible change of sintering mechanisms from conventional size to nanopowders and with sintering temperature. Finally, as in any extrapolation, a careful inspection of the initial assumptions is in order. Herring cautioned that the degree of contamination on the surface should be kept the same for all particle sizes. This is definitely questionable when conventional and nanosize particles are compared. Not only higher contamination levels have been reported (Andrievski 1994), but also different type of oxides have been identified on nanoparticles (Luo et al. 1995, Trudeau et al. 1995). In addition, when sintering occurs at lower temperatures, the complete reduction of oxides or hydroxides may be hindered. The specific effect of the surface contamination on sintering will be discussed in the next section. Technologically, there are significant benefits from the lower sintering temperatures of nanopowders, such as, avoiding sintering aids, decomposition of phases, deleterious interfacial interactions, and undesirable phase transformations (e.g., the monoclinic to tetragonal ZrO2 transformation that inevitably results in cracking (Hahn 1993)).

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13.3.3 Powder c o n t a m i n a t i o n

As already mentioned, surface contamination is inevitable for any powder material and may become significant in nanopowders due to their large surface area. Metastable powders obtained by mechanical alloying or milling are a special case. Their structure is either amorphous or nanocrystalline but the powder particle size is in the micron range. In this case, the surface area is less than for nanometer powder particles but contamination is due to a large impurity intake during the milling process. Oxides, nitrides and other compounds are often found in consolidated parts made of attrition-milled nanopowders. Although products of contamination, these compounds maybe used to advantage when present as fine dispersions to prevent grain coarsening upon further densification (Perez et al. 1996). Presently, distinct efforts have aimed at minimizing contamination in attrition milled powders (Goodwin and Ward-Close 1993 and Koch 1997). Large amounts of adsorbates have been documented on both ceramic and metal nanopowders (Ying 1993, Andrievski 1994). The influence of contamination on sintering was studied by in-situ TEM under controlled oxygen levels as compared to ultrahigh vacuum conditions for both metal and ceramic nanopowders (Bonevics and Marks 1992, Rankin and Sheldon 1995 and Olynick et al. 1995). All results indicate that ultra-clean nanoparticles sinter very rapidly even at room temperature. In contrast, when sintering of nanoparticles takes place in the presence of controlled oxygen traces, very little or no neck growth was observed (Olynick et al. 1996). This oxygen contamination decreases the surface energy and slows down the sintering kinetics. More specifically, TEM studies showed that sintering of Fe-Ni nanoparticles starts only when oxide layers are reduced (Eldrup et al. 1993). This occurred above 500 K in hydrogen. When vacuum was used, no sintering took place up to 725 K. As already mentioned, reducing the amount of oxygen in TiN nanopowders decreases the sintering temperature (Rabe and Wasche 1995). The sintering environment appears to have a similar effect in nanopowders as in conventional ones. Dominguez and Bigot (1995) calculated a lower activation energy for sintering nano-iron in hydrogen as compared to vacuum. Consistently, ceramics sintered bettermhigher densities at same temperatures or similar densities at lower temperatures--in vacuum than in air (Hahn 1993, Skandan 1995). To avoid further contamination, the as-produced nanopowders may be either consolidated in-situ or handled in controlled environment prior to sintering. This adds large costs to the actual component production. In a manufacturing environment, powders are consolidated in a remote place from the powder producing area. Therefore, contamination of small particles is likely to occur. The role of oxygen in nano-iron sintering was addressed by Bourell and Kaysser (1994). The powders were produced by an evaporation-condensation method. Although sintering was carried out in a reducing hydrogen atmosphere, oxygen was found as fine oxides in the densified product. Alternatively, consolidation processes that are less susceptible to the contamination levels are recommended (e.g., high pressure consolidation and severe plastic deformation), although not always practical. Some electrical-field-assisted sintering processes that claim an initial oxide reduction may be also helpful (Groza 1998).

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13.4. C O N S O L I D A T I O N M E T H O D S

The selection of the consolidation method for far-from equilibrium powders depends on the acceptable final degree of metastability, initial powder particle size and conditions and, certainly on the available equipment. Rapidly solidified powders are generally in the micron range and therefore, the consolidation method is less dependent on the initial powder handling procedures. The restrictions arise from the degree of metastability that should be retained in the final product. Therefore, consolidation methods that trade pressure for high temperatures or long exposures have been traditionally applied (Flinn 1985, Grant 1985, Suryanarayana 1991). When the powder particle size decreases below the micron range, consolidation efforts are faced with additional problems related to agglomeration, interparticle friction and contamination. In spite of these new challenges, both conventional (pressureless) and pressure-assisted sintering methods have been applied for nanopowder consolidation.

13.4.1 Conventional sintering Conventional sintering has been applied for full densification of both ceramic and metal nanopowders. The initial step in most densification processes is to compact the powders at room temperature, or cold compaction, to form a green body. Final sintering results are largely dictated by the green compact microstructure (German 1996). A larger number of initial point contacts, smaller pores in a high green density compact and a uniform pore distribution favor a higher final density. Sintering time may be shorter and lower sintering temperatures may be used. Conversely, many sintering defects may be traced to the green compact structure. Inhomogeneities in density, packing, and grain/particle size in the green compact will limit the final sintered density. Generally, nanocrystalline powders are more sensitive to green compact defects as compared to conventional powders. The most common problem is the elimination of large pores that originate from the green compact. This elimination requires high temperatures upon subsequent sintering thereby promoting unwanted grain growth and loosing the desired nanosize features. 13.4.1.1 Cold compaction This comprises specific stages that involve sliding and rearrangement of particles, elastic compression at particle contact points, plastic yielding for metals or fragmentation for brittle materials. Similar steps are seen in nanopowder compaction. The principal differences in nanoparticles arise from the specifics of small particle sliding and friction, as well as from the changes in dislocation assisted deformation behavior. Sliding and rearrangement in nanopowders are severely restricted owing to large frictional forces among powder particles. Mechanical friction resistance is substantial due to numerous interparticle contact points. Irregular grain boundaries favor agglomerate formation. As a result, particle rearrangement is hindered and lower green densities are likely to be achieved as compared to conventional micron size powders. For example, Bourell et al. (1993) achieved 45% density for 70 nm Y-TZP and only 29% for 40 nm. Conversely, when particle sliding is facilitated, high green values are obtained. Examples of green densities obtained in cold compaction of nanopowders are shown in Table 13.4. For comparison,

Powder Consolidation

33'/

Table 13.4. Cold compaction of nanopowders.

Powder

Particle size (nm)

Compaction method

Pressure (GPa)

Green density (%)

Reference

v-A1203

20

Cryogenic UP

2

72

Wu and De Jonghe (1996)

v-A1203

20

Dry UP

2

64

Chen et al. (1994)

v-A120 3 SiO2

13 15

HP HP

5.6 5.6

90 80

Gallas et al. (1997) Gallas et al. (1997)

ZrO2 Si3N 4 Si3N 4

6 50 50

UP UP CIP

1 1 1

55 54 60

Skandan (1995) Andrievski (1994) Andrievski (1994)

Si3N 4

15

UP

1

47

Andrievski (1994)

CIP UP UP

1 1 1

50 90 49

Andrievski (1994) Bonetti et al. (1993) Andrievski (1994)

Si3N 4 Fe3A1 Ni

15 10-15" 15

Ni

50

UP

1

58

Andrievski (1994)

Ni

5000

UP

1

73

Andrievski (1994)

Ni

15

CIP

1

61

Andrievski (1994)

Ni Ni Fe

50 5000 26

CIP CIP UP

1 1 2

68 77 68

Andrievski (1994) Andrievski (1994) Dominguez et al. (1995)

UP - uniaxial pressing (compaction), CIP - cold isostatic pressing, HP -- high pressure compaction. For the mechanically alloyed specimens only the grain size was reported (the particle size was not reported but, is assumed to be in the order of microns).

some results for micron size powders are also shown. For cold compaction of nanopowders, the use of lubricants, such as in wet processing or using liquid nitrogen, and particle coatings results in higher density (Chen et al. 1994, Mayo 1996). A better initial particle rearrangement facilitated by good wetting enhances sinterability even without increasing green density. For instance, Yan and Rhodes (1993) achieved 99% density by pressureless sintering of only 41% dense green compact of TiO2 nanopowders prepared by the centrifugation method. Similar full densification with 80 nm final grain size after wet compaction is reported by Mayo (1996) in ZRO2-3 mol%. Y203. These results are to be contrasted to dry compaction of the same powders that requires 0.5-1 GPa pressure to achieve similar densification characteristics. Therefore, techniques that take advantage of wet processing of ceramics such as centrifugation, tape and slip casting, and osmotic and pressure filtration are given preference. More results on cold compaction of nanoceramics are discussed in a recent review by Mayo (1996). In dry compaction, ceramic powders are more difficult to compact than metals and therefore, very large pressure levels are applied (Table 13.4). Experimentally, the literature reports indicate room temperature densities reaching >95 % for nanometals but only 7590% for nanoceramics (Gutmanas 1990, Gutmanas et al. 1994, Gallas et al. 1997). For instance, Gutmanas et al. (1994) obtained high densities up to 97% at room temperature and 99% at 575 K in Ni (65 nm) and Fe (30 nm) at 3 GPa. In cold isostatic pressing at very high pressures (up to 5.6 GPa), densities in excess of 90% for nanocrystalline A1203

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J. R. Groza

and 80% for SiO2 have been achieved by Gallas et al. (1997). A major inconvenience in applying large pressures, other than equipment limitation, is the high level of residual stresses in ceramic powders that may result in fracture upon subsequent handling (Mayo 1996). A high value of the green density may not always reflect a uniform pore size and distribution. When large agglomerates are present, the green density may be larger than in non- agglomerated powders. For instance, a high density of 75% in as-compacted nanoTiO2 at 425 K under 2 GPa pressure for 2 hours was explained by easier packing of large agglomerates (50 nm) rather than individual smaller particles (14 nm) (Hahn et al. 1990). For practical purposes, transparency of green ceramics may be used as a quick check of the compaction results (Mayo 1996, Gallas et al. 1997). A translucent green compact is an indication that pores are smaller than the visible light wavelength and consequently may reach high densities upon further densification. A particular case is represented by mechanically alloyed powders that most commonly consist of micron size particles with multigrains in the nanometer range. In this case, the particle size dictates the cold compaction behavior and high green densities are, therefore, achieved. For instance, a green density of ~90% has been reached under 1 GPa in mechanically alloyed Fe3A1 powders (Bonetti et al. 1993) (see Table 13.4). As already shown, particle surface condition becomes increasingly important in nanopowders due to the dramatic increase in the surface area. Literature examples illustrate the interdependence of green density value on surface chemistries, adsorption characteristics and reactivity. Andrievski (1994) measured high absorbed gas levels on 50 nm Ni powders and, consequently, attained low final green densities even at high pressures. As seen in Table 13.4, the green density values reached in these powders at 1 GPa pressure by uniaxial pressing are only ~58% and 68% by cold isostatic pressing (CIP). Even lower green densities were obtained for smaller Ni powders (15 nm in Table 13.4). In contrast, adsorbates in ceramics are at significant lower levels. As a result, most ceramics are by far less sensitive to storage in humid conditions although the level of adsorbed gases such as CO, CO2, N2 and H2 may be quite high (Andrievski 1994, Mayo 1996). 13.4.1.2 Warm c o m p a c t i o n

To eliminate adsorbates and enhance interparticle bonding, warm compaction at temperatures up to 675 K has been largely applied to nanopowders, particularly for those processed by inert gas condensation (Fougere et al. 1995, Sanders et al. 1997a,b). These powders were densified under high-pressure application (1.4 GPa) for 5-900 min at temperatures between 375 and 575 K. Densities obtained exceeded 95%. Higher densities were obtained when volatile contaminants were carefully eliminated based on studies of gas desorption temperatures and kinetics (Sanders et al. 1997b). The main purpose of such studies was to minimize the consolidation artifacts, which hindered the accurate characterization of mechanical properties. Higher compaction temperatures and, consequently, slightly larger grain sizes may also allow plastic deformation to contribute to nanometal densification (Gutmanas 1990, Sanders et al. 1997b). In late sintering stages, the sinterability of nanosize powders has been greatly enhanced by improvements in nanopowder synthesis, particularly the agglomeration control (Messing and Kumagai 1994, Boutz et al. 1994, Mayo 1996). Similar to conventional

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powders, full density or rapid sintering of nanosize powders is achieved when the green structure contains a narrow pore size distribution. Conversely, densification is retarded or inhibited when pore distribution is wide. In this case, big pores become larger and only smaller pores shrink. The removal of large pores is a lengthy process and requires higher temperatures. This way, the overall effect of large pores is to slow densification and induce undesirable coarsening. For nanoparticles, this coarsening is critical since it may compromise the overall notion of nanograined materials. Large pores usually originate from agglomerated powders and nanoparticles are inevitably linked to a high agglomeration tendency. Generally, the smaller the particles, the larger the agglomerates. Agglomerated powders have a bimodal pore distribution with small inter-agglomerate and large intra-agglomerates pores. The removal of large interagglomerates pores based on vacancy diffusion requires significantly higher temperatures and longer sintering times. Mayo (1996) developed a modified sintering law that directly accounts for pore size effects on densification rate: 1 dp 1 1 (-Q) c~ m _ exp p(1 - p ) dt dn r --~

(13.5)

where p is the density, d is the particle size, n is a constant dependent on the sintering mechanism, r is the pore radius, Q is the activation energy, R is the gas constant and T is the absolute sintering temperature. This equation predicts that the highest densification rate occurs for the finest pore size. It was found to hold throughout the sintering process and for a large range of pore sizes, at least when pore size distribution is uniform. This relationship has two implications. First, the pore size, in addition to grain size should be controlled during sintering. Fast sintering kinetics occur when pore sizes are small. Second, the densification rate is dictated by the instantaneous pore size, not only the initial pore size. Therefore, to maintain a fast sintering rate in late sintering stages, the pores should remain small even in late sintering stages. A small pore size throughout the sintering process is also critical in controlling the final grain size based on the pore pinning effect. For these purposes, a small and uniform pore population is desired in the green compact. Most often, such a pore distribution is associated with a high green density in non- or weakly agglomerated powders (Van de Graaf et al. 1985, Hahn 1993, Mayo 1996). The benefits of a small pore size and narrow size distribution in reaching high densities and reduced grain coarsening have been shown by many researchers in non- agglomerated A1203 (Messing and Kumagai 1994, Wu and de Jonghe 1996) Y203 and ZrO2 (Hahn 1993, Skandan 1995), SnO2 (Ahn et al. 1997) and TiN (Rabe and Wasche 1995) nanopowders.

13.4.1.3 Conventional sintering of nanopowders As seen in Table 13.2, sintering of nanosize powders starts at significantly lower temperatures than in conventional ones. The benefits of fine grain size in ceramic sintering have been reported 30 years ago and extensively used in the last decade and a half (Morgan 1968, Yan and Rhodes 1983, Winnubst et al. 1983, Mayo 1996). Initially, grain sizes in conventionally sintered parts exceeded 200 nm (Yan and Rhodes 1983, Winnubst et al. 1983). Lately, intensified efforts on sintering nanograined ceramic powders (ZrO2, A1203,

w

a\

0

Table 13.5. Density and grain sizes in nanoceramics consolidated by conventional (pressureless) sintering.

Material Zr02 ZrO2 zro2-3 mol% Y2O3 zro2-3 mol% Y203 zro2-3 mol% Y2O3 CeO2 CeO2-6 at.% Ca Ti02 TiO; TiO2-6% Y TiAl Ni

Sintering parameters Initial grain Temperature size (nm) (K)

6 to 9 6 to 9 NIA 8 8 10 to 15 10 to 15 14

-6 !I5 I (r204 25

1400 1250 1270 1273 1323 2 1623 1173 893 1 I73 723-773 13005

Final properties Time

Atmosphere

(h)

1.3 0.67 0.67 100

7 Air 10 C h i n 15 to 20 NR

15-20 2 31 Wmin

Air Vacuum Vacuum Air Air Near full Air Oxygen NR Oxygen Vacuum Vacuum

Density (%no)

Grain size (nm)

Full Near full Full >95 96 -100 Full 90 99 -98 >95 99

80 60
-

Reference Skandan ( 1995) Skandan (1995) Skandan (1995) Thuenissen et al. (1993) Thuenissen et al. (1993) Rahaman (1996) Hahn ef al. (1990) Kumar et al. (1992) Hahnetal. (1991) Chang et al. (1992) Ragulya and Skorokhod (1995)

1. Centrifugally consolidated; 2. Sol-gel prepared; 3. NR = Not reported; 4.Partial amorphous structure; 5. Rate controlled sintering.

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361

CdO and TiO2) by pressureless sintering resulted in fully dense ceramics with restrained grain growth (Rhodes 1981, Hahn et al. 1990, 1991, Kumar et al. 1992, Thuenissen 1993, Rahaman 1996). Examples of dense materials that have retained nanosize grains _<100 nm by conventional sintering are provided in Table 13.5. As seen in this table, full ceramic densification without grain coarsening has been achieved predominantly in the ZrO2 system. Conventional sintering has been less successful when applied to nanosize metals and intermetallics. For instance, only rate controlled sintering of pure Ni resulted in 99% dense specimen with 70-80 nm grain size (Ragulya and Skorokhod 1995).

13.4.1.4 Nano-composite densification Handwerker et al. (1989) summarized the sintering behavior of ceramic composites. Three different behaviors are identified in the sintering of viscous matrices. The common rule of mixture applies only when the volume fraction of the second phase is less than 10%. Scherrer's theory for viscous sintering with rigid inclusions applies between l0 and 12%. At higher volume fractions, deviations from both models are observed. For crystalline matrices, densification rates are lower than theoretical at all volume fractions indicating that the second phase retards matrix densification. Sintering of nanocomposites seems to adhere to the same mechanisms (Prabhu and Bourell 1995, Wittenauer 1996). The densification rate of the nanocomposite is not a linear function of the densification rate of the two components and each phase in the mixture has its own densification and coarsening rate. The nanocomposites sinter to densities less than monolithic ceramics and often final grain size is in the micron range (Wittenauer 1996, Anya and Roberts 1997). In a metal composite such as Cu-Nb, however, the Nb phase successfully restricted grain growth upon conventional sintering that achieved a high density and retained a grain size of 100 nm (Holtz and Provenzano 1994). 13.4.1.5 Grain-size control

The success in the consolidation of morphologically metastable powders is intimately related to the control of the competition between densification and coarsening. The driving force for densification is large due to the extreme curvature of nanoparticles (Eq. (13.2)). As analyzed by Raj (1987), when sintering progresses and grain boundaries are created, an additional term in the sintering driving force should be considered: 2 y b / d , where Yb is he grain boundary energy and d is the average grain size. This term is also the driving force for grain growth. When particles bond and grain sizes are small, the driving force for coarsening may be quite significant. The thermal stability of nanocrystalline materials has been the subject of recent reviews (Lange 1989, Suryanarayana 1995, Malow and Koch 1996a,b) or covered by a number of more extensive reviews (Flinn 1985, Mayo 1996, Weissmuller 1996, Lu 1996, Groza 1998). In general, grain coarsening is similar to that of conventional materials. The kinetics and activation energies are close to those in large grained materials with only a few exceptions (Thuenissen et al. 1993, Malow and Koch 1996a,b). The simplest way to control grain coarsening during sintering is to take advantage of the pinning effect of open pores. The influence of pores on grain coarsening has been well documented theoretically and experimentally, particularly in ceramics (Lange 1989, Hahn et al. 1990, Mayo 1996, Hague and Mayo 1997). The presence of open pores inhibits grain growth.

362

J. R. Groza

150 Ti02 TiO2~% Y

125 A

E 100-

C

.N

75-

C

.m

=m

50

Ni 25

. . . . . . . . . . .

20

l

1

30

40

. . . . . . .

-I- . . . . . . . . . . . .

T . . . . . . . . . . .

50 60 Density (%)

f

. . . . . . . . . .

70

T

. . . . . . . . . . .

80

-] .

.

.

.

90

.

I

100

Figure 13.1. Exaggerated grain growth in Ni and TiO2 when porosity becomes of closed type (from Hahn et al. 1990 and Ragulya and Skorokhod 1995). The effect of dopants on grain growth is shown for TIO2-6% Y (from Hahn et al. (1991)).

Experimentally, a strong correlation between the closure of open porosity and the onset of exaggerated grain growth was noted in both ceramic and metal nanopowders (Hahn et al. 1990, 1991, Andrievski 1994, Messing and Kumagai 1994). As seen in Fig. 13.1 for TiO2 and Ni, accelerated grain growth occurs at densities above ~90% or when the pores become closed. Often, the final grain size upon full densification is close to 1/zm, or the materials are no longer nanocrystalline. In contrast, dopants suppress pore breakaway by providing a drag force on the moving boundaries or, forming secondary phase particles (Hahn et al. 1991, Thuenissen et al. 1993, Rahaman 1996). For instance, Hahn et al. (1991) found that 6% Y addition to TiO2 retarded grain coarsening to >95% density, thus retaining a grain size less than 100 nm (Fig. 13.1). Perhaps one of the most dramatic effects of dopants is the addition of 6% Ca to CeO2 that retained a 30 nm grain size after full sintering by heating to 1625 K with a heating rate of 10 K/min (Rahaman 1996) (Table 13.5). In the same sintering condition, undoped CeO2 reached a grain size of 400 nm. Grain boundary chemical segregation or ordering have been very effective in suppressing grain growth of milled intermetallics (Gao and Fultz 1994). If no dopants are desired because of their negative effect on the final properties (e.g. on dielectric or mechanical properties), sintering under pressure may be another effective means to suppress grain growth. This approach will be discussed in the next section.

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363

13.4.2 Pressure consolidation of metastable powders The classical temperature-pressure trade-off for sintering brings distinct advantages to the consolidation of far-from equilibrium powders by enabling a low-temperature exposure and thereby making the retention of the metastable condition more likely. Pressure assisted methods for consolidation of metastable powders have been mainly developed for consolidation of RS materials. A schematic of the temperature-pressure combination for various pressure consolidation methods is presented in Figure 13.2. For RS powders, conventional processes such as extrusion, pressing, rolling or even hot isostatic pressing (HIP) have been less successful in the retention of metastable phases since high temperatures and long times are involved. Higher pressures are usually desirable to promote densification by plastic flow, shorten the diffusional densification and thus minimize microstructural changes. Therefore, high pressure consolidation methods have been favored. The pressure levels span a large range from low (< 100 MPa) in hot pressing to moderate (100-500 MPa) in hot pressing and HIP and high (>0.5 GPa) in high pressure methods such as piston cylinder and belt type apparatuses. Generally, the capability to collapse the pores in pressure-assisted sintering scales with the shear stress level. A shear stress is also beneficial for the mechanical disruption of surface oxide layers which provides better interparticle bonding for subsequent densification. The applied mean stress, known as the hydrostatic stress, induces new plasticity-driven sintering mechanisms, as well as stress-assisted diffusion. Both mechanisms enhance the densification rate. The shear component of the applied stress causes particle rearrangement and collapse of large pores. The particle rearrangement and macroscopic deformation of pores increase the number of particle contacts. This different pore size and distribution, which is more amenable to pore removal than in pressureless sintering, is the key factor in using pressure-assisted methods for the consolidation of metastable powders (Groza 1993, Hague and Mayo 1997). A schematic description of the shear stress and pressure components for various consolidation methods is illustrated in Figure 13.3. Among the pressure-assisted compaction methods, hot isostatic pressing (HIP) has been well developed to produce fully dense microstructures under simultaneous application of pressure and temperature with minimal additions of sintering aids. However, the typical thermal exposure during HIP is around two-thirds of the material melting point (Fig. 13.2) for a time on the order of hours. In the context of preserving the metastable structures, this high temperature exposure almost always causes significant microstructural changes. Therefore, further reducing the high temperature excursion while densifying materials has been a major goal for developing new consolidation processes. These new processes combine higher pressures with lower temperatures and shorter times than HIP and are aimed at preserving the unique microstructures and properties of metastable materials. Some of these new processes are quasi-HIP (Q-HIP) processes, field activated sintering technique (FAST) or high pressure methods (Fig. 13.3). Hot pressing (HP) and sinter forging involve a uniaxilly applied pressure in a die for hot pressing or with no die in sinter forging. In HIE a fluid transmits a hydrostatic stress. Variants of HIP with quasi-isostatic stress distribution may be obtained if a fluid-like pressure transmitting medium is used such as in the Ceracon process (Anderson and Groza 1987). Extremely high pressure levels (>5 GPa) are characteristic for diamond anvil, torsion-pressure method (sometimes

364

J. R. Groza

0.8TM--

(TEMP.)

RT_

I .1

1.0

I PRESSURE (GPa)

HP = Hot Pressing. HIP = Hot Isostatic Pressing. Q - HIP = Quasi Hot Isostatic Pressing. SPDC = Severe Plastic Deformation Consolidation.

10.0 FAST= Field Activated Sintering Technique. RT= Room temperature. TM- Melting temperature.

Figure 13.2. Schematiccomparisonof pressure consolidation techniques based on the temperature-pressure combination.

referred to as severe plastic deformation consolidation--SPDC) and shock compaction, also known as dynamic consolidation. The latter method is presented in the next chapter (Non-Conventional Sintering Methods).

13.4.3 Pressure-assisted consolidation methods

Most pressure-assisted consolidation methods have been applied for metastable powder densification: hot pressing, sinter forging, HIP, quasi-isostatic pressing, extrusion, and high pressure techniques. To illustrate the high pressure effects, Hahn (1993) fully sintered

365

Powder Consolidation SHEAR

PRESSURE

SHEAR

PRESSURE

m,

"

Extrusion Sinter Hot Q-HIP Forging pressing and FAST *Dynamic Compaction.

HIP

High *D,C, Pressure SPDC

Figure 13.3. Schematic diagram of shear and pressure stresses in various pressureassisted consolidation methods. For clarity, only average pressure values have been plotted in some cases (e.g., extrusion, forging).

TiO2 at 725-825 K ("~0.35Tm) by applying 1 GPa while retaining the morphologically metastable structure (i.e., with no grain growth). Similarly, Araki et al. (1995) densified mechanically alloyed Al-10.7at% Ti powders to 98% density under 2 GPa at 573 K with virtually no grain growth and retention of the initial solution supersaturation. Select examples of pressure-assisted methods applied to nanopowders that were densified to near (>95%) or full densities while retaining metastable features are given in Table 13.6. Some limitations in using pressure consolidation are related to ceramic specimen fracture at high

366

J. R. Groza

strain rates at low temperatures in sinter forging or the lack of adequate encapsulation materials with softening temperatures close to ceramic sintering temperatures for HIP (Mayo 1996, Hague and Mayo 1997). The classical pressure-assisted sintering diagrams developed by Ashby and coworkers have been used to describe metastable powder consolidation (Arzt et al. 1983, McKimpson 1996). In nanocrystalline powders, a change in the sintering mechanism from stress-assisted densification to curvature driven grain boundary diffusion has been observed (Bourell et al. 1993, Shaik and Milligan 1997).

13.4.4 Non-conventional sintering methods A number of non-conventional consolidation methods have been applied to metastable powder densification: microwave sintering, shock or dynamic consolidation, field assisted sintering. The main purpose in using these methods is to enhance densification, thus reducing the sintering temperature or time with the ultimate benefit of preserving the initial metastable structure such as a final fine grain size.

13.4. 4.1 Microwave sintering The processing time in microwave sintering is reduced as compared to conventional heating from external sources due to the direct energy coupling with electric dipoles within the heating body (Clark and Sutton 1996). These shorter times bring about benefits in energy savings and final properties. Temperature gradients are reduced and, most importantly, an overall short sintering time minimizes grain growth. In addition, rapid heating rate can bypass the low-temperature region where the rate of grain growth is higher than the rate of densification. Since grain boundaries are the primary sites for electric dipoles, microwave sintering appears particularly attractive for the densification of nanocrystalline powders. The method has been applied to ceramic particles both micron and nanocrystalline size, such as TiO2 and A1203 (Freim et al. 1994, Lewis et al. 1997). In these materials, only densities less than 95% have been achieved when the grain size was less than 100 nm. Contrary to expectations, the density-grain size combinations in y-A1203 were almost identical to those achieved in more conventional pressureless sintering experiments at the same temperature (Freim et al. 1994). The sequence of phase transformation in A1203 was also similar to that in conventional sintering.

13.4.4.2 Field-assisted sintering Electric feld application has been shown to enhance powder densification thereby achieving higher densities at lower temperatures or in shorter times. Variants of field sintering that involve pulsed current discharge and resistance heating, sometimes termed plasma activated sintering (PAS) or field-assisted sintering technique (FAST), plasma pressure consolidation (PPC), spark plasma sintering or pulse electro-discharge consolidation, have been effectively applied to metastable powder consolidation (Kimura and Kobayashi 1994, Yoo 1996, Mishra et al. 1996, Sudarshan and Yoo 1997). The enhanced densification in FAST is most noticeable at lower temperatures or when multiple discharges of current through the powder compact are applied. Generally, it is believed that pulsed current application promotes the removal of adsorbates or surface oxides due to a pos-

Table 13.6 Pressure-assisted consolidation methods for sintenng (near) fully dense metastable materials

Temperature

Pressure

Material

Consolidation method

(K)

(MPa)

Zr02 Ti02 Ti02 Ti02 CeO2 TiAl Fe- 10% Cu Al- 10%Ti Fe(Co, Ni)B Cu-NbC

Sinter forging Sinter forging Hot pressing Hot pressing High pressure HIP Sinter forging Hot pressing Hot pressing HIP

1223 923 753 873 873 1000 800 623 670 1273

300 < 100

700 500 1100 3 10 525 500 1000 17.5

Time (h)

Metastable feature or final grain size (nrn)

Reference

2.5 to 3 -1 NR 2 NR 4 0.03 0.5 0.5 1

45 60 40 46 10 25-50 45 -75 Amorphous <60

Skandan (1995) Averback et al. (1993) Hahn and Averback (1992) Tenvilliger and Chang (1993) Lavik et al. (1996) Froes et al. (1995) Shaik and Milligan (1997) Guoxian et al. (1995) Saida et al. (1994) Murphy and Courtney (1994)

368

J. R. Groza

sible plasma generation (Yoo 1996). This enhances the sinter bonding of particulates. For non-conductive ceramics, a tentative explanation of the enhanced densification by multiple discharges is related to the increase in the dielectric constant with the temperature level at which the discharge was applied (Mishra et al. 1996). FAST sintering for 3 minutes at 725 K of nanosize Fe-85% Fe3C achieved 99% density with a final grain size of 45 nm (Goodwin et al. 1997). This is to be compared to HIPping of the same powders at 1025 K for 60 minutes which yielded the same density but with a final grain size of 87 nm. Kimura and Kobayashi (1994) fully sintered mechanically alloyed TiA1 powders and retained nanosize grains by spark sintering at 1051-1312 K under modest pressures (29-147 MPa). Similar to conventional sintering, the densification temperature inversely scaled with the applied pressure. An alternate electric pulsed power method induces short high pressure pulses that result in powder densification such as in dynamic magnetic compaction (DMC). Ivanov et al. (1995) achieved up to 5 GPa pressure and thus consolidated alumina powders to 83% density while retaining a nanocrystalline structure. The DMC method was applied to sintering of large parts of nanocrystalline iron and alumina (Chelluri et al. 1997). 13.4.4.3 Shockwave consolidation Dynamic or shockwave consolidation proceeds with the passage of a large-amplitude compressive stress generated by plate impact or explosion without any external heating (Kondo and Sawai 1990, Korth and Williamson 1995). The peak pressure values may be on the order of tens of GPa, thus providing densification by plastic yielding for both metals and ceramics. Localized heating possibly up to melting temperatures due to particle friction occurs and enables good interparticle bonding. In nanosize powders the heat may transfer throughout the entire particle, thus providing an advantage over coarser materials where the heating is only superficial. Best results are achieved when high temperatures are reached before the shockwave passes. If particles are heated, they may deform rather than fracture upon stress application. This very short high temperature exposure provides the best means to retain fine grain size or far-from equilibrium conditions such as amorphous structures (Glade and Thadhani 1995), or supersaturated solid solutions (Dogan et al. 1994). Full densification was also reported in mechanically alloyed TiA1 specimens with final grain sizes of 15 nm (Suryanarayana et al. 1997). These latter results are compared with HIPping that provided full consolidation at 1075 K/207MPa/2h but grain sizes were about 100 nm. The major drawback in the shockwave consolidation is the difficult coordination of the short stress and heat application events that result in frequent specimen fracture (cracking).

13.5. C O N C L U D I N G R E M A R K S

Densification of far-from equilibrium powders presents challenges related to the need to retain the metastable microstructures in addition to usual pore elimination. Lower sintering temperatures or shorter durations, pressure application and a more rigorous control of both powder synthesis (agglomeration and contamination) and pre-sintering stages (cold compaction) have been applied to successfully densify and prevent undesirable transformations of the initial metastable structures. For the most recent far-from equilibrium

Powder Consolidation

369

powders with particles in the nanometer range, the main features influencing the sintering process are the high driving force and enhanced kinetics due to large curvature effects. Densification of nanopowders takes place at temperatures consistently below those of larger grained powders by up to several hundreds of degrees. There are numerous benefits from using lower sintering temperatures: small final grain sizes, a better retention of metastability and elimination of sintering aids. Studies of nanopowder densification have shed more light into sintering issues such as powder agglomeration, surface condition or contamination, pore role in sintering and grain growth. The resultant control of synthesis (e. g., non-agglomerated nanopowders) and processing has enabled fabrication of fully dense nanocrystalline parts particularly ceramic, even by conventional sintering. The recent efforts to improve the understanding of pressure effects in nanopowder consolidation have reflected positively on their final densification results. One notable example is sinter forging. Field assisted sintering seems to offer benefits in nanopwder sintering with restricted grain growth. Although some distinct differences in the densification of nano- versus micron grained powders seem to emerge and a better understanding of nanosintering has been accomplished, the specific effect of densification variables on final density and properties of nanomaterials is not well understood yet. The interdependence of the sintering behavior on size, chemistry and structure of nanoscale particles is an area of worthwhile future research. More fundamentally, there is a tremendous need for specific materials data in the nanoregime, such as diffusion coefficients and surface energies.

ACKNOWLEDGMENTS

The author's research on nanocrystalline materials densification has been supported by National Science Foundation, Chemical and Transport Systems (grant # CTS-9632280) and Manufacturing Processes and Equipment Program (grant # DMII - 9532072). The author wishes to thank all those who provided her with unpublished results or preprints of their recent work and helpful discussions.

REFERENCES Ahn, J.-P., Huh, M.-Y. and Park, J.-K. (1997) Nanostr. Mater., 8, 637. Alymov, M. I., Maltina, E. I. and Stepanov, Y. N. (1992) Nanostr. Mat., 4, 737. Anderson, R. L. and Groza, J. R. (1987) Metal Powder Report, 42, 678. Andrievski, R. A. (1994) Internat. Powder Metall., 30, 59. Anya, C. C. and Roberts, S. G. (1997) J. Euro. Ceram. Soc., 17, 565. Araki, H., Saji, S., Okabe, T., Minamino, Y., Yamane, T. and Miyamoto, Y. (1995) Mater. Trans. JIM, 36, 465. Arzt, E., Ashby, M. E and Easterling, K. E. (1983) Metall. Trans. 14A, 211. Averback, R. S., H6fler, H. J. and Tao, R. (1993) Mater. Sci. Eng., A166, 169. Averback, R. S., Zhu, H., Tao, R. and H6fler, H. (1996) in Synthesis and Processing of Nanocrystalline Powder, ed. Bourell, D. L. (TMS, Warrendale), p. 203. Balachandran, U., Siegel, R. W., Liao, Y. X. and Askew, T. R. (1995) Nanostr. Mater., 5, 505. Bonetti, E., Valdre, G., Enzo, S., Cocco, G. and Soletta, I. (1993) Nanostr. Mater., 2, 369. Bonevics, J. E. and Marks, L. D. (1992) J. Mater Res., 7, 1489. Bourell, D. L. and Kaysser, W. A. (1994) Metall. Mater. Trans., 25A, 677. Bourell, D. L., Parimal and Kaysser, W. A. (1993) J. Amer. Ceram. Soc., 76, 705.

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Chapter 14 Bulk Amorphous Alloys 14.1. History of Bulk Amorphous Alloys 14.2. Dominant Factors for High Glass-Forming Ability 14.3. Continuous Cooling Transformation of Alloys with High Glass-Forming Ability 14.4. Preparation Methods and Elemental Effect 14.5. Structural Relaxation and Glass Transition 14.6. Physical Properties 14.6.1 Density 14.6.2 Electrical Resistivity 14.6.3 Thermal Expansion Coefficient 14.7. Mechanical Properties 14.8. Viscoelasticity 14.9. Soft Magnetic Properties 14.9.1 Formation and Soft Magnetic Properties of Bulk Amorphous Alloys 14.10. Viscous Flow and Micro-Formability of Supercooled Liquid 14.10.1 Feature of Phase Transition of Bulk Amorphous Alloys 14.10.2 Deformation Behavior of Supercooled Liquid 14.10.3 Micro-Forming of Supercooled Liquid 14.11. Applications and Future Prospects References

375 376 380 383 387 392 392 393 394 395 399 403 403 406 409 409 412 413 413

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Chapter 14 Bulk Amorphous Alloys AKIHISA INOUE

14.1. H I S T O R Y OF B U L K A M O R P H O U S A L L O Y S

Since the synthesis of an amorphous phase in the Au-Si system by rapid solidification by Klement et al. in 1960 (Klement et al. 1960), a great number of scientific and engineering data for amorphous alloys have been accumulated. As a result, it has been shown that amorphous alloys have the features of new alloy compositions and new atomic configurations that are different from those for crystalline alloys. These features have enabled the exhibition of various characteristics such as good mechanical properties, useful physical properties and unique chemical properties (Chen 1980, Masumoto, 1982, Luborsky 1983, Liebermann 1993) which have not been obtained for conventional crystalline alloys. Recently, amorphous alloys have also attracted increasing interest as a precursor to produce nanocrystalline alloys by crystallization because of the appearance of good mechanical properties (Kim et al. 1990, Chen et al. 1991), soft magnetism (Yoshizawa et al. 1988, Suzuki et al. 1991), hard magnetism (Croat 1980, Kneller and Hawig 1991, Inoue et al. 1995), high magnetostriction in low applied fields (Inoue et al. 1994a), and high catalytic properties (Yokoyama et al. 1981). Based on these results, scientific and engineering importance of the amorphous and nanocrystalline alloys has steadily increased for the last three decades. When attention is paid to bulk amorphous alloys with low critical cooling rates (Rc), it is known that Pd40Ni40P20 (Chen 1976) and Pd76Cu6Si18 (Chen 1978) amorphous alloys can be produced in bulk form with diameters up to 3 mm and 0.3 mm, respectively, by water quenching. Subsequently, a flux treatment using a B203 medium for the Pd-Ni-P alloy was found to be effective for increasing the maximum sample thickness, tmax to about 10 mm (Kui et al. 1984, 1985). Thus prior to 1990, except for the Pd-metalloid systems, no other bulk amorphous alloys have been synthesized because Re needed to be above 105 K/s. Recently, a number of bulk amorphous alloys with Re below 103 K/s have been synthesized in multicomponent systems of Mg-Ln-TM (Inoue et al. 1988), Ln-A1-TM (Inoue et al. 1989) (see also some other studies of such alloys (Inoue et al. 1989), Zr-A1-TM (Inoue et al. 1990, 1993), Ti-Zr-TM (Amiya et al. 1994), Zr-(Ti, Nb, Pd)-A1-TM (Inoue et al. 1995a), Zr-Ti-TM-Be (Peker and Johnson 1993), Pd-Cu-Ni-P (Inoue et al. 1996a), Fe-(A1, Ga)-(P, C, B, Si, Ge) (Inoue and Gook 1995, Inoue et al. 1995b), Pd-Ni-Fe-P (He and Schwarz 1998), Fe-(Co, Ni)-(Zr, Nb, Ta)-B (Inoue et al. 1997e), Fe-(Co, Ni)-(Zr, Nb, Ta)-(Mo, W)-B (Inoue et al. 1997f) and CoFe-(Zr, Nb, Ta)-B (Inoue et al. 1997e) (Ln = lanthanide metal, TM = transition metal). 375

A. Inoue

376

Table 14.1. Alloy systems, year of discovery, and maximum sample thickness (tmax)of new multicomponent alloys with high glass-forming ability Alloy systyem Non-Ferrous systems Mg-Ln-M (M = Ni, Cu, Zn) Ln-A1-TM (TM = VI-VIII group transition metals) Ln-Ga-TM Zr-A1-TM Ti-Zr-TM Zr-Ti-TM-Be Zr-(Ti, Nb, Pd)-AI-TM Pd-Cu-Ni-P

Year

tmax (mm)

1988 1989

10 10

1989 1990 1993 1993 1995 1996

10 30 3 30 30 75

Alloy system

Year

tmax (mm)

Pd-Cu-B-Si Co-(A1,Ga)-(P, B, Si) Co-(Zr,Hf, Nb)-B

1996 1996 1996

10 1 1

Ni-(Zr,Hf, Nb)-B

1996

1

Ferroussystems Fe-(AI,Ga)-(P, C, B, Si, Ge) Fe-(Nb,Mo)-(A1, Ga)-(P, B, Si) Fe-(Zr,Hf, Nb)-B

1995 1995 1996

2 2 5

Ln = Lanthanide metal; TM = Transition metal.

Table 14.1 summarizes the alloy systems of the bulk amorphous alloys, the year of their discovery, and tmax. The tmax reached about 30 mm (Inoue and Zhang 1996) for the Zr-A1TM alloys, 25 mm (Johnson 1996) for the Z r - T i - T M - B e alloys developed by Johnson's group (see also Johnson 1996), 72 mm (Inoue et al. 1997c) for the P d - C u - N i - P alloy, 25 mm (He and Schwarz 1998) for the P d - N i - F e - P alloy and 6 mm (Inoue and Zhang 1996) for the F e - C o - Z r - M o - W - B alloy. This chapter reviews recent results on the fabrication, fundamental properties, workability and applications of the new bulk amorphous alloys with thicknesses above several millimeters that were mainly obtained by our group during the last eight years.

14.2. D O M I N A N T F A C T O R S F O R H I G H G L A S S - F O R M I N G

ABILITY

For the formation of an amorphous phase by rapid solidification it is essential to suppress the nucleation and growth reactions of a crystalline phase in the supercooled liquid region between the melting temperature (Tm) and the glass transition temperature (Tg). Rc has been reported (Davies 1983) to be above 104 K/s for Fe-, Co- and Ni-based amorphous alloys and above 102 K/s for Pd- and Pt-based amorphous alloys, being much higher than those for oxide glasses. Recently, we have succeeded (Inoue 1995a,b, 1996, 1997a,b) in finding new multicomponent amorphous alloys with much lower Rc ranging from 0.1 K/s to several hundred K/s, as shown in Fig. 14.1 (a). Simultaneously, the tmax increases drastically from several millimeters to about one hundred millimeters. It is to be noted that the lowest Rc and the largest tmax are comparable to those for oxide and fluoride glasses. Figure 14.1 (a) also shows that the new amorphous alloys have a high reduced glass transition temperature (Tg/Tm) of above 0.60. Furthermore, the new amorphous alloys have a much wider supercooled liquid region which is defined by the difference between crystallization temperature (Tx) and Tg, (ATx = Tx - Tg), as shown in Fig. 14.1(b). From the empiri-

Bulk Amorphous Alloys

377

cal relations among Rc, tmax and Tg/Tm or A Tx, it may be noted that high glass-forming ability (GFA) is obtained when an alloy exhibits (1) high Tg/Tm, and (2) large A Tx. As summarized in Table 14.1, the bulk amorphous alloys consist of multicomponent systems containing Mg, Ln, Zr, Pd, Fe or Co as the main constituent element. Based on the multicomponent amorphous alloys with high GFA, we have proposed (Inoue 1995a,b, 1996, 1997a,b) three simple empirical rules for achieving high GFA for metallic alloys. These are: (1) the alloy system should contain more than three elements, (2) a significant difference in atomic size ratios (above about 12%) should exist among the main constituent elements, and (3) the component elements should exhibit negative heats of mixing among themselves. The importance of the above three empirical rules can be understood from a number of experimental data as well as theoretical aspects. Various models have been proposed till to-date to describe the glass formation and GFA (Davies 1983). These models can be grouped into three categories of thermodynamics, kinetics, and structure, depending on the factors that are viewed as decisive in the formation of amorphous alloys. From these factors it is possible to understand why the above-described ternary and higher-order amorphous alloys have low Re. Firstly, we investigate the reason for the high GFA from the thermodynamical point of view. The high GFA is obtained when the free energy AG(T) for the transformation of liquid to crystalline phase is low. In the relation A G = AHf - T ASf for Gibbs free energy, a low A G value is obtained when A Hf is small and/or A Sf is large. Here &Hi and ASf are the enthalpy of fusion and the entropy of fusion, respectively. Large ASf is expected to be obtained in multicomponent systems because ASf is proportional to the number of microscopic states. The free energy at a constant temperature also decreases in the case of low chemical potential caused by the low enthalpy and high Tg/Tm as well as of high liquid/solid interfacial energy. Based on these thermodynamical aspects, it is concluded that the multiplicity of alloy components leading to an increase of A Sf increases the degree of dense random packing which is favorable for the decrease of A H f and the increase of solid/liquid interfacial energy. This interpretation is consistent with the present result that a high GFA has been obtained for the above-described multicomponent systems. Furthermore, the multicomponent amorphous alloys have been recently shown to have high packing fractions as is evidenced from low specific volumes shown in section 7. Secondly, the reason for the high GFA is investigated from the kinetic point of view. The homogeneous nucleation and growth of a crystalline phase with a spherical morphology from the supercooled liquid are expressed by the following relations (Turnbull 1956): I = 103~149 exp[-bot3fi/(Tr(1 U = 1 0 2 f / r ] 9[1 - exp[-fl

-

Tr) 2]

[cm -3 s -1]

ATr/Tr(T/Tm)]]

[cms -1]

(14.1) (14.2)

Here Tr is the reduced temperature (T/Tm), ATr is the difference in temperature from Tm, b is a shape factor (e.g., 16Jr/3 for a spherical nucleus), r/is viscosity and f is the fraction of nucleus sites at the growth interface, ot and fi are dimensionless parameters

378

A. Inoue

108

i

'

107 106

I

'

i

'

i

~ ~ o 2 . " 0.1

105 Fe-,Co-,Ni-base ~ "7

~'-"

104-103 Pd',Pt'base " I w ' ~ Ln-AI-TM Mg-TM-Ln Zr-AI-TM Zr-AI-Ni-Cu

102 _

,0'-

~ \ \ ~,

Fe-A1-Ga-P-C-B Pd-Cu-Ni-P Fe-Zr-Nb-B Co-Zr-Nb-B Ni-Zr-Nb-B

10° _ 10-1 10-2 _

1

\~o \ ~ \ ~ \

,.a

G~.

10

\ w

\

100

10-3 10"4

105

0.4 t

'

i

'

,

0'.5

i

,

' 016 Tg/Tm

i

,

i

'017

'

i

0.8

,

i

O.1

O FesoPt3C7 104 103

~

102 _ 101

_

10 o _

Sii8 d77Cu68il7 ~ oFe-A1-Ga-P-C-B Pd40Cu40P20 P_t60N!20P20 "X~ oFe-Co-Zr-Nb-B LassAIEsNi2o ..4~ Zr60AllsNi25/ 0 Mg-Cu-Y ~ 10 Zr-AI-Ni-Cu-Pd~ o Pd-Cu-Ni-P ~ o a/ , o ~x (non-fluxed) Zr-Al-eu-Ni Pd-Cu-Ni-P o (fluxed)

10 -1 _

10-2 0

2 "0

' 4 "0

'

6 "0

'

8"0

'

100 '

E

100

'

120 '

ATx(=Tx-Tg)/K F i g u r e 14.1. Relation between the minimum critical cooling rate for glass formation (Rc), maximum sample thickness (tmax) and the reduced glass transition temperature (Tg/Tm) or the temperature interval of supercooled liquid region A T x ( = Tx - Tg) for the new multicomponent amorphous alloys in the lanthanide (Ln)-, Mg-, Zr-, Pd-, Fe-, Co- and Ni-based systems.

Bulk Amorphous Alloys

379

related to the liquid/solid interfacial energy (or). A H f and ASf c a n be expressed as c~ = ( N o V ) I / 3 a / A H f and fl = A S f / R . Here No, V and R are Avogadro number, atomic volume and gas constant, respectively. In these relations, the important parameters are r/, ot and ft. The increase in the three parameters decreases the I and U values, leading to an increase of GFA. The increase of ot and/3 also implies the increase of cr and A Sf and the decrease of A H f , being consistent with the interpretation on the achievement of high GFA derived from the thermodynamical point of view. Furthermore, one can notice that r/ is closely related to Tg/Tm value and ot3fl reflects the thermal stability of the supercooled liquid. The importance of the ot3fl parameter can be understood from the following two examples: (1) when the value of or/31/3 exceeds 0.9, unseeded liquid would not in practice crystallize at any cooling rate, and (2) when ~fll/3 is below 0.3, it would be impossible to suppress crystallization. Furthermore, the reason for the high GFA of the particular multicomponent alloys is investigated from the structural point of view. As described above, the present multicomponent alloys consist of elements with a significant difference in atomic sizes and negative heats of mixing. The atomic sizes of the constituent elements in the ternary amorphous alloys can be classified into three groups of larger, intermediate and smaller, as shown in Fig. 14.2. The combination of the significant difference in the atomic sizes and negative heats of mixing is expected to induce a higher dense random packing in the supercooled liquid which enables the achievement of high liquid/solid interfacial energy as well as the difficulty of atomic rearrangements due to low atomic diffusivity and high viscosity. The formation of a higher dense random-packed structure has also been confirmed from the X-ray diffraction profiles of the ternary La-A1-Ni (Matsubara et al. (1990), M g - N i La (Matsubara et al. 1992a) and Zr-A1-Ni (Matsubara et al. 1992b) amorphous alloys. Neither a prepeak at the lower angle side nor any splitting of the second peak is seen in the X-ray interference functions (Matsubara et al. 1990). The features of the X-ray interference functions are analogous to those for a liquid phase, indicating that the ternary amorphous alloy has a homogeneously mixed atomic configuration corresponding to the formation of a higher dense random packed structure. The increase in the random packing density has also been supported from the very small change in density of the multicomponent amorphous alloys upon crystallization (Inoue et al. 1998a). The densities of the Zr60All0Nil0Cu20 and Pd40Cu30Nil0P20 alloys are 6.72 and 9.27 Mg/m 3, respectively, for the amorphous state and 6.73 and 9.31 Mg/m 3, respectively, for the crystalline state. The change in density upon crystallization is evaluated to be 0.30 and 0.54%, respectively, which is much smaller than that (about 2%) (Chen 1980, Masumoto 1982, Masumoto et al. 1976) for ordinary amorphous alloys. We have also evaluated the changes in the atomic distances and coordination numbers among the constituent elements upon crystallization of the ternary amorphous alloys from the radial distribution function (RDF) data obtained by the anomalous X-ray scattering technique (Matsubara et al. 1990, 1992a,b), Table 14.2 summarizes the atomic distances and coordination numbers of each atomic pair for an amorphous Zr60A115Ni25 alloy in the as-quenched and the crystallized states (Matsubara et al. 1992). Although no significant changes in the coordination numbers and atomic distances of Ni-Zr and Zr-Zr pairs upon crystallization are recognized, one can notice the drastic change in the coordi-

A. Inoue

380 5

~

4

Ln-A1-TM

~.

~.

Zr-A1-TM

"--

-"

-"

Mg-Ln-TM

--

---

--

Fe-Zr-B

.

~ 3

~,

.

.

.

L.

b

~

1

Figure 14.2. Atomic radii of the constituent elements for the new ternary amorphous alloys.

nation number of Zr-A1 atomic pair upon crystallization. The significant change indicates that the local atomic configurations in the amorphous alloy are significantly different from those in the corresponding equilibrium crystalline phase. This result indicates the necessity of long-range rearrangement of A1 around Zr in the progress of crystallization. The long-range atomic rearrangement is difficult in the higher dense random packed amorphous phase. Based on the above-described results and discussion, Fig. 14.3 summarizes the reasons for the achievement of high GFA of the multicomponent amorphous alloys in Ln-, Mg-, Zr-, Ti- and Hf-based systems (Inoue 1995a,b, 1996, 1997a,b). It is again confirmed that the high GFA is attributed to the formation of a new kind of supercooled liquid with a higher dense random packing density, new short-range atomic configurations and long-range atomic interactions resulting from the multicomponent elements with different atomic sizes and negative heats of mixing. The peculiar liquid structure can have a higher liquid/solid interfacial energy leading to the suppression of nucleation of a crystalline phase, a higher viscosity leading to the increase of Tg/Tm and a lower atomic diffusivity leading to the difficulty of growth of a crystalline phase.

14.3. C O N T I N U O U S GLASS-FORMING

COOLING

TRANSFORMATION

OF ALLOYS

WITH HIGH

ABILITY

It is important to know Rc for evaluating the GFA and the production of bulk amorphous alloys. The use of multicomponent alloys has enabled the measurement of continuous cooling transformation (CCT) curve by using different kinds of instruments, e.g., thermocouple, pyrometer, and DTA. In particular, a new Pd-Cu-Ni-P bulk amorphous alloy

Bulk Amorphous Alloys

381

Table 14.2 Atomic distances (r) and co-ordination numbers (N) calculated from (a) the ordianry radial distribution function (RDF), and from the environmental RDFs for (b) Zr and (c) Ni in the as-quenched amorphous and crystallized Zr60All5 Ni25 amorphous alloys

As-quenched amorphous Crystalline (Annealed for 180 ks at 730 K

(a) (b) (c) (a) (b) (c)

rZrNi (nm)

NZrNi

rZrZr (nm)

NZrZr

NZrAI

0.2674- 0.002 0.2674- 0.002 0.269+ 0.002 0.2684- 0.002 0.2674- 0.002 0.2734- 0.002

2.3 4- 0.2 2.1 4- 0.5 2.3 4- 0.2 3.0 4- 0.2 3.0 4- 0.5 2.3 4- 0.2

0.317 4- 0.002

10.3 4- 0.7

-0.1 4- 0.9

0.322 4- 0.002

8.2 4- 0.9

0.8 4- 0.7

Reasons for the Large glass-Forming Ability of the Mg-, Ln-, Ti-, Zr- and Hf-based Alloys The constituent elements more than three kinds with large atomic size ratios above 12 % and neg]ative heats of mixing

V I n c r e a s e in the degree of dense r a n d o m p a c k e d s t r u c t u r e (topological and chemical points of view) F o r m a t i o n of liquid with new atomic configurations and m u i t i c o m p o n e n t interactions on a s h o r t - r a n g e scale

Increase of solid/liquid interfacial e n e r g y

Suppression of nucleation phase o f a crystalline . ~ ~

Difficulty of atomic r e a r r a n g e m e n t (decrease of atomic diffusivity, X, increase of viscosity )

~

~

.[. V ~

e

~

Necessity of atomic r e a r r a n g e m e n t on a long-range scale for crystallization

!

Suppression of g r o w t h line phase

Decrease of Tn~ Increase of T ~rF m Figure 14.3. Reasons for the achievement of a high glass-forming ability for some temary alloy systems such as Ln-A1-TM, Mg-Ln-TM and Zr-A1-TM, etc. (Ln = lanthanide metal, TM = transition metal).

382

A. Inoue

was produced by slow cooling in the DTA where the cooling behavior of the molten alloy was easily measured in a wide cooling rate range and the precipitation of a crystalline phase could be detected on the DTA curve. Figure 14.4 shows the CCT curve of the Pd40Cu30Nil0P20 alloy obtained at various cooling rates in the DTA (Nishiyama and Inoue 1996). The amount of the ingot was 2 g and the cooling rate was changed in the range from 0.033 to 3.74 K/s. On the assumption that Re is the minimum cooling rate necessary to avoid the intersection with the CCT curve for the liquid, Re of the Pd-CuNi-P alloy was measured to be 1.58 K/s (Nishiyama and Inoue 1996) for the non-fluxed melt and 0.1 K/s (Nishiyama and Inoue 1997) for the BzO3-fluxed melt. The flux treatment was effective in decreasing Re because the molten alloy was cleaned by the flux treatment and heterogeneous nucleation sites were reduced (Chen 1978, Kui et al. 1984, 1985). The Re value is believed to be the lowest for all metallic amorphous alloys reported up to date. The effect of the flux treatment on the suppression of heterogeneous nucleation was evaluated by comparing the present experimental CCT curve with the predicted time-temperature-transformation (TTT) curve which can be calculated from the homogeneous nucleation and growth theories (Nishiyama 1997, Christian 1965). When the homogeneous nucleation and growth rates of the crystal are represented by Ivhomo and Uc, respectively, the volume fraction (X) crystallized in time (t) is expressed by Eq. (14.3) in the Johnson-Mehl-Avrami treatment of transformation kinetics. X - (rc/3)IvhomoU3t 4

(14.3)

The Ivhomo and Uc for glass-forming systems are given by Eqs. (14.1) and (14.2) (Uhlmann 1969), respectively. Beside, the relation between D and h is expressed by the following Stokes-Einstein equation. D = kT/3rraorl

(14.4)

By inserting Eqs. (14.1), (14.2) and (14.4) into Eq. (14.3), we obtain

t - - ( 9 . 3 0 / k T ) [ ( X a 9 / f 3 N v ) e x p ( 1 . O 7 / T r 3 ATr2)/{1- exp(--AH f

ATr/RT)}3] 1/4 (14.5)

t was calculated using the following experimental (Nishiyama 1997) and reference (Uhlmann 1972, Christian 1965) values: X = 10 -6, f = 0.2ATr, Nv = 8.01 x 102 atoms/cm 3, a0 -- 0.133 nm, A H f -- 4840 J/mol and r/-- 9.34 x 10 -3 e x p [ 4 1 3 5 / ( T 447)]. Here Nv and a0 are the calculated values from the atomic sizes and atomic percentages of the constituent elements for Pd40Cu30Nil0P20. The A H f and r/ are the experimental data for the Pd40Cu30Nil0P20 amorphous alloy. The relation between t and T obtained thus is also shown in Fig. 14.4, where the experimental data (CCT curve) of the Pd40Cu30Nil0P20 alloy are also shown for comparison. The calculated TTT curve lies at longer times and lower temperatures as compared with the experimental CCT curve. The nose point on the calculated TTT curve is located at T = 670 K and t = 104 s and Re is evaluated to be 1.34 z 10 -3 K/s which is smaller by two orders than that (Re =

383

Bulk Amorphous Alloys

1000

I

I

I

I

I... 800

-- ~ - ~

i"~/~.,~

7oo-

E

4)

!-

I

Te= 8 0 4 K

m L.

4) Q.

I

Pd4~176176176

900 4) L.

I

600 -

"

Calculated

.......

-

(TTT, Homog. nucleation) , Rc=3

30 x 10 -~ K/s

, ~

, /'~..

/

9

-

Non-fluxed (DTA) Fluxed (DTA)A) .57K/s Rc=0.100K/s . . . . . Tg= 592K . . . . .

Rc=t I

10 0

Ni3P

101

I

10 2

I

I

10 3

10 4

I

I

10 5

10 6

I

10 7

10 8

Time, t / s Figure 14.4. Continuous-cooling-transformation (CCT) curve of the Pd40Cu30Nil0P20 alloy subjected to the B20 3 flux treatment, in comparison with that for the same alloy without the flux treatment. The temperature-timetransformation (TTT) curve calculated on the basis of the homogeneous nucleation and growth theories is also presented for comparison.

0.1 K/s) for the fluxed Pd40Cu30Nil0P20 alloy. The significant difference in Rc implies that the B203 flux treatment cannot completely eliminate the heterogeneous nucleation sites in the Pd-based alloy liquid. If we can eliminate the heterogeneous nucleation sites completely, trnax is predicted to reach as large as about 1.5 m (Nishiyama and Inoue 1997) even for the Pd40Cu30Nil0P20 amorphous alloy. CCT curves were similarly obtained for the Zr-A1-Ni-Cu amorphous alloys. The Rc values have been reported to be about 10 K/s for the arc-melted ingot (Inoue et al. 1995f), 40 K/s for the unidirectionally solidified ingot (Inoue et al. 1994b) and 80 K/s for the ingot cast into a wedge-shape copper mold (Inoue et al. 1995c). The difference in Rc among the solidification processes reflects the difference in the ease of heterogeneous nucleation of a crystalline phase.

14.4. P R E P A R A T I O N M E T H O D S AND E L E M E N T A L E F F E C T

By choosing appropriate compositions with large A Tx, high Tg/Tm and low Rc, bulk amorphous alloys have been produced by various techniques, which are classified into two groups of solidification and consolidation. The different solidification techniques employed include water quenching, high pressure die casting, arc melting, copper mold casting, unidirectional casting and suction casting. The bulk amorphous alloys are also produced by hot pressing and warm extrusion of atomized amorphous powders in the su-

384

A. Inoue

percooled liquid state (Inoue 1995a,b 1996, 1997a,b). As examples, Fig. 14.5 shows the shape and morphology of bulk amorphous alloys in a cylindrical form with 17 mm diameter and 600 mm long for the Zr-alloy system and 75 mm diameter and 80 mm long for the Pd-alloy system. Here, it is important to understand the high thermal stability of the supercooled liquid during casting to examine the ease of the movement of the amorphous/crystal interface in the alloy. The quaternary and quinary Zr-A1-TM (TM = Co, Ni, Cu) bulk amorphous alloys can be made into buttons by arc melting on a copper hearth (Inoue and Zhang 1994). However, in this melting method, it is difficult to completely suppress the precipitation of crystalline phases because of the ease of heterogeneous nucleation resulting from incomplete melting at the bottom side in contact with the copper hearth. Here, it is to be noticed that only an amorphous phase is formed in the inner region where the heat flux reinforced from the copper hearth disappears and the cooling capacity is reduced. That is, the low cooling rate obtained by arc melting is high enough to form an amorphous phase in the Zr-A1-Co-Ni-Cu system. The crystalline phase can be divided into three different zones of outer equiaxed zone, columnar dendritic zone and metastable single phase zone, as illustrated in Fig. 14.6. It has further been confirmed that the interface between the amorphous and crystalline phases is very smooth, because of the satisfaction of the highly supercooled solidification condition defined by the criterion of Vi >> Di/(~i (Kurz and Fisher 1989). Here Vi is the velocity of the moving liquid/solid interface, Di is the

Figure 14.5. Outer morphology of Zr60All0Nil0Cu20 and Pd40Cu30Nil0P20 bulk cylindrical amorphous alloys.

Bulk Amorphous Alloys

385

Inner equiaxed zone

DF

~.i= 2~r( VATo

Ir 2 , G
)

O Amorphous phase zone I

C-

-Metastable single phase zone

Columnar dendritic zone

Outer equiaxed zone

'

V >>

61

!_O2)o

( unstable interface )

Figure 14.6. Schematicillustration showing the features and dominantparameters of the solidification structure of the arc-melted ingot.

diffusivity of the constituent elements at the interface and ~i is the distance of the atomic rearrangement required for the formation of the solid phase from the liquid. On the other hand, in the inner region, one cannot take account of the effect of reinforced heat flux from the copper hearth. It is therefore expected (Flemings 1974) that the direction of heat flux agrees with that of grain growth and the interface becomes unstable and irregular. However, we can always observe a smooth interface even in the inner region (Inoue and Zhang 1994). Considering that the wave length ()~i) of the interface can be expressed by ~.i = 2rc(DF/V AT0) 1/2, (1-" = o ' / A S f , AT0 is the temperature interval between liquidus and solidus temperatures) (Trivedi and Kurz1986), the inconsistency is presumably because the supercooled liquid has a high liquid/solid interface energy. This presumption is consistent with the interpretation (section 2) derived from the thermodynamic aspects. The Zr-based bulk amorphous alloys are expected to develop as an important engineering material. It is therefore important to search for additional elements that increase the GFA by suppressing the heterogeneous nucleation and growth of the crystalline phase. The Tg, Tx and ATx of the Zr65-xA1]0Cu15Ni10Mx (M = Ti or Hf) amorphous alloys were examined as a function of M content (Inoue et al. 1995a). No distinct changes in Tg, Tx and A Tx are seen for Hf, while the addition of Ti significantly decreases Tx and slightly increases Tg, leading to the decrease of A Tx from 107 K at 0% Ti to 38 K at 10% Ti. The influence on A Tx is significantly different between Ti and Hf. Similarly, the Tg, Tx and A Tx values as a function of additional M elements were examined for (Zro.65Alo.lCuo.x5Nio.l)lOo-xMx (M = V, Nb, Cr or Mo), Zr65AlloCu15-xNiloMx (M = Fe, Co, Pd or Ag) and Zr65AlloCulsNilo-xMx (M = Fe, Co, Pd or Ag) alloys. As these additional elements increase, Tx decreases gradually and Tg increases slightly, leading to a significant decrease in A Tx. This tendency is independent of the kind of the M ele-

386

A. Inoue

ments. No effective element on the increase in A Tx was found. Since the increase in Tg by the addition of M is slight, the significant decrease of A Tx for the M-containing alloys is concluded to be due to the decrease of Tx. Subsequently, the formation tendency of an amorphous phase in the arc-melted Zr65Al10Cu15Ni10 alloy ingots with tmax of about 8 mm was examined in comparison with the magnitude of A Tx for the melt-spun amorphous ribbons. Figure 14.7 shows the change in the ratio of amorphous area to the total cross sectional area with Ti or Hf content for the arc-melted Zr65-xAl10Cu15Ni10Mx (M = Ti or Hf) ingots. The area ratio of the amorphous phase remains almost constant (70%) for M = Hf. However, the ratio for the Ti-containing alloys increases from 67% for the quaternary alloy to about 90% for the 5 % Ti alloy and then decreases very rapidly with further increase in the Ti content. It is to be noticed that the area ratio of the amorphous phase increases significantly for the Ti-containing alloys with smaller A Tx values. This striking result suggests that the formation tendency of an amorphous phase in the arc-melted ingot is not directly related to A Tx. A similar increase in the area ratio of the amorphous phase to about 90% is recognized only for Nb in the (Zro.65Alo.lCuo.15Nio.1)lOo-xMx (M = V, Nb, Cr or Mo) alloys and Pd in the Zr65AlloCUls-xNiloMx and Zr65AlloCulsNilo-xMx (M = Fe, Co, Pd or Ag) alloys. We also examined the relation between the composition range in which an amorphous phase is formed in the arc-melted ingots and the temperature interval of the supercooled liquid region (A Tx). The amorphous phase was formed even at compositions with smaller ATx values only for the Ti-, Nb- and Pd-containing alloys. The minimum ATx value for the formation of the amorphous phase in the arc-melted ingots is 54 K for the Ticontaining alloy, 65 K for the Nb-containing alloy and 47 to 49 K for the Pd-containing alloy, indicating that the effectiveness of the additional elements is the greatest for Pd, followed by Ti and then Nb. It is again to be noticed that the Zr65AlloCu15-xNiloPdx (x _< 7.5 at%) and Zr65Al10CulsNi10-xPdx (x _< 5 at%) alloys with ATx below 50 K are made in the bulk amorphous form with a tmax of about 8 mm by arc melting. These results indicate clearly that the formation tendency of an amorphous phase by arc melting is not directly related to A Tx values, being different from the previous tendency (Inoue 1995a,b, 1996, 1997a,b,c) for Rc to decrease with increasing ATx. The inconsistency is thought to result from the incomplete melting of the arc-melted ingot in the region in contact with the copper mold. The temperature in the region in contact with the copper mold seems to lie around Tm because the temperature of the copper mold is usually below the Tm of Cu metal. The incomplete melting at the contact side induces the residual existence of crystalline phases that act as crystal nucleation sites in the subsequent solidification process. The arc-melted ingot always contains a number of crystalline nuclei in the region in contact with the copper mold. It is therefore thought that the GFA of the arc-melted ingot is not dominated by the difficulty of the generation of a crystalline nucleus, but is attributed to the difficulty of the growth of the crystalline nucleus. The evaluation of GFA by A Tx contains simultaneously both parameters of the nucleation and growth reactions. On the other hand, the evaluation of glass formation from the area ratio of amorphous phase in the arc-melted ingots contains only the parameter of the growth reaction. The

Bulk Amorphous Alloys

387

Zr65-xAI10Cu 15 Nil 0Mx

(9 r

t~ r CL >,, r

U'3 i

r

"6

0.5 ---O--

o

M=Ti M=Hf

~o <( t,,.

0

0

5 10 M content, X (at.%)

15

Figure 14.7. Changes in the area ratio of amorphous phase with M content for the arc-melted Zr65_xAl10Cu15Ni10Mx (M = Ti or Hf) ingots with a maximum thickness of 8 mm.

addition of Ti, Nb or Pd to the Zr-A1-Cu-Ni alloy is useful for suppressing the growth reaction of the crystalline nuclei.

14.5. S T R U C T U R A L R E L A X A T I O N A N D G L A S S T R A N S I T I O N

Figure 14.8 shows the change in the thermograms of a Zr65A17.5Cu27.5 amorphous alloy with Ta for an annealing period, ta, of 12 h (Inoue et al. 1992). On heating the as-quenched sample, the apparent specific heat (Cp,q) begins to decrease, indicative of a structural relaxation at 380 K, and shows a minimum value at 575 K. With a further increase in temperature, Cp,q, increases gradually up to 620 K and then rapidly in the glass transition range from 630 to 675 K, reaching 47.4 J/mol.K for the supercooled liquid at around 680 K. It then decreases rapidly due to crystallization at 731 K. The heating curve for the annealed sample shows a Cp,a(T) behavior which closely follows the Cp curve of the reference sample, C p,s, up to each Ta and then exhibits an excess endothermic relative to the reference sample before merging with that of the as-quenched sample in the supercooled liquid region above 642 K. The main features of Fig. 14.8 are summarized

A. Inoue

388

9

1

i

80

I

I

I

II 600K

0.67K/s

j'i

Zr65 AtT.sC,U 27.5

I

I

ta=12h

J

i J 6~

i

s

w

O

E

-b

O.

t9

40-

/ ./.~ii

500K 550K/.

Ta =400K ....

-.

.

/.-~.,,,~ ~,~-t

450K "\ \ ~

.,

Cp.\ \

X

"'-".~~ .

20 ,-

-i

\',..,/, X

\

.,,

//

Cp,q J

400

i

1

1

1

500 600 Terrc~rature TIK

. .

J

1

700

F i g u r e 14.8. Thermograms for an amorphous Zr65A17.5Cu27.5 alloy annealed for 12 h at temperatures from 400 to 620 K. The solid line represents the thermogram of the sample heated to 690 K.

Bulk Amorphous Alloys

389

as follows: (1) The sample annealed at Tg shows an excess endothermic specific heat beginning at Ta, implying that the Cp value in the temperature range above Ta is dependent on the thermal history and consists of configurational contributions as well as those arising from purely thermal vibrations. (2) The magnitude of the endothermic peak increases rapidly at Ta just below Tg. (3) The excess endothermic peak is recoverable, while the exothermic broad peak is irrecoverable and the Cp,a(T) curves for the annealed samples couple the recoverable endothermic and irrecoverable exothermic reactions. The excess endothermic peak has been thought (Chen 1981) to occur by the rearrangement from a relaxed atomic configuration caused by annealing to an atomic configuration which is more stable at temperatures above Ta. The excess endothermic reaction reflects the atomic configuration caused by annealing and hence we can obtain information on the structural relaxation during annealing by examining the change in the excess endothermic peak with Ta and ta. The temperature dependence of the difference in Cp between the annealed and the reference states, [ACp,endo = Cp,a(T) - Cp,s(T)], for La55AlzsNi20 and Zr65A17.5Cu27.5 alloys is shown in Fig. 14.9. With increasing Ta, A Cp,max for the two alloys initially increases gradually, followed by a rapid increase at temperatures slightly below Tg. The rapid increase in A Cp,maxis interpreted as corresponding to glass transition phenomenon. Similarly, the rapid decrease in A Cp,max above Tg is due to the achievement of internal equilibrium resulting from the very short relaxation times in the supercooled liquid region. Although the change in A Cp,max with Ta is similar to that for Zr-Cu and Zr-Ni amorphous alloys (Inoue et al. 1985a), it is different from the two-stage change which has been observed (Inoue et al. 1985b) for all amorphous alloys in the metal- metalloid systems containing more than two types of metallic elements. It has been suggested (Inoue et al. 1985b) that the appearance of the two-stage relaxation process is due to the difference in the relaxation time between the metal-metal pairs with weaker bonding and the metal-metalloid pairs with stronger bonding. Since no splitting of ACp,max into two stages is seen as a function of Ta for the La-A1-Ni and Zr-A1-Cu amorphous alloys, the relaxation times seem to be nearly the same for the La-A1, A1-Ni and La-Ni pairs and for the Zr-A1, A1-Cu and Zr-Cu pairs, which have large negative enthalpies of mixing. This result suggests that there is no appreciable difference in the attractive bonding between the atomic pairs and that the constituent atoms in these amorphous alloys are in an optimum bonding state. The optimum bonding state prevents easy atomic movement, even in the supercooled liquid, and suppresses the nucleation and growth of a crystalline phase, leading to the appearance of the extremely wide supercooled liquid region. The glass transition behavior was examined from the change in Cp(T) by the transition of an amorphous solid to a supercooled liquid and the temperature dependence of the specific heat in the amorphous solid and supercooled liquid. Figure 14.10 shows the thermograms of an amorphous Zr60AllsNi25 alloy (Inoue et al. 1990). The Cp increases gradually and begins to decrease, indicating an irreversible structural relaxation at 440 K. With a further increase in temperature, Cp reaches its minimum at 640 K, then increases rapidly in the glass transition range from 660 to 720 K and reaches 33.7 J/mol K for the supercooled liquid around 730 K. With further increase in temperature, the Cp of the supercooled liquid decreases gradually and then rapidly due to crystallization at 770 K. It

A. Inoue

390 t

"r

20[ / /

LassAI2sNi2o

' 10l j

30 ta =6h

t

425K

r

Zr65 AIT,5 Cu27.5

0.67K/s 620K ~~ 640K

ta=lh 575K 600K

~, 0

Ta =375K

400K

// I\

E -~ 50

d' <3 4o

J

J ta=96h

"I ~ 2o I

.10 0 375

425

450

475

Temperature TIK

ta=12h

'~ I

400K

400

.

500

525

o 400

/~~

550K

~ 500

~

!1 IIII

600

700

Temperature T/K

Figure 14.9. The differential specific heat, ACp(T), between the reference and annealed samples, for amorphous La55A125Ni20 and Zr65A17.5Cu27.5 alloys annealed for different periods at temperatures ranging from 375 to 620 K.

is seen in Fig. 14.10 that the transition of the amorphous solid to the supercooled liquid is accompanied by a large increase of ACp,sl of 6.25 J/mol K. The difference in Cp(T) between the as-quenched and the reheated states, [ACp(T)], manifests the irreversible structural relaxation which is presumed to arise from the annihilation of various kinds of quenched-in "defects" and the enhancement of the topological and chemical short-range ordering by the atomic rearrangement. The Cp,s curve of the reheated (control) sample is unaffected by thermal changes and consists of configurational contributions as well as those arising from purely thermal vibrations. Therefore, the vibrational specific heat, C p,v for the amorphous alloy is extrapolated from C p values in the low temperature region and is a linear function of temperature, viz.,

Cp,v-29.9+1.29x 10-2(T-340)

340
(14.6)

Similarly, the equilibrium specific heat, Cp,s, of the supercooled liquid, including the vibrational and configurational specific heat, can be expressed by Eq. (14.7) based on the data shown in Fig. 14.10, Cp,e

-

33.7 + 2.59 x 10-2(755 - T)

735 < T < 760

(14.7)

Bulk Amorphous Alloys I

50

I

391

1

I

I

0.67 K/s

Zr60 AI15 Ni25

.- 4O I

7 ~

m

.C.p q 20

J

'

t

Tg I

4OO

I

|

500 600 Temperature,

I

700

,

I

800

T~ K

Figure 14.10. The thermogram Cp,q(T) of an amorphous Zr60Al15Ni25 alloy in the as-quenched state. The solid line represents the thermogram Cp,s(T) of the sample heated to 750 K.

The ACp,sl for the Zr-A1-Ni amorphous alloys was examined as a function of composition. It was found that the A Cp,sl value is about 6.5 J/mol in the vicinity of Zr60A120Ni20 and tends to decrease with a deviation from the alloy composition, being similar to the compositional dependence of ATx and AHx. There is a tendency for ACp,sl to increase with increasing ATx and AHx. However, the ACp,sl values are considerably smaller than those (10 to 20 J/mol K) for Pt-Ni-P (Chen and Jackson 1978), Pd-Ni-P Chen, H.S. and Jackson, K.A. (1978), Mg-Ni-La (Inoue et al. 1989) and La-A1-Ni (Inoue et al. 1990) amorphous alloys. The small ACp,sl value may be related to the result that the glass transition appears at temperatures much higher than the Tg values of the Pt-, Pd-, Mg- and La-based amorphous alloys.

392

A. I n o u e

14.6. P H Y S I C A L P R O P E R T I E S

14.6.1 D e n s i t y ( I n o u e et al. 1 9 9 8 )

The anomalous X-ray scattering analyses have previously pointed out that when the three empirical rules for achievement of high GFA are satisfied, a higher dense random packed structure is formed in topological and chemical atomic configurations. Furthermore, the short-range atomic configurations are also different from those for the corresponding crystalline alloys. These data allow us to conclude that the supercooled liquid in the multicomponent systems has a new atomic configuration that has not been obtained for other metallic alloys. It is expected that the unique atomic configuration changes the density of the new bulk amorphous alloys, though the density for previously reported amorphous alloys with low GFA decreases by about 2% upon amorphization (Chen 1980, Masumoto 1982, Masumoto et al. 1976). Table 14.3 summarizes the densities in the as-cast, relaxed, and fully crystallized states for the Zr-A1-Cu, Zr-A1-Ni, Zr-A1-Ni-Cu, Zr-Ti-A1-Ni-Cu and Pd40Cu30Nil0P20 amorphous alloys (Inoue et al. 1998a). The densities increase steadily upon structural relaxation and crystallization. The increasing ratios of density defined by (Pc - Pa)/Pa are as small as 0.30 to 0.54% which are much smaller than those (about 2%) (Chen 1980, Masumoto 1982, Masumoto et al. 1976) for previously reported amorphous alloys. The small ratios indicate that the bulk amorphous alloys have a higher dense random packing state of the constituent atoms. This result is consistent with the tendency obtained from the structural analyses. In the higher packing density state, the atomic rearrangements of the constituent elements are suppressed, leading to the decrease in atomic mobility and the increase in viscosity. Furthermore, the frequency at which the atoms are trapped at the liquid/solid interface decreases in the higher dense random packing state, leading to the decrease in the nucleation frequency. These changes in the mobility and viscosity are interpreted to cause the increase in Tg/Tm and the decrease in the growth rate of a crystalline phase.

Table 14.3 Densitiesof Zr- and Pd-based bulk amorphous alloys in the as-cast, relaxed, and crystallized states

Alloy Zr60A110Cu30 Zr60A115Ni25 Zr55All 0Cu30Ni5 Pd40Cu30Ni10P20 Zr55Ti5A110Cu20Ni 10 Zr52.5Ti5A112.5Cu20Ni10

As-cast amorphous Pa (MG/m3) 6.72 6.36 6.82 9.27 6.62 6.52

Relaxed amorphous ,Or,a (Mg/m3) 6.83 9.28 -

Pc (Mg/m3)

Crystalline

Ap = (Pc -- Pa)/Pa %

6.74 6.38 6.85 9.31 6.64 6.55

0.30 0.31 0.44 0.54 0.30 0.45

Bulk Amorphous Alloys

393

14.6.2 Electrical resistivity (Haruyama et aL 1996) Figure 14.11 shows the DSC thermogram and the corresponding electrical resistance curve, R(T)Ro, around the supercooled liquid region for the Zr60AllsNizs alloy, where R0 represents the electrical resistance at 300 K. The temperature derivative curve is also plotted, because the onset of the glass transition is not obvious on the R(T) curve. Since Tg and Tx estimated from the R(T) measurement are in agreement with those of the DSC thermogram, the R(T) measurement was also used to investigate the thermal stability of amorphous alloys. The R value around room temperature decreased with increasing Ta. It has been reported (Okumura et al. 1991) that an amorphous (Nio.33Zr0.67)g5A115 alloy exhibits a negative temperature derivative of the resistivity in the temperature range from 3 to 300 K. Before the glass transition occurs, the R value increases temporarily, starting from 450 K, although the temperature derivative d(R(T)/Ro)/dT remains still negative. This is responsible for the onset of structural relaxation (Inoue et al. 1990). In the supercooled liquid region, d(R(T)/Ro)/dT is positive and is almost linear with temperature (see also Fig. 14.11). It is predicted that the electron transport properties are significantly changed after the glass transition. A similar R(T) curve has been obtained for the Zr60Al15Ni7.5Co2.5Cu15 alloy.

A

9

.A

i ...... J

i

>,

i

,i ~ I

~._..,,,,..,.,.,~..,......~..,x,..._...,~._..,..,.,...v,.,.,~,,,~.

E

.":~..

9 electrical resistance ......... differential DSC

E

"........

I

. . . . . . .

!I I./ / " ii

i/

i

IV

Tg

Tx jj Tx i ~176 ~176176 i ti

600

S4o

I

680

a

I

720

,

I

760

!1

800

,

Temperature, 77N Figure 14.11. DSC thermogram and electrical resistance curve along with its temperature derivative curve around the supercooled liquid region. The samples were prepared with Vs = 40 m/s in an argon atmosphere. All measurements were carried out in an argon atmosphere.

1

840

A. Inoue

394

14.6.3 Thermal expansion coefficient (Kimura et aL 1997) Figure 14.12 shows the thermal dilatation (AL/L) curve under a small tensile stress of 0.7 MPa for the melt-spun Pd-Cu-Ni-P amorphous ribbon, together with the data of Tg and Tx determined from the DSC curve at the same heating rate of 0.17 K/s. No appreciable change in A L / L is seen in the temperature range below 580 K, indicating that the alloy has an extremely low coefficient of thermal expansion (a) in the amorphous condition. With increasing temperature, the alloy begins to elongate at 580 K, followed by a rapid increase in elongation in the range of 600 to 615 K and then instantaneous stop of elongation around 620 K. In comparison with the Tg and Tx values marked with arrows in Fig. 14.12, it is seen that the elongation starts at a temperature 30 K higher than Tg and then stops at a temperature 30 K lower than Tx. Considering that the onset temperature of elongation agrees with the temperature where the supercooled liquid reaches the truly internal equilibrium state, the extremely large elongation is obtained only in the supercooled liquid state. Since the elongation stops at 620 K, due to the start of crystallization, the difference is attributed to the acceleration of crystallization reaction in the supercooled liquid caused by application of the small tensile load. The ot value is evaluated to be 8 x 10 -6 K -1 for the amorphous solid in the range below 500 K and 2.6 x 10 -2 K - 1 for the supercooled liquid in the range from 602 to 613 K. The change in the ot value in the amorphous solid and supercooled liquid with Cu content was also examined for the melt-spun Pd40Ni40-xCuxP20 amorphous ribbons. The ot for the amorphous solid

'

I

I

I

"Ix

Pd4o NiloCu30 P zo

1

Ribbon 0.7 MPa Tension

T Tg ,

500

Figure

I

1

1

550

600 Temp.,/K

650

_

14.12. Temperature dependence of relative elongation (AL/L) u n d e r a tensile stress of 0.7 MPa for a melt-spun Pd40Cu30Nil0P20 amorphous ribbon.

700

Bulk Amorphous Alloys

395

in the range from room temperature to 500 K is 1.0 x 10 -5 K -1 at 0% Cu, shows a minimum value of 0.8 x 10 -5 at 10% Cu and then increases gradually to 1.8 x 10 -5 K -1 with increasing Cu content to 40% Cu. The c~ is 1.3 x 10 -5 K -1 for the Pd40Cu30Nil0P20 alloy with the largest GFA. The lowest ot (8 x 10 -6 K -1) is the same order as that (1.2 x 10 -6 K -1) (Guillaume 1920) for a typical invar Fe-36.5 wt. % Ni alloy, indicating that the Pd-Cu-Ni-P amorphous alloys have very low coefficients of thermal expansion amongst metallic materials. On the other hand, the ot value for the supercooled liquid is as large as 3.4 x 10 - 2 K-1 at 0% Cu. 14.7. MECHANICAL PROPERTIES (INOUE E T A L . 1995)

The cylindrical amorphous alloys prepared by copper mold casting formed at different cooling rates in the inner and outer surface regions. The different cooling rates are expected to cause differences in the mechanical properties, though no appreciable difference is detected for Tg, Tx, and A Tx. We measured the Vickers hardness (Hv) as function of distance from the central point in the transverse cross section for the cylindrical Zr60A110Co3Ni9Cu18 amorphous alloys with diameters of 5 and 7 mm. The average Hv value was about 470 for the 5 mm cylinder and 475 for the 7 mm cylinder, in agreement with the value of 460 for the melt-spun ribbon. It is expected that the bulk amorphous Zr-A1-M alloys exhibit high tensile fracture strength (crf) comparable to that for the melt-spun amorphous ribbons. The test specimen of the amorphous Zr60All0Co3Ni9Cul8 alloy was made by mechanical working from the as-cast amorphous cylinder of 5 mm ~ x 50 mm. The gauge dimension has a diameter of 2 mm and a length of 10 mm as shown in Fig. 14.13(a). The tensile stress-strain curves obtained for the two test specimens are shown in Fig. 14.13(b). It is seen that the alloy deforms elastically, followed by yielding, slight plastic deformation accompanying a slight work hardening, distinct plastic flow accompanying serration and then final fracture. The feature of the stress-strain curve without significant work hardening and plastic elongation is in agreement with that for conventional amorphous metallic alloys with good bending ductility. It is to be noticed that the distinct plastic flow accompanying serration takes place before final fracture. The observation of the serrated flow phenomenon during tensile deformation has hitherto been reported only for noble metal-based amorphous alloys of Pd-based systems (Chen 1980). The occurrence of the serrated flow phenomenon implies that shear deformation typical of ductile amorphous alloys does not occur catastrophically over the whole cross section and is stopped at the deformation site. Momentary interruption of the shear deformation also implies the suppression of the final adiabatic fracture just after the shear deformation, indicating that the bulk amorphous alloy has a highly ductile nature comparable to that for the noble-metal base amorphous alloys. The Young's modulus (E), yield stress defined by the deviation from the proportional relation, elastic elongation, plastic elongation and O'f are 97 GPa, 1390 MPa, 1.6%, 0.4% and 1510 MPa, respectively. The values of ey,t(= ~ry/E), ~y,c(= 3Hv/E) and ~rf/Hv are 0.14, 0.15, and 3.2, respectively. The ratio of O'f to O v indicates that the amorphous solid deforms by an ideal elastic-plastic mode without work hardening. Thus, the present bulk amorphous alloy has good ductility and exhibits the tensile deformation and fracture behavior which agrees with that for an ideal elastic-plastic material.

396

A. Inoue

4

, ~.,3

2000

2

-I

0

ZrsoAIloCu18NigCo3

I

2

3

4

I

E

~ = 5.5 x 1 0 - " s - 1

13_ :E 1500 r,n r,/) t,., 4,-,

u')

1000

o ~

e -

~

500 0

oo5

_

Figure 14.13. (a) Outer surface appearance of an amorphous Zr60All0Co3Ni9CUl8 cylinder used for tensile testing, and (b) its tensile stress-strain curves at room temperature.

Figure 14.14 shows the tensile fracture surface appearance for the cast Zr60All0Co3Ni9Cul8 cylinder. Fracture occurs along the maximum shear plane which is inclined by about 45 degrees to the direction of tensile load and the fracture surface consists of a well-developed vein pattern. In comparison with the fracture surface appearance for the corresponding melt-spun amorphous ribbon, the diameter of the veins was about 1 to 2 mm for the cast cylinder, being about 10 times as large as that (about 0.2 mm) for the melt-spun ribbon and comparable to that for the ribbon subjected to creep deformation at elevated temperatures near Tg. The development of the vein pattern and the increase in the diameter of the veins suggest that the temperature during final adiabatic fracture increases because of the suppression of the final fracture resulting from the good ductility. Furthermore, the larger diameter of the veins also implies the increase in the thickness of the shear deformation region that induces an increase in the energy required for plastic deformation and final fracture.

397

Bulk Amorphous Alloys

(al

I

1

9

i'

l 2 Figure 14.14. (a) Tensile fracture behavior and (b) fracture surface appearance of an amorphous Zr60A110Co3 Ni9 Cu 18 cylinder.

The Charpy test specimen with a U-shape notch has the shape and dimension that agree with those for the JIS-No. 3 (2.5 mm in width) type. The fracture stress and displacement at the impact fracture and the impact fracture energy were measured with the Charpy testing machine (Inoue and Zhang 1996b). The initial angle of the arm bar was 143.5 degrees and the moving velocity at impact was 3.761 m/s. The load-displacement curve was stored in the computer and the impact fracture energy was simultaneously calculated from the load-displacement curve. Figure 14.15 shows the Charpy impact load-displacement curves for the Zr55A110Cu30Ni5 and Pd40Cu30Nil0P20 amorphous alloys. The maximum load is measured to be 4.3 kN for the Zr-based alloy and 4.4 kN for the Pd-based alloy and hence the maximum impact fracture stress is evaluated to be 1615 and 1630 MPa, respectively. The impact fracture stress is nearly the same as the static crf (about 1570 MPa (Inoue et al. 1995e) for the Zr-based alloy and 1670 MPa (Nishiyama 1997) for the Pd-based alloy obtained in the wide strain rate range of 10 -4 to 10-1 s-1 for the cast bulk amorphous alloys and the maximum displacement is as large as 0.44 to 0.48 mm. The impact fracture energy was 63 kJ/m 2 for the Zr-based alloy and 70 kJ/m 2 for the Pd-based alloy. In comparison with the Charpy fracture energies for other metallic crystalline alloys, the present impact fracture energies exceed largely those (12 to 36 kJ/m 2) (Ohtera 1996) for newly developed Al-based high-strength alloys that have already been used in some application fields. The comparison allows us to conclude that the present bulk amorphous alloys possess the impact fracture energies that are high enough to enable the application as a structural material. Three distinct peaks in the impact load-displacement curve have frequently been observed in the Charpy test and is presumed to result from the harmonic vibrations of the specimen at the impact fracture. It is of importance to clarify the impact fracture mode of Zr- and Pd-based amorphous alloys with the high impact fracture energies. The fracture surface consisted of a typical vein pattern in the region just near the U-shape notch and changed to an equiaxed dimplelike pattern which occupies most of the region in the fracture surface. Neither shell-like

A. Inoue

398

5001

2000

I

- Zr55Al~oCuaoNis

Z u.

4

~>

3

to 3

2

,

0

0

;

.

01.2

016 Displacement, d/ mm 0|.4

I

0.8

,

~="Pmax )/~

,

Stress = 1630MPa

300

200

,ooF/ i/

E

-1

Z

looo ~; ~

-6

'

I Pd4ONll 0Cu30P20 Ann L.

I^

o

-10

0

0.2

0.4

0.6

Displacement, d

0.8 /mm

Figure 14.15. Charpy impact fracture load displacement curves of the cast amorphous Zr55 A1!0Cu30 Ni5 and Pd40 Cu30 Ni !0 P20 specimens.

pattern nor featureless cleavage-like pattern typical for embrittled amorphous alloys was seen over the whole fracture surface. The feature of the fracture surface indicates that the cast bulk amorphous alloy has good ductility leading to severe deformation even in the impact fracture mode. Besides, secondary failure was observed and the change in the fracture surface appearance from equiaxed dimple-like pattern to vein pattern was recognized near the failure. The secondary failure appears to occur in close relation to casting defects. That is, the casting defect corresponds to the interface between the amorphous regions formed at two different stages. The amorphous phase region near the interface is subjected to subsequent heating. The heating decreases the ductility as compared with the as-cast amorphous phase without subsequent heating, leading to the generation of secondary failures. Although an interface between the formerly and latterly solidified amorphous alloys cannot be completely avoided in the casting process, a decrease in the number of the interfaces is expected to further raise the impact fracture energy. Based on the above-described data, in comparison with the corresponding crystalline alloys, bulk amorphous alloys have (1) higher of and lower E, (2) much larger elastic strain of about 2%, exceeding a yield limit of 0.2% for crystalline alloys, (3) much higher elastic energy up to yielding, (4) absence of distinct plastic elongation at room temperature due to inhomogeneous deformation mode, and (5) relatively high impact fracture energy of about 70 kJ/m 2. When the stress-strain curves of Zr-based bulk amorphous alloys are compared with those for Ti-based crystalline alloys, the elastic energy required up to yielding for the bulk amorphous alloy is larger by about 25 times. The relations between E and crf or Hv for the bulk amorphous alloys are shown in Fig. 14.16, where the data of conventional crystalline alloys are also plotted for comparison (Inoue 1997). When the E values of the bulk amorphous alloys are compared with those of the crystalline alloys with the same o'f, the moduli are smaller by about 60% than those for the crystalline alloys. The relation between E and O'f or By can be classified into two different groups. The distinct classification indicates that the features of mechanical properties of bulk amorphous alloys are different from those for the crystalline alloys. The distinct

399

Bulk Amorphous Alloys 5000

,

u

"

1500 ....

i /

'

'

i

.... l /

/

4500

4000 _

~aO3000

Fe-B base bulk glass D//

Fe-B base bulk glass l / /

/

/

i000

/

3500 /ll

-

2000 _ _

/

dPd-based

Young's Modulus (GPa) Figure

/

.~

-

/m /

Zr-based bulk glass j bulk glass / . 1500 / X Ti alloy,,,,,- :~ Stainless steel 0 ['~ lO00 ~" ~ Miz-based bulk ~lass ('~ " [ / - - " La-based bulk glass q ..... hiohv- : 5OO V / _ , , z . x tmrammm stremgth steel : 0 ~-~ .,~jMg allo~r (T7075-T6) I ' 0 i00 200 300 o

.,..~

-

Fe-P base bulk glass -

/

2500

'

Fe-P base bulk glass

/ j~ Zr-based bulk glass ,~_ Pd-based bulk glass r / X Ti alloy . ~ : ~ Stain'i'ess steel" ]" ~ M.-based Im~ ~lass 0 l- / Ua-b~d balk glass Super high/ .,~ ~ ~ stremgth steel , 0 L f ,,QMg allo~r (TT075-T6) I 100 200 300 0

500

Young's Modulus ( G P a )

Relation between Young's modulus (E) and tensile strength ( o f ) or Vickers hardness (He) for bulk amorphous and crystalline alloys.

14.16.

difference in the fundamental mechanical properties is very important for applications. Some engineering applications by the use of the difference are in progress because of the unique characteristics that cannot be obtained for crystalline alloys in each application field. 14.8. V I S C O E L A S T I C I T Y

Dynamic mechanical properties of storage, E ~, and loss, E", moduli were measured with a dynamic mechanical analyzer by a forced oscillation method. The complex modulus, E*, is given as the ratio of an oscillatory stress to the oscillatory strain. The real and imaginary parts of E* are defined as E ' and E" and described by Eq. (14.8) (Nielsen 1962) E* -- E' + i E " -- Eco2r2/(1 + co2z'2) -Jr- iEcor/(1 + co2.g2)

(14.8)

Here, E is the elastic modulus, co is the measurement frequency and z is the relaxation time. At coz = 1, E' and E" show an inflection point of a reduction curve and a loss peak, respectively. The temperature-dependent measurement was carried out every less than two degrees at heating rates from 0.0083 to 0.25 K/s at frequencies of 0.628, 6.28 and 62.8 rad/s. Figure 14.17 shows the temperature dependence of E ~ and E" of the La55AlzsNi2o amorphous alloy with heating rate (or) at 6.28 rad/s (Okumura et al. 1993). At 0.017 K/s, E ' starts to reduce at 466 K due to the lower relaxation system. The effect of the higher temperature relaxation system appears in U ( T ) at 481 K as a plateau which indicates the transition of the main relaxation system in the measurement from the lower temperature relaxation to the higher temperature relaxation, leading to a second rapid reduction

400

A. Inoue

of E t starting at 485 K with increasing temperature. From 503 K, E t increases due to crystallization. At a higher ot of 0.083 K/s, E ~ follows the same E t(T) curve of the first reduction at 0.017 K/s but E ~ shows more reduction superimposed on the first reduction curve at 0.017 K/s until 487 K. After the transition of the main relaxation system, E ~reduces again due to the higher temperature relaxation and rejoins the E t (T) curve at 0.017 K/s above 493 K. The appearance of the second relaxation system does not depend on the stress level but temperature suggesting the ot dependence, that is, shifting to higher temperature with increasing or. E ~shows more second reduction to a lower value as compared with the U(T) curve at 0.017 K/s because the Tx shifts to a higher temperature of 517 K at 0.083 K/s. With further increasing c~ to 0.13 K/s, the change in the E~(T) curve has the same tendencies. The first and the second reduction E ~(T) curves show the same and the superimposed curves at lower ot but the appearance shifts to the higher temperature of 493 K. The Tx also shifts to 521 K. In the temperature dependence of E" shown in Fig. 14.17(b), two relaxation peaks for the lower and the higher temperature relaxation and the crystallization peak are clearly observed at 473,490 and 511 K, respectively, at 0.017 K/s. The valley temperature of 479 K between the two relaxation peaks corresponds to the temperature for the appearance of the effect of the second relaxation in the E ~(T) curve as shown in Fig. 14.17(a). The peak temperatures, Tpl and Tp2, for the first and the second relaxation, respectively, show contrary results by the increase of or. Tpl slightly shifts to lower temperatures from 473 to 471 K and the height of the peak slightly decreases while Tp2 clearly shifts to higher temperatures from 490 to 501 K and the height rapidly decreases with increasing c~ from 0.017 to 0.13 K/s. At the higher c~, the appearance of the effect of the second relaxation shifts to higher temperature marking the shift of Tp2, and the overlap of the two relaxation peaks becomes less. Thus, the first peak distinctly appears with a slight shift to lower temperature and a slight decrease in height. Because the initial E" value for the second relaxation is lower, caused by the progress of the first peak, the height of the second peak decreases rapidly with increasing or. The E t~(T) curve merges with that of a lower ot at temperatures above Tp2 showing the same tendency as seen in the E ~(T) curve. The onset temperatures, To l and To2, for the first and the second reduction of U , are defined as the points of intersection of two tangent lines of the Et(T) curve. The ot dependence of To1 and To2 at three frequencies of 62.8, 6.28 and 0.628 rad/s is shown in Fig. 14.18(a). No different tendency is observed in the results of different frequencies. With increasing or, Tol negligibly increases to a saturated value while To2 distinctly increases. For example, in the case of 0.628 rad/s, Tol increases from 466 to 467 K and To2 increases from 481 to 493 K with increasing ot from 0.0083 to 0.083 K/s. Because the appearance of the second relaxation in E t (T) has the oe dependence, To2 shows the distinct ot dependence. The second relaxation starts to appear before the first relaxation shows a steep reduction in the E~(T) curve at lower ot (see Fig. 14.17(a)). Thus, the tangent line of the reduction side for To1 becomes sluggish, leading to the slight shift to a lower temperature for To1. At sufficiently higher ot as the first relaxation shows steep reduction, To1 shows very little ot dependence and saturated value. The shifting behavior of Tpl and Tp2 against ot at three frequencies is summarized in Fig. 14.18(b). As the two relaxation peaks become more distinct at the lower frequency

Bulk Amorphous Alloys

401

102 I

a~AI2~Ni..,,

"

.-.'T" o

o2

10'

o

,,,=o 9 9

9

o

9

~Po 9

A

;

9

o

o

,

90 . 1 3

440

, ,-

9 9

6,

90 . 0 8 3

10 -I

9

'= 9

o 0.017 K/S

9

9

o

6.28 rad/s

10v

9

*

~

l

*

,

480

I

J

520

i

TIK

I

6.28 rad/s

La.~sAl:.~Ni:, Tp2

0.017 K/S

o

9 0.083

10'|

.e#~ Tpl o~ ~,

!

~,

9 0.13

o

&Aee 6 6

ee %

6

9

...."

9

9 A

1r

9

9

9

A ee

lO-i

44O

~,

I

~0

,

I

5~

i

I

5~

T/K

Figure 14.17. (a) Change in the temperature dependence of storage modulus (E I) with heating rate at a frequency of 6.28 rad/s. (b) Change in the temperature dependence of loss modulus ( E ' ) with heating rate at a frequency of 6.28 rad/s.

402

A. Inoue

Las.~Ai2~Ni2~,

510

0 62.8rad/s

zx 6.~

To2

El 0.628

5OO

490

48O

470

460

L

I

!

10-2

10"

IlY'

Heating rate, a / K s-'

510

0 62.8r a d / s Z~6.28 Tr,2

5OO

490

480

Tpl

470

460

~ 10-2

i 10-~

_

,

10o

Heating rate, a / K s" F i g u r e 14.18. Heating rate dependence of Tol and To2 at three frequencies of 0.628, 6.28 and 62.8 rad/s. (b) Heating rate dependence of Tpl and Tp2 at three frequencies.

Bulk A m o r p h o u s Alloys

403

(Okumura et al. 1992), Tpl and Tp2 at the lower frequency shift more drastically than those at the higher frequency. The endothermic reaction for the glass transition in the DSC curve also shifts to higher temperature with increasing or. To correlate with the change in Ef'(T) with or, the apparent activation energies, Q, are evaluated by the Kissinger method (Kissinger 1957). The value of Q obtained from Tp2 at three frequencies varies slightly from 300 to 340 kJ/mol which agrees with the value of 340 kJ/mol from the DSC measurements. It should be mentioned that Q of about 340 kJ/mol represents the value not only for the relaxation but also the value including the effects of the phase separation and relaxation. The real activation energy for the higher temperature relaxation has been obtained as 550 kJ/mol (Okumura et al. 1992a). From this result, it is clear that the shift of the DSC curve is mainly affected by the higher temperature relaxation. Making a comparison of the ratio of E" to E', tan ~ (= E " / E ' ) , between the two relaxations, tan 6 for the higher temperature relaxation exceeds 10 ~ which is one order of magnitude higher than that for the lower temperature relaxation (Okumura et al. 1991). It suggests that the higher temperature relaxation occurs more drastically than the lower temperature relaxation. The appearance of the ot dependence of Tg is mainly caused by the phase separation as evidenced by two E ~ relaxation peaks. In order to examine the ot dependence of the normal glass transition without the effect of the phase separation, poly-methyl methacrylate, PMMA (Iwayanagi and Hideshima 1953) is used. As amorphous alloys with one relaxation peak, to our knowledge, do not show stable supercooled liquid region, the crystallization peak of E ~ overlaps with the relaxation peak. The ot dependence is hardly observed in mechanical measurements considering the margin of errors. Almost all amorphous alloys with a wide supercooled liquid region exceeding 40 K in the DSC curve at 0.67 K/s, such as Zr60Al15Ni25, Mg65Cu25Y10 and Pd48Ni32P20 amorphous alloys, have two relaxation systems suggesting the phase separation in the glass transition region (Okumura et al. 1991/92). These alloys show distinct ot dependence of Tg like the La55AlzsNi20 alloy. On the other hand, the amorphous alloys with one relaxation system and smaller supercooled liquid region, such as Zrv0Cu30, La55A145 and Pd80Si20, show little ot dependence of Tg. A similar two-stage change in the dynamic mechanical properties has been noted for the Zr60AllsNi25 amorphous alloy (Okumura et al. 1992b). 14.9. S O F T M A G N E T I C P R O P E R T I E S

Based on the three empirical rules for the achievement of high GFA, ferromagnetism at room temperature has been found in Fe-(A1, Ga)-(R C, B, Si) (Inoue and Gook 1995, Inoue et al. 1995b, Inoue et al. 1996b), Co-Cr-(A1, Ga)-(P, B, Si) (Inoue and Katsuya 1996), Fe-(Co, Ni)-(Zr, Nb, Ta)-B (Inoue et al. 1997e, Inoue et al. 1998b, Inoue et al. in press) and Co-Fe-(Zr, Nb)-B (Inoue et al. 1998b) systems. This section describes the soft magnetic properties of the Fe- and Co-based bulk amorphous alloys.

14.9.1 Formation and soft magnetic properties of bulk amorphous alloys The Fe-based amorphous alloys with large A Tx (>60 K) are expected to have a high GFA which enables the production of bulk amorphous alloys with diameters larger than

404

A. I n o u e

1.5

I

as-Q.

1.0

Y_. tO

,,,,,,,,

N

.

_

_ FeTaAI5Ga2P11CsB4

(lmm ~)

............. Fe72AIsGa2P10CeB4Si1 (2mrn~)

0.5 o.o

,,,,,,m

,i,,,,i

~E

-0.5

e"~"§ -1.0

-1.5

"o .~

-16oo

u

9 . . . . . . . . .

I

-eoo

....

o

Applied Field, H(A/m)

I

800

,

16oo

Figure 14.19. Hysteresis B-H loops of the cast amorphous Fe73AI5Ga2PllC5B4 and Fe72A15Ga2PIoC6B4Si 1 cylinders.

1 mm by the casting processes. The cast Fe72A15Ga2PllC6B4 (Inoue et al. 1995b) and FevzA15GazP10C6B4Sil (Inoue et al. 1996b, Inoue et al. 1997e) amorphous cylinders with diameters of 1 to 2 mm had smooth surface and metallic luster. Besides, good castability of these Fe-based alloys is indicated by noting that the outer shape of the cast samples mirrors the inner cavity of the copper mold. Even after an appropriate etching treatment, no appreciable contrast corresponding to a crystalline phase is seen over the whole transverse cross section of the cast Fe72A15Ga2P10C6B4Si] cylinders with diameters of 1 and 2 mm. In addition, neither cavities nor shrinkage holes are observed, indicating that the amorphous alloy has a good castability owing to the lack of the discontinuous change in the specific volume-temperature relation which is ordinarily recognized for crystalline alloys obtained from the liquid. The cast FevzA15Ga2P]]C6B4 amorphous cylinders with diameters of 0.5 to 1.5 mm (Inoue et al. 1995b) show the sequential changes of the glass transition, supercooled liquid and then single-stage exothermic peak which agree with those of the corresponding melt-spun amorphous ribbon. No distinct changes in the thermal stability and crystallization mode of the supercooled liquid are seen between the cast amorphous cylinders and the melt-spun amorphous ribbon.

Bulk Amorphous Alloys

405

The hysteresis B - H loop of the cast Fe72A15Ga2P11C6B4 amorphous cylinder with a diameter of 1 mm was examined in the as-cast and annealed (723 K, 600 s) states (Inoue et al. 1995b, Inoue et al. 1996b). The Bs, He and Br/Bs of the annealed sample are 1.07 T, 5.1 A/m and 0.37, respectively, indicating that the cast amorphous cylinder has good soft magnetic properties. Besides, the ~e at 1 kHz for the annealed amorphous cylinder also shows a high value of 7000. The soft magnetic properties were further improved for the 1 at% Si-containing alloy (Inoue et al. 1996b, Inoue et al. 1997b). Figure 14.19 shows the hysteresis B - H loop of the cast FevzA15Ga2P10C6B4Si1 amorphous cylinder with a diameter of 2 mm, together with the data of the cast Fev3A15Ga2PllCsB4 amorphous cylinder with a diameter of 1 mm. The Bs, He and Br/Bs for the Si-containing cylinder are 1.14 T, 0.5 A/m and 0.38, respectively. In comparison with those (Inoue et al. 1995b) for the Fe-A1-Ga-P-C-B cylinder, Hc decreases and B~/Bs increases for the Si-containing amorphous cylinder. From the thermomagnetic data, it is confirmed that the cast Fev3A15GazP11CsB4 amorphous cylinder has a Curie temperature (Tc) of 600 K which is lower by 185 K than the Tx. The Tc value agrees with that determined from the endothermic peak on the DSC curve. Thus, the cast Fe-based amorphous cylinders with diameters up to 2 mm exhibit good soft magnetic properties of 1.1 T for Bs, 2 to 6 A/m for He and 7000 for #e at 1 kHz. When these soft magnetic properties are compared with those for the corresponding melt-spun amorphous ribbons, no distinct changes in Bs and He are seen. However, the #e value is degraded for the cast cylinder presumably because of the increase in the influence of the demagnetization resulting from the significant change in the sample morphology. Figure 14.20 shows the outer morphology of the bulk Fe61CovZrl0Mo5W2B15 cylinders with diameters of 3 and 5 mm (Inoue et al. 1997f). These samples also have smooth surface and metallic luster and no contrast of a crystalline phase is seen over the outer surface. The X-ray diffraction patterns showed a main halo peak with a wave vector Kp(= 47r sin 0/)Q around 29.6 nm -1 and no crystalline peak is observed even for the 5 mmO sample. Besides, the optical micrographs of the cross section of the two samples also revealed a featureless contrast in a sample etched with HE Considering that the bulk cylinder of 7 mm diameter consists of an amorphous phase in an outer surface region with a thickness of about 2 mm and of amorphous and crystalline phases in the inner region, the

Figure 14.20. Outermorphologyof cast amorphous Fe61Co7Zrl0Mo5W2B15 cylinders with diameters of 3 and 5 mm.

406

A. Inoue

tmax for the Fe61Co7Zrl0Mo5W2B15 alloy is determined to be about 6 mm. The tmax is 3 times larger than the largest value (2 mm for FevzA15Ga2P10C6B4Sil) (Inoue et al. 1996b, Inoue et al. 1997b) for Fe-based amorphous alloys reported to date. Figure 14.21 shows the DSC curves of the bulk amorphous Fe61Co7NivZrsNbzB15, Fe56CovNi7ZrgTazB20, Fe60CosZr8NbzMosWzB 15 and Fe60Co8Zrl0MosWzB 15 cylinders with diameters of 1 to 3 mm (Inoue et al. 19970. These amorphous alloys exhibit the sequential transition of glass transition, supecooled liquid and crystallization. The ATx is as large as 55 to 88 K and crystallization occurs by a single exothermic reaction. The crystallites were identified to consist of ot-Fe, Fe2Zr, Fe3B, MoB and W2B phases for the Fe60Co8Zrl0Mo5WzB15 sample heated to a temperature just above the exothermic peak. Thus, the crystallization is due to the simultaneous precipitation of the five crystalline phases. This crystallization mode is in agreement with that (Inoue 1995a, Inoue 1995b, Inoue 1996, Inoue 1997a, Inoue 1997b) for other bulk amorphous alloys. The largest ATx is 88 K for Fe56Co7Ni7Zr8Ta2B20, being larger than the largest values (57 to 67 K) for Fe-(A1, Ga)-(P, C, B, Si) (Inoue and Gook 1995, Inoue et al. 1997b) and nonferrous Pd- and Pt-based amorphous alloys (Chen 1980, Chen 1976, Chen 1978). The Tm was 1420 K for Fe56Co7Ni7Zr10B20 and 1416 K for Fe61Co7Zrl0Mo5W2B15 and the Tg/Tm was evaluated to be 0.60 for the former alloy and 0.63 for the latter alloy. Considering that Tg / Tm is 0.54 for Fe80P12B4Si4 (Inoue and Park 1996) and 0.57 for Fe73A15Ga2P11C4B4Sil (Inoue et al. 1997b), the present Tg / Tm values are believed to be the highest among all Fe-based amorphous alloys. Table 14.4 summarizes the tmax, Tg, A Tx, Tg/Tm, compressive fracture strength (Crc,f), Hv, Bs, He, # e at 1 kHz and )~s for the new amorphous Fe-(Co, Ni)-(Zr, Nb, Ta)-B (Inoue et al. 1997e, Inoue et al. in press) and Fe-Co-Zr-(Mo, W)-B alloys (Inoue et al. 1997f). These amorphous alloys exhibit good soft magnetic properties in an annealed (800 K, 300 s) state, i.e., high Bs of 0.74 to 0.96 T, low He of 1.1 to 3.2 A/m, high #e of 12000 to 25000 and low )~s of (10-14) • 10 -6. The He and #e are superior to those for conventional Fe-Si-B amorphous ribbons (Kikuchi et al. 1975, Zhang et al. 1991), presumably because of the lower )~s. Furthermore, the Fe60CosZrl0MosWzB 15 bulk amorphous alloy has high crc,f and Hv of 3800 MPa and 1360, respectively, which exceed largely those (crf = 1000-2000 MPa, Hv = 300-800) for high-carbon high-alloy tool steels and 25 wt%Ni maraging steel. No bulk alloys with high strength above 3000 MPa for Crc,f and 1000 for Hv have been obtained for any Fe-based alloys including amorphous and crystalline phases. Besides, no weight loss is detected after immersion for 3.6 ks at 298 K in aqua regia. Thus, the Fe-based bulk amorphous alloys possess simultaneously high GFA, high strength, high corrosion resistance and good soft magnetic properties which cannot be obtained for other amorphous and crystalline alloys. 14.10. V I S C O U S FLOW AND M I C R O - F O R M A B I L I T Y OF S U P R C O O L E D LIQUID

In addition to the low Rc and large tmax, bulk amorphous alloys have a wide supercooled liquid region before crystallization (Inoue 1995a, Inoue 1995b, Inoue 1996, Inoue 1997a, Inoue 1997b). The ATx reaches as large as 127 K (Zhang et al. 1991) for the Zr-A1-NiCu alloys and 98 K (Inoue et al. 1996a) for the Pd-Cu-Ni-P alloys. In the supercooled liquid region with such large A Tx values, the viscosity decreases significantly and large

Bulk Amorphous Alloys

9

r

9

i

'

i

....

9.....

407

=

9

i

0 . 6 7 K / s

Fe6oCo~ZrloMo5W2B15 d=3mm~ .

.

.

.

r-

.

FesoCosZraNb2MosW2B15 d--3mm+

-

,

.

.

.

.

.

.

.

.

.

F

.

r-"

._o

v

(--

E FessCo7Ni7ZrsTa2B20 d=lmm t,,.._

(!) r 0 X

Ill

t

FeslCo7Ni7ZreNb2Bls d=lmm~, .

.

.

.

.

Tg

.

I

9

600 Figure 14.21. Fe56C~

1

9

1

9

I

700 800 900 Temperature, T/K

,

1

1000

DSC curves of cast amorphous Fe61Co7Ni7Zr8Nb2B15, Fe60C~ and Fe60Co8Zr10Mo5W2B15 cylinders.

P 0 to

Table 14.4 Thermal stability, mechanical strength, and magnetic properties of the cast bulk amorphous Fe-(Al, GaHP, C, B, Si), Fe-(Co, Ni)-(Zr, Nb, T a t B , and Fe-CeZr-(Mo, W)-B alloys Thermal stability h a x

Tg

Vickers

Compressive

Hc

Pe

As

(T)

Wm)

(1 kHz)

(lop6)

-

0.96 0.75 0.85 0.74

2.0 1.1 3.2 2.6

19100 25000 12000 12000

10 13 14 14

3800

-

-

-

14

AT, (K)

Tgl

hardness

strength

(K)

Tm

HU

q c . f (MW -

FeyjC07Ni7ZrloB20 FeyjCo7Ni7ZrgNb~B20 Fe61Co7Ni7ZrgNb2B15 Fe56Co7Ni7Zr8TazB20

2 2 2 2

814 828 808 827

73 86 50 88

0.60

-

1370 1370 1340 1360

Fe60CogZr10MogW2B15

6

898

64

0.63

1360

-

Soft magnetic properties

-

B,

? 3 0 fi

m

Bulk Amorphous Alloys

409

viscous flow is obtained (Inoue and Zhang 1997). By utilizing the low viscosity and large viscous flow, the bulk amorphous alloys can be deformed to various complicated shapes while retaining the good mechanical properties (Inoue et al. 1997d, Inoue et al. 1996c, Kawamura et al. 1996, Nishiyama and Inoue to be submitted). This section intends to describe the micro-formability in the supercooled liquid region for the La-A1-Ni, Zr-A1Ni-Cu and Pd-Cu-Ni-P bulk amorphous alloys and to introduce some practical examples of micro-formed bulk amorphous alloys. 14.10.1 Feature of phase transition of bulk amorphous alloys

The phase transition of bulk amorphous alloys heated at a scanning rate of 1 K/s is different from that of conventional amorphous alloys which require high cooling rates above 104 K/s. The bulk amorphous alloys show the sequential transition of amorphous --+ glass transition --+ supercooled liquid region ~ crystallization upon continuous heating, while neither glass transition nor supercooled liquid region is observed for the conventional amorphous alloys. Thus the bulk amorphous alloys have a wide supercooled liquid region before crystallization. The appearance of the wide supercooled liquid region also implies that the supercooled liquid has a high resistance against crystallization. The high thermal stability is the origin for the achievement of the high GFA. The reason for the high stability of the supercooled liquid in the new systems has been described (Inoue 1995a, Inoue 1995b, Inoue 1996, Inoue 1997a, Inoue 1997b) in Section 14.2 on the basis of the formation of a higher dense random packed structure in the multicomponent alloys which satisfy the three empirical rules. 14.10.2 Deformation behavior of supercooled liquid

The Tg and ATx values of the bulk amorphous La55A125Nie0, Zr65A110Ni10Cu15 and Pd40Cu30Nil0P20 alloys measured at a heating rate of 0.67 K/s were respectively 480 and 70 K for the La-based alloy, 650 and 85 K for the Zr-based alloy and 575 and 95 K for the Pd-based alloy. Fig. 14.22 shows the relation between the maximum flow stress (O'max) and the strain rate (~) in the supercooled liquid for the La55A125Ni20 and Zr65A110Ni10Cu15 amorphous alloys (Inoue and Saotome 1993, Inoue et al. 1996c, Kawamura et al. 1996). The O'max increases with increasing ~ and the slope becomes significant in the higher temperature range. The slope corresponds to the strain-rate sensitivity exponent, m. As seen in the figure, the m-value is measured to be approximately 1.0 in the supercooled liquid region, indicating that the supercooled liquid has an ideal Newtonian flow, i.e., ideal superplasticity. Fig. 14.22 also shows that the deformation in the supercooled liquid takes place in a homogeneous mode over the whole strain rate range from 1.5 • 10 -4 to 7 • 10 -1 s -1 (Kawamura et al. 1996). This result also indicates that the high strain-rate superplasticity is obtained in the supercooled liquid. The amorphous alloy exhibits an inhomogeneous deformation mode in the temperature range well below Tg and the flow stress decreases slightly with increasing strain rate. Thus, the deformation behavior differs significantly between the amorphous solid and supercooled liquid. Figure 14.23 shows the change in the viscosity of the supercooled liquid for the Pd40Cu30Nil0Pe0 amorphous alloy during isothermal heating (Hata et al. to be submit-

410

A. Inoue 1000

~ . .I

~aK

100

~

_ f" P" / ~,f~ f " " " .,,eo ,..=..r'" .'~" /'653K o~/" ,/ik'" /

B~s~ w

=E

~

r 473K

fl'" 9

100

/

,

r

//:,/,"//= ,/

4

~ lO rr

~

~,"

/ IV/713K

10

~

10 "4

1 -3

LassAI2oNi25 , 10-1 100

' 10-2

ZresAiloNiloCU15 ,,

1o::4

I

10-3

....

I

I

10-2

10-1

Strain rate, ~ / s'l

Strain rate, ~ I s "1

Figure 14.22. Relation between the maximum true flow stress (O'max) and strain rate (~) in the supercooled liquid for the amorphous La55A125Ni20 and Zr65Al10Ni10Cu15 alloys.

16 14 12 o

>,

4-,

~

o o .(a >

......... qTm = 103 Pa.s q Tm -- 102 Pa. s

Pd40Nil oCu30P20

- ' - - qTm = 101 Pa. s

9TMA result

10

-

Fulcher t y p e e q u a t i o n

8

....... q

6

0 0.75

( Calculated ) _

9.34 x l O 3 e x p

0,,,~0

i 0.80

I 0.85

I 0.90

4135

T - 447

0.95

1.00

Figure 14.23. Temperature dependence of viscosity in the supercooled liquid for the amorphous Pd40Cu30Nil0P20 alloy. The solid circles represent the experimental data and the three lines denote the calculated values on the assumption of different viscosities of the liquid at Tin.

____

100

Bulk Amorphous Alloys

"'4 "'\N ,

~

~

)i!j !!

..1 :,'

"

........

"-';

,

'fj,,

"/;..,',

,.. ,'//

'...i~(:"/ // /'

:7.",~-~'. """ '~'

.,/" "

Figure 14.24. Wire- and gear-shaped amorphous La55A125Ni20 samples prepared by tensile stretching and die forging treatments, respectively, in the supercooled liquid region.

411

412

A. Inoue

. ....

~.~,lb;mm

Figure 14.25. Fine precision amorphous alloy mirrors prepared by die forging an amorphous Zr60All0Nil0CU20 alloy in the supercooled liquid region.

ted). Considering that the flow deformation becomes extremely easy in the viscosity range below 108 Pa.s, the viscous flow deformation can occur during the long time period of 102 to 104 s in the temperature range between 600 and 650 K. The times are long enough to form various complicated shapes because these amorphous alloys have the high strain-rate superplasticity.

14.10.3

Micro-formingof supercooled liquid

By utilizing the ideal superplasticity which can be achieved over the wide strain-rate range in the supercooled liquid region, various micro-forming treatments have been conducted on the La-, Zr- and Pd-based amorphous alloys. Figure 14.24 shows the outer shapes of the La55A125Ni20 amorphous wire obtained by tensile deformation at an initial strain rate of about 1 x 10 -2 s-1 and 500 K (Zhang et al. 1991 b) and the La-based amorphous gear with an outer diameter of 1 mm prepared by die forging into the gear-shape silicon die for 103 s at 500 K (Inoue and Saotome 1993). We have also obtained (Nishiyama and Inoue to be submitted) small amorphous gears by a die extrusion process in the supercooled liquid region for amorphous Zr65Al10Ni10Cu15 alloy with much higher Tg as compared with that for the La-based amorphous alloys. In the relation between extrusion pressure and extrusion velocity at 711 K in the supercooled liquid region for the Zr-A1-Ni-Cu amorphous alloy, there is a linear relation and the m-value is evaluated to be 1.0, indicating that the supercooled liquid also exhibits ideal superplastic behavior even in the die extrusion treatment. Figure 14.25 shows the outer morphology of the precision optical mirrors prepared by die forging the Zr60All0Nil0Cu20 amorphous alloy for 600 s at 673 K. In the as-forged state, the forged sample exhibits a smooth outer surface and good metallic luster. The outer surface keeps good smoothness on a nanometer scale from the data obtained by a surface roughness indicator. The three-stage amorphous gears with outer diameters of 5, 6 and 7 mm were also prepared by press die-forging of the Pd40Cu30Nil0P20 amorphous alloy for 600 s at 630 K (Hata et al. to be submitted). All micro-formed alloys retain the amorphous phase and no distinct difference in the thermal stability and mechanical properties is seen between the as-cast and micro-formed samples (Inoue et al. 1997d, Inoue and Zhang 1997, Kawamura et al. 1996, Nishiyama and Inoue to be submitted). The combination of

B u l k A m o r p h o u s Alloys

413

high GFA, good micro-formability and good mechanical properties has already enabled some practical uses of these bulk amorphous alloys as various micro-forged materials such as fine machinery parts and fine precision optical parts.

14.11. A P P L I C A T I O N S A N D F U T U R E P R O S P E C T S

As described above, the discovery of a number of new alloys with a high stability of the supercooled liquid against crystallization has led to a drastic decrease by 6 to 8 orders in the critical cooling rate for glass formation and enabled the production of bulk amorphous alloys with thicknesses ranging from several to about 80 millimeters by various casting processes. This drastic increase of the stability of the supercooled liquid also implies the advent of new basic science and materials. Table 14.5 summarizes the fundamental characteristics and application fields of the bulk amorphous alloys produced by conventional casting processes. It is confirmed that the bulk amorphous alloys exhibit various characteristics such as high mechanical strength, high elastic energy, high impact fracture energy, high wear resistance, high corrosion resistance, good soft magnetic properties, high frequency permeability, fine and precise viscous deformability, good castability and high consolidation tendency into bulk forms. By utilizing these characteristics inherent to the bulk amorphous alloys, the new alloys are expected to open up new application fields in the near future. Table 14.5 Fundamentalcharacteristics and application fields of bulk amorphous alloys Fundamental characteristic

Application field

High strength High hardness High impact fracture energy High elastic energy High corrosion resistance High wear resistance High viscous flowability Good soft magnetism High magnetostricfion

Machinery structural materials Optical precision material Tool materials Cutting materials Corrosion resistant materials Ornamental materials Composite materials Writing appliance materials Sporting goods materials Soft magnetic materials High magnetostrictive materials

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Bulk Amorphous Alloys

415

Johnson, W. L. (1996a) Current Opinion in Solid State and Material Science, 1, 383. Johnson, W. L. (1996b) Mater. Sci. Forum, 225-227, 35. Kawamura, Y., Shibata, T., Inoue, A., and Masumoto, T. (1996) Appl. Phys. Lett., 69, 1208. Kikuchi, M., Fujimori, H., Obi, T. and Masumoto, T. (1975) Jpn. J. Appl. Phys., 14, 1077. Kim, Y. H., Inoue, A. and Masumoto, T. (1990) Mater. Trans. JIM, 31, 747. Kimura, H. M., Inoue, A., Nishiyama, N., Sasamori, K., Haruyama, O. and Masumoto, T. (1997) Sci. Rep. Res. Inst. Tohoku Univ., A43, 101. Kissinger, H. E. (1957) Analyt. Chem., 29, 1702. Klement, W., Willens, R. H. and Duwez, E (1960) Nature, 187, 869. Kneller E. E and Hawig, R. (1991) IEEE Trans. Mag., 27, 3588. Kui, H. W. and Tumbull, D. (1985) Appl. Phys. Lett., 47, 796. Kui, H. W., Greer, A. L. and Tumbull, D. (1984) Appl. Phys. Lett., 45, 615. Kurz, W. and Fisher, D. J. (1989) Fundamentals of Solidification. (Trans Tech Publications, Zeurich). Liebermann, H. H. (1993) (ed.) Rapidly Solidified Alloys. (Marcel Dekker, Inc., New York). Luborsky, E E. (1983) (ed.) Amorphous Metallic Alloys. (Butterworths, London). Masumoto, T. (1982) (ed.) Materials Science of Amorphous Metals. (Ohmu Pub., Tokyo). Masumoto, T., Kimura, H. M., Inoue, A. and Waseda, Y. (1976) Mater. Sci. Eng., 23, 141. Matsubara, E., Tamura, T., Waseda, Y., Inoue, A., Kohinata,M. and Masumoto, T. (1990) Mater. Trans. JIM, 31, 228. Matsubara, E., Tamura, T., Waseda, Y., Inoue, A., Zhang, T. and Masumoto, T. (1992a) Mater. Trans. JIM, 33, 873. Matsubara, E., Tamura, T., Waseda, Y., Zhang, T., Inoue, A. and Masumoto, T. (1992b) J. Non-Cryst. Solids, 150, 380. Nielsen, L. E. (1962) Mechanical Properties of Polymers, Chap. 7, Reinhold, New York. Nishiyama, N. (1997) Doctor Thesis, Tohoku Univ., Sendai, Japan. Nishiyama, N. and Inoue, A. (1996) Mater. Trans. JIM, 37, 1531. Nishiyama, N. and Inoue, A. (1997) Mater. Trans. JIM, 38, 464. Nishiyama, N. and Inoue, A., Mater. Trans. JIM, to be submitted. Ohtera, K. (1996) Doctoral Thesis, Tohoku Univ., Sendai, Japan. Okumura, H., Inoue, A. and Masumoto, T. (1991) Mater. Trans. JIM, 32, 593. Okumura, H., Inoue, A. and Masumoto, T. (1991/92) Sci. Rep. Res. Inst. Tohoku Univ., A36, 239. Okumura, H., Chen, H. S., Inoue, A. and Masumoto, T. (1992a) J. Non-Cryst. Solids, 142, 165. Okumura, H., Inoue, A. and Masumoto, T. (1992b) Jpn. J. Appl. Phys., 31, 3403. Okumura, H., Inoue, A. and Masumoto, T. (1993) Acta Metall. Mater., 41, 915. Peker, A. and Johnson, W. L. (1993) Appl. Phys. Lett., 63, 2342. Suzuki, K., Makino, A., Kataoka, N., Inoue, A. and Masumoto, T. (1991) Mater. Trans. JIM, 32, 93. Trivedi, R. and Kurz, W. (1986) Acta Met., 34, 1663. Turnbull, D. (1956) Solid State Phys., 3, 225. Uhlmann, D. R. (1969) Materials Science Research, Vol. 4, (Plenum Press, New York, NY), p. 172. Uhlmann, D. R. (1972) J. Non-Cryst. Solids, 7, 337. Yokoyama, A., Komiyama, H., Inoue, H., Masumoto, T. and Kimura, H. M. (1981) J. Catalysis, 68, 355. Yoshizawa, Y., Oguma, S. and Yamauchi, K. (1988) J. Appl. Phys., 64, 6044. Zhang, T., Inoue, A. and Masumoto, T. (1991a) Mater. Trans. JIM, 32, 1005. Zhang, T., Tsai, A. P., Inoue, A. and Masumoto, T. (1991b) Boundary, 7(9), 39.

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Author Index

Ananthapadmanabhan, E V., 121 Bhaduri, S. B., 289 Bhaduri, S., 289 Cahn, R. W., 1 Colligon, J. S., 225 Dollet, A., 257 Groza, J. R., 347 Inoue, A., 375 Ishihara, K. N., 5

Jones, H., 23 Kobayashi, K. E, 89 Koch, C. C., 313 Lavernia, E. J., 153 Li, B., 153 Liu, B. X., 197 Suryanarayana, C., 49, 313 Teyssandier, E, 257 Venkatramani, N., 121

417

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Subject Index

ABS s e e arc bond sputtering process (ABS) AC plasma torches, 126 Acetone, as process control agent, 52 Adhesion, ion-induced, 239 AFM s e e atomic force microscopy (AFM) Age-hardening, early studies, 1 Alloy phase formation, 41 ion mixing, 200-3, 217-18 Alloy phases free energy, 7-8 new, 41 nucleation, 15-17 Alloys disordering, 64 eutectics, 29 microstructure selection maps, 97-8 precipitation-hardened, 76 solid solution-strengthened, 76 solidification microstructure, 97 spray-formed, 169-70 supercorroding, 80 s e e a l s o amorphous alloys Alonso's method, 200-1 Alumina coatings, 133 plasma-sprayed, 131 scanning electron micrograph, 143-4 transmission electron micrograph, 143, 145 ultra fine powder, thermal plasma synthesis, 142-7 or-Alumina, synthesis, 306 y-Alumina, formation, 144-7 Aluminum powder plasma oxidation, 142-3 solutes, solid solubility

extensions, 31 aluminum-base alloys, applications, 79-80 aluminum copper alloys, mechanical properties, 185-6 aluminum iron alloys, solidification microstructure selection map, 35 Aluminum-iron-vanadium-silicon alloys, tensile properties, 42 Aluminum-nickel-manganese alloys, HallPetch plots, 41 Aluminum oxide ultrafine, 140 X-ray diffraction, 143 Aluminum-zinc alloys, mechanical properties, 187 AM s e e amorphous materials (AM) Amorphization conditions, 207 kinetics, 98 and milling temperature, 69 solid-state, 67-71, 199, 218 Amorphous alloys, 58, 199, 319 corrosion resistance, 98 formation, 70, 314-16 via ion mixing, 198, 207-8 metal-metal, formation, 197 metal-metalloid, formation, 197 metallic, 332 ternary, 380 s e e a l s o bulk amorphous alloys Amorphous materials (AM), structure, 268 Amorphous phase formation and atomic size, 69-70 conditions, 69

419

420

Subject Index

ion mixing, 200, 205-12 kinetics, 96-7 laser quenching, 98-9 in mechanical alloying, 49, 50, 61, 68, 81 in mechanical milling, 68 predicted composition ranges, 207-8 processes, 67-8 rapid solidification, 376 Amorphous phases enthalpy of formation, 201 free energy of, 6-7 heats of crystallization, 9 processing, 141 Anti-site disorder, 64 APS see atmospheric plasma spraying (APS) Arc bond sputtering process (ABS), 236-7 Arc discharge, 147 Arc melted ingots, 385-7 Arc plating system, 232 Arrhenius parameters, 272, 354 Atmospheric plasma spraying (APS), 134, 135 Atmospheric plasmas, 122 Atomic clusters, nanoscale, 313 Atomic collision cascade, use of term, 199 Atomic force microscopy (AFM), applications, 278 Atomic size and amorphous phase formation, 69-70 and solid solubility extensions, 32 Atomization electro-hydrodynamic, 147 melt, 23, 41 processes, 24-5 see also spray atomization Atomization gas, in spray forming, 160-3 Atomizers circular, 158 in spray forming, 158-63 Attritor mills, 53 Auto-ignition synthesis, of nanocrystalline oxides, 304-6

Automobile emissions, 340 Ball milling, nanocrystalline materials, 67 Ball mills ball-to-powder weight ratios, 57 in mechanical alloying, 49, 53-4, 71 Bearing steels, 41 Borides, reaction temperatures, 294 Brazing foils, 43 Bulk amorphous alloys, 375-413 alloy systems, 376 applications, 413 characteristics, 413 continuous cooling transformation, 3803 density, 392 electrical resistivity, 393 endothermic reactions, 403 glass-forming ability, 376-83 factors, 376-80 glass transition, 387-91 historical background, 375-6 hysteresis, 404-5 mirrors, 412 phase transition, 409 preparation, 383-7 properties, 408 magnetic, 403-6 mechanical, 395-9 physical, 392-5 specific heats, 387-91 structural relaxation, 387-91 supercooled liquids deformation behavior, 409-12 micro-formability, 406-13 synthesis, 403-6 tensile fracture behavior, 396-7 thermal expansion coefficient, 394-5 thermal stability, 384 thermograms, 388-91,393-4 viscoelasticity, 399-403 viscous flow, 406-13 Young's modulus, 399

Subject Index

c-BN s e e cubic boron nitride (c-BN) Calculating the phase diagram s e e CALPHAD CALPHAD, 7, 8 applications, 12 Carbides, reaction temperatures, 293 Carrier concentration, active, 103, 104 Catalysts heterogeneous, 340 nanostructured materials, 340-1 rapidly solidified materials, 41 CCPS s e e counter current plasma synthesis

(ccps) CCT s e e continuous cooling transformation

(CCT) Ceracon processing, 57, 326 Ceramics coatings, 133 nanocrystalline, 330 non-oxide, production, 140-1 processing, thermal plasmas, 140 CFUBMS s e e closed field unbalanced magnetron sputtering (CFUBMS) Charpy test, 397-8 Chemical vapor deposition (CVD), 225, 257-82 applications, 257, 316 approaches solid-gas thermodynamic equilibrium, 267-8 thermodynamics, 263-4 crystal growth, 270-1,278-81 morphology, 279-81 diamond deposition, 141-2 early studies, 2 fluid flows non-reactive, 260-1 reactive, 261-7 gas-phase mechanisms, 264 gas-phase transport, 259-67 growth rates, 263 kinetics, 264, 271-3 theoretical method, 273 layers, structure, 268-70

421

mechanisms, 258-9 modified, 265 nucleation enhancement, 276-7 epitaxial, 275-6 heterogeneous, 273-8 multi-component, 276 solid particle, 265-7 studies, 277-8 theories, 274-5 operating parameters, 261 plasma-assisted, 248 reactors, 257-8,260-1 research, 282 selective deposition, 277 solid phase formation, 267-81 supersaturation, 270-1 surface mechanisms, 271-3 thermodynamics, 263-4 Chemical vapor infiltration (CVI), 257 CHEMKIN/SENKIN package, 264 CIP s e e cold isostatic pressing (CIP) Closed field unbalanced magnetron sputtering (CFUBMS), 236, 237 Co-injection, and spray forming, 163-6 Coatings adhesion, 134 anti-reflective, 243,244 ceramic, 133 corrosion protective, 245 decorative, 225 hard, 246-8 ion-assisted, 249 nanocomposites, 248,249 nanostructured materials, 339 optical, 243-5 plasma spraying, 131,132-4 thermal barrier, 133, 134, 339 wear-resistant, 132, 339 Cobalt/tungsten carbide powder, production, 318 Cold isostatic pressing (CIP), 358 Cold-wall reactors, 257,258 Cold welding, 52, 55, 79

422

Subject Index

Combustion synthesis (CS), 63,289-307 advantages, 290 apparent activation energy, 295 applications, 290 auto ignition, 294 "azide" process, 294, 297-9 conversion efficiency, 297 development, 2, 289-90 enthalpy of formation, 291 field-activated, 295-6 gas phase reactions, 299 green density, 303 ignition chemical oven, 293 explosion mode, 293 kinetics, 294-5 mechanisms, 290 non-equilibrium effects, 307 propagating reactions, 296 reaction temperatures, 291-4, 299, 3012 thermodynamics, 291-4 use of term, 304-5 Combustion wave propagation, 304 Containerless processing, 41 Continuous cooling transformation (CCT) bulk amorphous alloys, 380-3 diagrams, laser processing, 96, 112, 113 Continuous wave lasers, 89, 94-5 cooling rates, 94 Copper-zirconium alloys, mechanical properties, 188 Corrosion protective coatings, physical vapor deposition, 243,245 Corrosive atmospheres, materials for, 77, 79 Counter current plasma synthesis (CCPS), 141 Crucibles, in spray forming, 158 Cryomilling, processes, 51 Crystal faces classification, 278-9 growth rates, 279-81 Crystal growth

chemical vapor deposition, 270-1,27881 laser processing, 96 mechanisms, 278 metastable phases, 12, 18 undercooling, 18 unstable, 280-1 Crystal-to-amorphous transformation, 205 Crystalline materials, synthesis, 268 Crystalline morphology, 279, 280 Crystalline phase formation laser quenching, 99-100 in mechanical alloying, 49 Crystallization explosive, 99 kinetics, 96 Cubic boron nitride (c-BN), coatings, 247 Cutting tools, nanostructured materials, 337 CVD see chemical vapor deposition (CVD) CVI see chemical vapor infiltration (CVI) CW lasers see continuous wave lasers DC plasma arc process, 139 DC plasma torches, 124-6 Dental amalgams, 43 Deposition dual ion beam, 233,234 electron beam vapor, 316 energy-assisted, 231,233,248 ion-assisted, 228,232, 238,249 ion beam, 233-5 ion beam induced chemical, 233,234 ionized cluster-beam, 226-7,233-5 laser-enhanced chemical vapor, 257 modified chemical vapor, 265 organometallic chemical vapor, 257 plasma-activated chemical vapor, 257 plasma-enhanced chemical vapor, 257, 268 rapid thermal processing chemical vapor, 257 thin film, 92, 230

Subject

chemical vapor deposition (CVD); electrodeposition; physical vapor deposition (PVD); pulsed laser deposition (PLD); spray deposition; vapor deposition diamondlike carbon (DLC), coatings, 246 Diamonds free energy, 5 plasma deposition, 141-2 Raman spectra, 142, 143 vapor-phase synthesis, 141 DIB s e e dual ion beam deposition (DIB) Differential calorimetry, 26 Differential scanning calorimetry (DSC), 9, 324 Direct reduced iron (DRI), 138 DISPAL, production, 50 Dispersion-strengthened materials, 162 research, 189 Dispersoids distribution, 71 in mechanical alloying, 50-1, 52 Displacement per atom (dpa), determination, 199-200 Displacement reactions, 49, 50 in mechanical alloying, 71-3 DMC s e e dynamic magnetic compaction (DMC) dpa s e e displacement per atom (dpa) DRAM s e e dynamic random access memory (DRAM) DRI s e e direct reduced iron (DRI) DSC s e e differential scanning calorimetry (DSC) Dual ion beam deposition (DIB), 233,234 Dynamic magnetic compaction (DMC), 368 Dynamic random access memory (DRAM), 100 see also

Index

423

Electromotive force (emf), measurement, 9, 10 Electron beam lithography, 341 Electron beam vapor deposition, applications, 316 Electron gun evaporation, 204 Electroplating, nanocrystalline materials synthesis, 328 Electrosleeve steam generator, 337 Elements, free energy, 5-7 Emf s e e electromotive force (emf) Emulsification, processes, 25-6 Enthalpy of mixing and glass formation, 30 mechanical alloying, 9-11 Miedema's model, 9 Epitaxy, 268 Equal-channel-angular technique, 318 Equilibrium phases, 29 Ethanol, as process control agent, 52 Evaporation, in physical vapor deposition, 228-9 EXCIMER lasers, 98-100, 103, 115 Explosive forming, 57

FACS s e e field-activated combustion synthesis (FACS) FAST s e e field-activated sintering technique (FAST) Fatigue crack growth, resistance, 79 Fe-6.3wt%Si, grain size, 32, 36 Feedstocks, rapid solidification in, 41 Ferromagnetic materials, nanoscale, 333, 334 Ferrous alloys s e e iron alloys Field-activated combustion synthesis (FACS), 294, 295-6 system, 298 Field-activated sintering technique (FAST), 363,364, 366-8 Field ion microscopy (FIM), 272 Electric arc spray process, 167-8 Filamentary structures, 313 Electro-discharge compaction, 57,326 Electrodeposition, nanostructured materi- Filtered arc system, 233 FIM s e e field ion microscopy (FIM) als preparation, 317, 335

424

Subject Index

FINEMET materials, 319, 333,335,340 Finite differences, and laser beam processing, 92, 94 Fluxing, in rapid solidification processes, 26 Frank-Read sources, 331 Frank-Van der Merwe mechanism, 275, 276 Free energy of alloy phases, 7-8 of elements, 5-7 interfacial, 219 of metastable crystalline phases, 216-18 see a l s o Gibbs free energy Free energy-concentration diagrams, 16 Free energy hierarchy, 7, 9, 13 Fuel cells, 43 Functional nanostructures, 313 Furnace racks, 79 Fusion welding, 91 heat-affected zones, 112

maps, 30 nominal, 211-12 Glass-forming composition range (GFR), and ion mixing, 198,205,206-7, 21112 Glass transition temperature, 7 Glassy phases, devitrification, 322 Glow discharge plasmas, 121 Glycine-nitrate process, 304 Golf club heads, 43 Grain growth abnormal, 324 inhibitors, 324 isochronal, 324 isothermal, 323 kinetics, 325-6 kinetics, 323 nanostructured materials, 323-6 Grain refinement, 64 Grain size and enthalpy of mixing, 9-11 and mechanical alloying, 71, 76 and rapid solidification, 32-3 Grinding balls, 53 Grinding media in mechanical alloying, 51, 74 velocity, 53

Gallium free energy hierarchy, 9 metastable phases, 15 Gas turbines combustion chambers, 79 vanes, 77, 78 Gasless combustion see combustion Hall-Petch equation, 183,240, 247, 329 synthesis (CS) Hall-Petch plots, 37, 41 GFR see glass-forming composition range Hall-Petch strengthening, nanocrystalline (GFR) materials, 67 Gibbs free energy, 6 Hard coatings, physical vapor deposition, determination, 201 246-8 Glass formation, 2, 32, 94 Hard facings, 43 bulk, 29 Hazardous wastes, treatment, 129, 141 critical cooling rates, 29 HE Mill, 54 and enthalpy of mixing, 30 Heat exchangers, components, 79 via combustion synthesis, 307 Heats of mixing Glass-forming ability (GFA) excess, 19 bulk amorphous alloys, 377-83 negative, 12, 69 and heats of formation, 206-8 Heptane, as process control agent, 52 intrinsic, 211 Heteroepitaxy, 269 and ion mixing, 198,205,206 Hexane, as process control agent, 52

Subject

Hierarchy of stability, 6 High-velocity oxyfuel thermal spray process, 167-8 HIP s e e hot isostatic pressing (HIP) Homoepitaxy, 268 Hot corrosion, resistance, 76, 79 Hot extrusion, 56-7, 326 Hot isostatic pressing (HIP), 56-7, 58, 290, 326, 363-6 Hot pressing, 290, 363,364 Hot-wall reactors, 257, 258 Hume-Rothery electron compounds, 62 Hydrogen, absorption, 337, 341 Hydrogen embrittlement, 104-5 Hydrogen storage materials nanostructured materials, 340-1 via mechanical alloying, 80 Hydrostatic extrusion, 326 Hysteresis effect, 237 IAD s e e ion-assisted deposition (IAD) IBC s e e ion beam induced chemical deposition (IBC) IBD s e e ion beam deposition (IBD) ICB s e e ionized cluster-beam deposition (ICB) ID s e e isothermal desorption (ID) Ideal solutions, phase diagrams, 16-17 IM s e e ion mixing (IM) Inert gas condensation, applications, 322 Ingot casting, and spray forming compared, 153-4 Integrated circuit pads, 115 Interaction parameter, 11 Intergranular corrosion, 335 Intergranular fractures, 185 Intermetallic matrix composites research, 189 spray-formed, 171, 172 synthesis, 168 Intermetallics disordering, 64-6 superlattice reflections, 65, 66 X-ray studies, 65

Index

425

ordered, 63 synthesis, 62-4, 67 Interstitials, redistribution, 64 Ion-assisted deposition (lAD), 228, 232, 238,249 Ion beam deposition (IBD), methods, 2335 Ion beam induced chemical deposition (IBC), 233,234 Ion bombardment, 228, 231 Ion implantation, 103, 153 applications, 197 early studies, 2 thermal plasmas, 122 Ion irradiation, 205 Ion mixing (IM), 197-220 alloy phase formation, 200-3 amorphous phase formation, 205-12 applications, 197 early studies, 2 experimentation, 203-5 glass-forming ability, 206-12 physics, 199-200 Rutherford backscattering spectrometry, 205 Ion-plating, 225,226, 227, 248 properties, 231 system, 232 Ionized cluster-beam deposition (ICB), 226-7,233,234-5 Iron alloys applications, 79 laser beam welding, 106-8 laser surface-alloying, 100-1 Iron-molybdenum multilayered films, backscattering spectra, 206 Iron oxide, ultrafine, 140 Iron-titanium carbide coatings, 136 Isothermal desorption (ID), 272 Joule effect, 92 Kiln furniture, 79 Kissinger method, 324, 403

426

Subject Index

Knock-out atom, 199 Lamellar structures see layered structures Lance pipes, coke injection, 79 Laser ablation, 92 Laser alloying, 98 Laser annealing, 90, 91, 102-5 applications, 103-4 Laser beam joining, 106-15 Laser beam processing, 94 and finite differences, 92, 94 Laser beam soldering, 90, 91-2 applications, 112 thermal cycles, 114 Laser beam welding, 90, 91 crack-flee, 112, 114 crack frequency, 113 ferrous alloys, 106-8 non-ferrous alloys, 108-12 weld fusion zones, 112 Laser-enhanced chemical vapor deposition (LECVD), 257 Laser glazing, 91 processes, 93-4 Laser-induced thermal desorption (LITD), 272 Laser irradiation applications, 99 for solution annealing, 104 Laser melting heat transfer, 92 processes, 92 and rapid solidification, 92, 95 Laser processing, 89-115, 122 classification, 89-92 continuous-cooling-transformation diagrams, 96, 112, 113 cooling rates, 96 crystal growth rate, 96 high-power, 98 partition coefficient, 97 solidification modes, 97 time-temperature-transformation diagrams, 96

Laser quenching amorphous phase formation, 98-9 crystalline phase formation, 99-100 nanostructured materials, 99 Laser scanning, in surface solutionizing, 105 Laser-solid interactions, 92-3 Laser surface alloying, 90, 91,100-2 Laser surface cladding, 90, 91, 100-2 Laser surface melting, 89-90, 99 Laser surface modification, 153 Laser surface quenching, 89-90 Laser surface solutionizing, 105 Laser surface treatment, 99 applications, 100-1 Lasers absorption phenomena, 92-3 absorption ratio, 92-3 early studies, 2 EXCIMER, 98-100, 103, 115 principles, 89 Q-switched sources, 89 spot size, 89 see a l s o continuous wave lasers; pulsed lasers Lattice damage, 103 Lattice parameter changes, 59 deviation, 12, 19 in disordered intermetallics, 65 of supersaturated phase, 11-12 Layer-by-layer mechanism, 275,277 Layered structures, 313 Lead shot, manufacture, 24 LECVD see laser-enhanced chemical vapor deposition (LECVD) Levitation (containerless) melting, in rapid solidification processes, 26 Ligaments, formation, 154 LIPS see liquid injection plasma synthesis (LIPS) Liquid crystal displays (LCDs), laser annealing, 103

Subject Index

Liquid injection plasma synthesis (LIPS), 128 applications, 141 Liquid melt quenching (LMQ), 197 LITD s e e laser-induced thermal desorption (LITD) LMQ s e e liquid melt quenching (LMQ) Low-friction materials, 132-3 Low-pressure plasmas, 121, 122

427

byproduct removal, 71-3 crystal defects in, 55-6 defect density, 59 deformations, 64 development, 5, 58 diffusion distances, 59 displacement reactions, 71-3 early studies, 2 enthalpy of mixing, 9-11 and grain size, 71, 76 MA s e e mechanical alloying (MA) grinding media, 51, 74 Magnesium-base alloys, applications, 80 interstitial contamination, 63 Magnetic materials mechanisms, 55-6 hard, 43,333-5 metastable phases, 51 permanent, 333-2, 335, 340 mills, 52-5 modeling, 75-6 remanence-enhanced, 73, 334, 335 nanocrystalline phases, 61 soft, 333,340 s e e a l s o ferromagnetic materials nomenclature, 50-1 Magnetic recording materials, 340 phase formation, 51 Magnetocaloric effect, 335 powders Magnetron sputtering, 230, 248-9 consolidation, 56-8 in physical vapor deposition, 235-8 contamination, 74-5 Magnetron systems, unbalanced, 227, 236, handling precautions, 81 particle size, 51 243 Marine corrosion, resistance, 79 processes, 51-5 Martensite phases, deformation-induced, and rapid solidification compared, 58 raw materials, 51-2 63 Martensitic transformation research, 2 early studies, 1 variables, 75 reverse, 215 Mechanical crystallization, 69 Material transfer, in mechanical alloying, Mechanical grinding (MG), 50 Mechanical milling (MM), 64, 66 50 Maximum possible amorphization range amorphous phase formation, 50, 68 (MPAR), 198,205 Medium-pressure plasmas, 122 MBRS s e e molecular beam reactive scatter- Melt atomization, 23, 41 ing (MBRS) Melt extraction, processes, 27 MCVD s e e modified chemical vapor depo- Melt extrusion s e e melt spinning sition (MCVD) Melt overflow, processes, 27 Mechanical alloying (MA), 49-81, 153, Melt spinning, 23, 41, 188, 340 340 cooling rates, 30 and enthalpy of mixing, 30 amorphous phase formation, 49, 50, 61, free-jet chill-block, 26, 27 68, 80 applications, industrial, 76-80 processes, 26-7 attributes, 49-50

428

Subject Index

Metal hydrides mechanical alloying, 51 applications, 341 nucleation, 13-15 production, 80 occurrence, 5 Metal matrix composites (MMCs), 43,293 physical vapor deposition, 240 synthesis, 163 processing, 141 Metal-metal systems, ion mixing, 209-10 spray forming, 178 Metallic glasses, 1,199 Metastable powders bulk modulus, 34 consolidation, 349-55 creep, 36 pressure, 363-6 ductility, 34 kinetics, 351 elasticity, 34 sintering fatigue limits, 36 mechanisms, 352-3 fracture, 34 scaling laws, 353-4 liquid melt quenching, 211 stability, 350 production, 141, 197,207, 212 thermodynamics, 350-1 properties, 34-7 Methanol, as process control agent, 52 electrical, 36 MG see mechanical grinding (MG) magnetic, 36 Microjoining, in laser beam processing, shear band offsets, 39 112-15 ultra-high strength, 26 Micromachining, and laser processing, 115 zirconium-based, 26 Microstructure formation, via rapid solidiMetalorganic vapor phase epitaxy process fication, 28-34 (MOVPE), 257 Microstructure selection maps (MSMs), alMetastabilities, types of, 348 loys, 97-8 Metastable alloy phases, 205 Miedema's theory, 9, 66, 198, 199, 200-1, Metastable alloys 218,219 formation, via ion mixing, 198, 199 Milling nanoscale composite, 247 ball, 67 Metastable crystalline alloys, formation, fracture rate-welding equilibrium, 55 199,213-18 high-energy, 55 Metastable crystalline phases, 58 intensity, 68, 74, 75 classification, 213-15 interrupted, 63 f.c.c.-I, 213-14 media, 74 f.c.c.-II, 215 reaction, 50 free energy, 216-18 research, 2 h.c.p.-I, 213 see also mechanical milling (MM) h.c.p.-II, 215 Milling atmosphere, 75 in intermetallic synthesis, 62 control, 50 Metastable phase formation and powder contamination, 74 kinetics, 5, 12-18 Milling temperature, and amorphization, 69 thermodynamics, 5-12 Milling/mixing see mechanical alloying Metastable phases (MA) crystal growth rate, 12, 18 Mills enthalpy of transformation, 9 commercial, 53-4 free energy, 9-11

Subject Index

high-energy, 53 low-energy, 53 new designs, 54-5 powder contamination, 74 types of, 52-5 s e e a l s o ball mills Mirrors, bulk amorphous alloys, 412 MM s e e mechanical milling (MM) MMCs s e e metal matrix composites (MMCs) Modified chemical vapor deposition (MCVD), 265 Molecular beam epitaxy (MBE), in physical vapor deposition, 229 Molecular beam reactive scattering (MBRS), 272 Moving point source model, 94 MOVPE s e e metalorganic vapor phase epitaxy process (MOVPE) MPAR s e e maximum possible amorphization range (MPAR) MSMs s e e microstructure selection maps (MSMs) Multilayered films interfacial free energy, 201-3 ion mixing, 197, 198 MX alloys s e e metastable alloys MX phases s e e metastable crystalline phases Nanocomposite refrigerants, 335 Nanocomposites, 338-9 applications, 339 coatings, 248,249 deformation behavior, 339 densification, 361 sintering, 361 synthesis, 306 Nanocrystalline alloys grain boundary mobility, 325 synthesis, 316 Nanocrystalline materials applications, 337-41 catalysts, 340-1

429

magnetic, 340 structural, 337-9 chemical reactivity, 335 corrosion behavior, 335 deformation mechanisms, 332 diffusivity, 328 dissolution rate, 335 early studies, 313 grain boundaries diffusion, 327 mobility, 325 segregation, 324 structure, 321-3 thickness, 320 grain growth, 323-6 grain size, 320-1 grain structure, 321 Hall-Petch strengthening, 67 hydrogen sorption, 337 magnetic, 319 mechanical twinning, 332 M6ssbauer spectra, 321 plastic deformation, 67 properties, 66, 327-37 chemical, 335-7 magnetic, 333-5 mechanical, 328-32, 337 shear banding, 332 structure, X-ray diffraction, 321 superplasticity, 332, 339 synthesis, 66-7, 104, 129, 306 gas condensation chambers, 316, 317 methods, 315, 319 spray conversion processing method, 318 thermal stability, 323, 361-2 three-dimensional, 317 volume fraction of atoms, 320 Nanocrystalline oxides, auto-ignition synthesis, 304-6 Nanocrystalline phases, in mechanical alloying, 61 Nanocrystalline powders, 319 Nanocrystalline spinel, synthesis, 306

430

Subject Index

Nanocrystallites, structure, 341 Nanodevices, 313 Nanofibers, production, 319 Nanoglasses, 314 Nanoparticles, 313,340 Nanoperm alloys, 333,335,340 Nanopowders ceramic, 355,356 compaction cold, 356-8 dry, 357-8 warm, 358-9 densification, 369 metal, 355,356 problems, 350 sintering, 328, 351,359-61 stability, 350-1 Nanoquasicrystals, 314 Nanorods, production, 319 Nanostructured materials, 313-42 applications, 337-41 catalytic, 340-1 electronic, 341 magnetic, 70, 340 structural, 337-9 atomic structure, 321-2 bulk, 313 classification, 313-14 composite, 70 corrosion behavior, 335 crystal lattices, atomic structure, 321 deformation mechanisms, 331-2 plastic, 318 diffusion, 327-8 diffusional creep, 330-1 ductility, 329, 331 early studies, 2 elasticity, 328 enthalpy, 9-12 ferromagnetic, 335 field of study, 313 grain boundaries, 67 atomic structure, 321-2

grain growth, 323-6 inhibition, 324-5 kinetics, 325-6 grain size, 318, 331 hardness, 329 junctions, 322-3 laser quenching, 99 microstructure, 320-1 optical, 341 powder consolidation, 326-7 properties, 327-37 catalytic, 335-7 chemical, 335-7 magnetic, 333-5 mechanical, 328-32 research, 341-2 retention, 57-8 sinterability, 327-8 stability, 323-6 strain rate, 331 strength, 329 structure, 319-23 superplasticity, 331 synthesis, 70, 314-19 via chemical reactions, 318 via devitrification, 319 via electrodeposition, 317 via inert gas condensation, 314-16 via mechanical attrition, 318 via rapid solidification, 316 toughness, 329 see also nanocrystalline materials New crystalline phases, 29 Ni-5wt%A1, grain size, 32, 36 Nickel, temperature-time curve, 29 Nickel aluminide, coatings, 133 Nickel-base alloys applications, 77-8 glass forming ability map, 30 properties, 77 Nickel-molybdenum system, free energy diagram, 217 Nickel-niobium system free energy diagram, 204

Subject Index

glass-forming composition range, 206-7 Nitrides, reaction temperatures, 299 Nitrogen, source of, 298 Non-equilibrium phases crystalline, 30 formation, 146, 147 nanoscale dispersions, 32 rapid solidification, 29-32 synthesis, 58-73 Non-equilibrium plasmas, 121 Non-equilibrium processing applications, 1 development, 5 early studies, 2 mechanical alloying, 49 in solid solubility extensions, 59 and Vegard's law, 11 Non-ferrous alloys laser beam welding, 108-12 laser surface-alloying, 101-2 Nopcowax-22DSP, as process control agent, 52 Nucleation frequencies alloy phases, 15 homogeneous, 13 metastable phases, 13-14 volume, 13 Octane, as process control agent, 52 ODS materials s e e oxide dispersionstrengthened (ODS) materials OMCVD s e e organometallic chemical vapor deposition (OMCVD) Optical coatings, physical vapor deposition, 243-5 Organometallic chemical vapor deposition (OMCVD), 257 Orowan looping mechanism, 184 Ostwald's step rule, 5, 13, 15 Oxalic acid, as process control agent, 52 Oxidation plasma, 142-3 resistance, 76, 78, 79, 101,133,301 surface, 341

431

Oxide dispersion-strengthened (ODS) materials, 49, 51, 59, 160-2 advantages, 77 applications, 76-80 sintering, 57 Oxide powders, production, 140 Oxides, single-phase, 306 Packaging, high-density, 112 PACVD s e e plasma-activated chemical vapor deposition (PACVD) Palladium-copper-nickel-phosphorus bulk amorphous alloys, 380-2 Palladium-copper-silicon alloys, constituent selection diagram, 35 PAS s e e plasma-activated sintering (PAS) Pattern transfer processes, 341 PCAs s e e process control agents (PCAs) PCM s e e polycrystalline materials (PCM) Pechini process, 304 PECVD s e e plasma-enhanced chemical vapor deposition (PECVD) Permanent magnets, materials, 41 Phase diagrams equilibrium, 7 metastable, 8 Phase separation, effects, 7 Physical vapor deposition (PVD), 225-49, 257 adhesion, 239-40 applications, 225,243-8 coatings, 243-8 corrosion protective coatings, 245 corrosion resistance, 243,245 deposition ion-assisted, 231-5 methods, 228-38 development, 225-7 early studies, 2 energy-assisted, 245 evaporation, 228-9 film morphology, 238-9 future trends, 248-9 hard coatings, 246-8

432

Subject Index

magnetron sputtering, 235-8 metastable phases, 240 microstructure, 240-3 molecular beam epitaxy, 229 novel phases, 243 nucleation, 239-40 optical coatings, 243-5 sputtering, 229-30 Planar diode sputtering, 230, 231 Planar flow-casting, 23 processes, 27 Planetary ball mills, 53 Plasma-activated chemical vapor deposition (PACVD), 257 Plasma-activated sintering (PAS), 57, 326, 366 Plasma cutting, 129 Plasma decomposition, 137-8 Plasma-enhanced chemical vapor deposition (PECVD), 257,268 Plasma furnaces, 137-41 Plasma machining, 129 Plasma melting, 122, 129 Plasma metallurgy, 138-9 Plasma oxidation, aluminum powder, 1423 Plasma-particle interactions, 128-9 heat transfer coefficient, 128 Plasma pressure consolidation (PPC), 366 Plasma processing applications, 316 development, 2 systems, 129 see also thermal plasma processing Plasma reactors, 137-4 1 Plasma spray forming, 167-8 Plasma spraying, 122, 128, 129-36, 147 advantages, 130 applications, 123, 132-4, 136 atmospheric, 134, 135 coatings, 131, 132-4 cooling rates, 132, 134 degree of flattening, 132 deposits, 131-4

low-pressure, 134 particle size, 131 plasma torches, 125 rapid solidification, 129, 134 reactive, 135, 136 spheroidization, 135-6 surface morphology, 131 vacuum, 133, 134 Plasma torches, 123-7, 129 AC, 126 applications, 138 DC, 124-6 electrothermal efficiency, 126 non-transferred arc mode, 125 plasma gases, 127 RE 126-7 transferred arc mode, 125 Plasmagen gases, 127-8 Plasmas DC, 138-9, 141-2 RE 138, 141 thermal conductivity, 128 types of, 121, 122 see also thermal plasmas Plastic deformation, of solids, 318 PLD see pulsed laser deposition (PLD) Polycrystalline films, morphology, 281 Polycrystalline materials (PCM), synthesis, 269 Powder consolidation, 347-69 densification, 347, 349 grain-size control, 361-2 in mechanical alloying, 56-8 metastability, 347, 348-9 compositional, 348-9 morphological, 348-9 topological, 348-9 methods, 356-68 Ceracon processing, 57 electrodischarge compaction, 57 explosive forming, 57 hot extrusion, 56-7 hot isostatic pressing, 56-7 plasma-activated sintering, 57

Subject Index

nanostructured materials, 326-7 and powder contamination, 355 pressure-assisted, 364-6 and rapid solidification, 347 shockwave, 368 sintering, 356-61 non-conventional, 366-8 Powder contamination in mechanical alloying, 74-5 and powder consolidation, 355 Powder formation, via rapid solidification, 24, 27 Powder metallurgy, and spray forming compared, 153-4 Powder-rolling, strained, 326 Powders agglomeration inhibition, 52 compaction, 357-9 handling precautions, 81 nanocrystalline, 319 oversprayed, 156 oxide, 140 particle size, in mechanical alloying, 51 ultrafine, 140, 305 s e e a l s o metastable powders; nanopowders PPC s e e plasma pressure consolidation (PPC) Precipitation-hardening, 49 Primary dendrite spacing, 33 Printed wiring boards, and laser processing, 114 Process control agents (PCAs), 50 applications, 52 and powder contamination, 74 as surface-active agents, 52 Pulsed laser deposition (PLD), 99 applications, 99-100 Pulsed lasers, 93-4, 98 cooling rates, 94 picosecond, 93 Pulsed plasma methods, 227 PVD s e e physical vapor deposition (PVD)

433

Q-HIP s e e quasi-hot isostatic pressing (QHIP) Q-switched sources, lasers, 89 Quantum Rice-Ramsperger-Kassel (QRRK) model, 264 Quasi-hot isostatic pressing (Q-HIP), 363 Quasicrystalline phases, 29, 49, 58 icosahedral, 62 in intermetallic synthesis, 62 non-equilibrium, 62 Quenching rate, determination, 94 Radial distribution function (RDF), 379, 381 Radiation, coherent, 89 Ranz-Marshall equation, 128 Rapid alloying, 115 Rapid solidification, 23-43 amorphous phase formation, 376 applications, 41 constitution changes, 28-33 development, 5 disordering processes, 64 in feedstocks, 43 free energy curves, 7 front velocities, 24, 28 and grain size, 32-3 growth temperature, 32 growth velocity, 23 and laser melting, 92, 95 materials coarsening, 40 corrosion resistance, 41 creep resistance, 41 properties, 34-42 and mechanical alloying compared, 58 methods, 24-8 droplet, 24-6 spinning, 24, 26-7 surface-melting, 24, 27-8 microstructure formation, 28-33 nanostructured materials preparation, 316 non-equilibrium phases, 29-32

434

Subject

in plasma spraying, 129, 134 and powder consolidation, 347 processes, 23 processing, 1 and spray forming, 168, 173-6 time-temperature-transformation diagrams, 29-32 and tungsten inert gas welding, 110 use of term, 23 Rapid thermal processing chemical vapor deposition (RTPCVD), 257 RBS s e e Rutherford backscattering spectrometry (RBS) R D F s e e radial distribution function (RDF) Reaction milling (RM), processes, 50 Reaction temperatures, in combustion synthesis, 291-4, 299, 301-2 Reactive plasma spraying, 135, 136 Reactive spray forming, 160-3, 168 Regular solution model, 11-12 Remanence, enhancement, 73,334, 335 RF plasma torches, 126-7 Rice-Ramsperger-Kassel-Marcus (RRKM) model, 264 R M s e e reaction milling (RM) Rod mills, 54 Rotating water bath, processes, 27 RRKM model s e e Rice-RamspergerKassel-Marcus (RRKM) model RSP s e e rapid solidification RTPCVD s e e rapid thermal processing chemical vapor deposition (RTPCVD) Rutherford backscattering spectrometry (RBS), in ion mixing, 205

Index

Scanning tunneling microscopy (STM), 272 Schaeffler diagram, 107-8 Secondary dendrite arm spacing, 33 Segregation-free solidification (SFS), 32 Self-propagating high-temperature synthesis (SHS) s e e combustion synthesis (CS) Self-supporting reaction waves, initiation, 294 Self-sustained high-temperature processing, 153 SEM s e e scanning electron microscopy (SEM) Severe plastic deformation consolidation (SPDC), 364 SFS s e e segregation-free solidification (SFS) Shaker mills, 53 Shear stress, critical resolved, 184 Shock waves, 290 SHS s e e combustion synthesis (CS) Silanes, reactive sticking coefficient, 265 Silica, ultrafine, 140 Silicides, reaction temperatures, 292-3 Silicon carbide amorphous, 99 powders, 140 Silicon-carbon films, 247 Silicon irons, applications, 40 Silver-copper alloys, solidification microstructure selection diagram, 34 Silver-copper system, 9, 10 lattice parameter deviation, 12 Silver-molybdenum system, free energy diSAP s e e sintered aluminum powder (SAP) agram, 208 SBE s e e single-beam plus evaporation Simultaneous combustion, 296 Single-beam plus evaporation (SBE), 233, (SBE) Scaling laws 234 Herring's, 354 Sinter-forging, 326 in powder consolidation, 353-4 Sintered aluminum powder (SAP), 162 Scanning electron microscopy (SEM), ap- Sintering plications, 278 conventional, 356-62 field-assisted, 366-8

Subject Index

metastable powders, 352-3 microwave, 366 nanopowders, 328, 351,359-61 non-conventional, 366-8 powder consolidation, 356-61 see also field-activated sintering technique (FAST); plasma-activated sintering (PAS) Skull melting, 158 Sliding wear, applications, 319 SMSD see solidification microstructure selection diagrams (SMSD) Sodium azide, as nitrogen source, 298 SOFCs see solid oxide fuel cells (SOFCs) Solid oxide fuel cells (SOFCs), 134 Solid phase formation, chemical vapor deposition, 267-81 Solid phases, interfacial energy, 9 Solid solubility extensions, 32, 59-62 aluminum solutes, 31 in spray forming, 176, 178-9, 185, 188 via non-equilibrium processing, 59 Solid solubility limits, 327-8 Solid solutions, 213,306 composites, 305 enthalpy of formation, 201 hardening, 36-7 in intermetallic synthesis, 64-6 in liquid immiscible systems, 61 phase diagrams, 17 To curves, 7 Vegard's law, 11-12 see also supersaturated solid solutions Solidification cooling rates, 24 kinetics, 96 models, 97 partitionless, 7 see also rapid solidification Solidification front, 23, 28 Solidification microstructure selection diagrams (SMSD), 32, 33-5 Solidification microstructures, alloy phase formation, 41

435

Solute partition coefficients, 33 Solute trapping, 32 Spark erosion, applications, 319 SPDC see severe plastic deformation consolidation (SPDC) Specific heat, measurement, 9 SPEX mills, 53 Spheroidization, 154-5 Splat formation, 25 Splat quenching, 23, 132 cooling rates, 30 early studies, 1 Spot welding, 27 Spray atomization, 154, 155, 166-7, 173, 176, 180, 188 cooling effects, 174-6 gas flow field, 155-6 gas jet nozzles, 160 use of term, 166 Spray conversion processing method, in nanocrystalline materials synthesis, 318 Spray deposition, 23, 25, 43, 154, 155,176, 180 processes, 156 Spray-formed materials grain size, 181 properties, 181-3 Spray forming, 41, 153-89 advantages, 168-72 applications, 168-72 atomizers, 158-63 circular vs. linear, 158-9, 161 close-coupled vs. free-fall, 159, 162 and co-injection, 163-6 coherent particles, 184-5 concept of, 154 cooling rates, 175 crucibles, 158 deposition, 176-8 development, 2 effects direct, 183-5 indirect, 185-8

436

Subject Index

environmental chamber, 156 grain size, 180-1, 183-5 and ingot casting compared, 153-4 limitations, 168 macrosegregation, 168, 172, 179, 183, 185, 187, 188 metastable phases, 178 microsegregation, 168, 172, 179, 183, 185 near net shape, 163 nomenclature, 166-8 non-equilibrium phenomena, 172-88 plasma, 167-8 and powder metallurgy compared, 153-4 principles, 154-6 and rapid solidification, 168, 173-6 reactive, 160-3, 168 research, 189 solid solubility extensions, 178-9, 185, 188 solidification rates, 175 surface tension, 155 undercooling, 173, 174 use of term, 157, 166 variations, 157-66 Spray pyrolysis process, 305 Sputtering, 225-6 applications, 316 ion-assisted, 225 in physical vapor deposition, 229-30 planar diode, 230, 231 thermal plasmas, 122 s e e a l s o closed field unbalanced magnetron sputtering (CFUBMS); magnetron sputtering Sputtering coefficients, 229-30 SRAM s e e static random access memory (SRAM) Stage turbine vanes, 78 Stainless steels, type 316, grain size, 32 Static random access memory (SRAM), 100 Stearic acid, as process control agent, 52

Steels bearing, 41 corrosion behavior, 139 creep strength, 139 laser beam welding, 106-8 stainless, 32 thermal plasma processing, 138-9 tool, 40 Steered arc system, 232 STM s e e scanning tunneling microscopy (STM) Stranski-Krastanov mechanism, 276 Stress corrosion cracking, resistance, 79 Structural difference rule, 198 Structural materials, welding of, 106-12 Subatmospheric-pressure plasmas, 122 Substrate-adsorbate interactions, 275 Sulfidation, resistance, 78 Super Misuni NEV-MA-8 (mill), 54 Superconducting alloys, 1 Superconducting materials, 100 ion implantation, 197 Supercooled liquids, 29 bulk amorphous alloys, 406-13 Superlattice reflections, in ordered intermetallics, 65, 66 Superparamagnetic materials, 333 Superplastic materials, 339 Supersaturated phases lattice parameter, 11-12 Ostwald's step rule, 15 Supersaturated solid solutions, 1, 2, 49, 58, 62,213 formation, 67, 109, 212, 307 Supersaturation, in chemical vapor deposition, 270-1 Surface-active agents, process control agents, 52 Surface alloying, 89 Surface melting, 41 non-equilibrium effects, 27 and rapid solidification, 27-8 s e e a l s o laser surface melting

Subject Index

437

rapid solidification, 29-32 Tip stability model, 33 Ti3SiCe, synthesis, combustion synthesis reactions, 299-301 To curves, metastable phase diagrams, 7-8 Titanium alloys, formation, 101-2 TAB s e e tape automated bonding (TAB) Tape automated bonding (TAB), and laser Titanium aluminides, nanocrystalline, 328 Titanium-aluminum-nitrogen films, added processing, 115 energy, 242 Target poisoning, 237 TEM s e e transmission electron microscopy Titanium-boron binary system, controlled (TEM) reactions, 301-4 Titanium dioxide, production, 140 Temperature-programmed desorption Titanium nitride, production, 140-1 (TPD), 272 Toluene, as process control agent, 52 TFTs s e e thin film transistors (TFTs) Thermal barrier coatings, 133, 134 Tool steels, 41 Thermal expansion mismatch, 133 TPD s e e temperature-programmed desorption (TPD) Thermal fatigue, resistance, 78 Thermal plasma processing, 121-48 Transformers core materials, 104 advantages, 121-3 laminations, 1 applications, 139 furnaces, 137-41 magnetic materials, 333 silicon irons, 41 homogeneous nucleation, 122, 147 Transgranular dimple rupture, 185 power sources, 129 Transmission electron microscopy (TEM), quench rates, 121, 122, 146-7 321 reactors, 130, 137-41 supercooling, 147 Traverse welding, 27 Tribological behavior, 99 systems, 129 Triple defects, 64 Thermal plasmas, 123-9 generation, 123 TTT diagrams s e e time-temperaturetransformation (TTT) diagrams power density, 123 Tungsten inert gas (TIG) welding, 108-11 Thermal shock, resistance, 133 cooling rates, 110 Thermal spike, concept, 200 hardness profiles, 108, 110, 111 Thermocouples, 79 Thin film deposition, 92, 230 weld fusion zones, 109-10, 111 Thin film transistors (TFTs), laser anneal- Turbine blades, 77, 78 ing, 103 Thin-strip casting, 41 Undercooled liquids Ti-6wt%A1-4wt%V, grain size, 32 frozen state, 6 homogeneous nucleation frequencies, 13 TiCN, chemical vapor deposition, 269 TIG welding s e e tungsten inert gas (TIG) Undercooling welding I and crystal growth rate, 18 Time-temperature-transformation (TTT) Ostwald's step rule, 15 and rapid solidification, 23, 24, 32 diagrams bulk amorphous alloys, 382 and spray forming, 173, 174 laser processing, 96 Uni-Ball-Mill, 54 Surface-mount technology, and laser processing, 112, 114

438

Subject Index

Vacuum furnaces, fixtures, 79 Vacuum plasma spraying (VPS), 133, 134 Vapor condensation, 147 Vapor deposition classification, 257 development, 5 see a l s o chemical vapor deposition (CVD); physical vapor deposition (PVD) Vegard's law, 19, 306 solid solutions, 11-12 Very large scale integration (VLSI) technology, ion mixing in, 197 Vibrating frame mills, 54 Vitrification, solid-state, 218 VLSI technology see very large scale integration (VLSI) technology Volmer-Weber mechanism, 276 VPS see vacuum plasma spraying (VPS) Wastes, hazardous, treatment, 129, 141 Welding

cold, 52, 55, 79 full penetration, 108 spot, 27 of structural materials, 106-12 traverse, 27 see a l s o fusion welding; laser beam welding; tungsten inert gas (TIG) welding Wiedemann-Franz law, 92 Yttrium-molybdenum system amorphous phase formation, 208 free energy diagram, 211 Yttrium-niobium system, free energy diagram, 204 Zirconia coatings, 133 production, 137 Zirconium oxide, ultrafine, 140 Zone I structure, 238