TRIBOLOGY OF CERAMICS AND COMPOSITES
TRIBOLOGY OF CERAMICS AND COMPOSITES
A Materials Science Perspective
BIKRAMJIT BASU Department of Materials Science and Engineering Indian Institute of Technology Kanpur, India Materials Research Center Indian Institute of Science Bangalore, India
MITJAN KALIN Faculty of Mechanical Engineering University of Ljubljana Ljubljana, Slovenia
A JOHN WILEY & SONS, INC., PUBLICATION
Copyright © 2011 by The American Ceramic Society. All rights reserved. Published by John Wiley & Sons, Inc., Hoboken, New Jersey. Published simultaneously in Canada. No part of this publication may be reproduced, stored in a retrieval system, or transmitted in any form or by any means, electronic, mechanical, photocopying, recording, scanning, or otherwise, except as permitted under Section 107 or 108 of the 1976 United States Copyright Act, without either the prior written permission of the Publisher, or authorization through payment of the appropriate per-copy fee to the Copyright Clearance Center, Inc., 222 Rosewood Drive, Danvers, MA 01923, (978) 750-8400, fax (978) 750-4470, or on the web at www.copyright.com. Requests to the Publisher for permission should be addressed to the Permissions Department, John Wiley & Sons, Inc., 111 River Street, Hoboken, NJ 07030, (201) 748-6011, fax (201) 748-6008, or online at http://www.wiley.com/go/ permissions. Limit of Liability/Disclaimer of Warranty: While the publisher and author have used their best efforts in preparing this book, they make no representations or warranties with respect to the accuracy or completeness of the contents of this book and specifically disclaim any implied warranties of merchantability or fitness for a particular purpose. No warranty may be created or extended by sales representatives or written sales materials. The advice and strategies contained herein may not be suitable for your situation. You should consult with a professional where appropriate. Neither the publisher nor author shall be liable for any loss of profit or any other commercial damages, including but not limited to special, incidental, consequential, or other damages. For general information on our other products and services or for technical support, please contact our Customer Care Department within the United States at (800) 762-2974, outside the United States at (317) 572-3993 or fax (317) 572-4002. Wiley also publishes its books in a variety of electronic formats. Some content that appears in print may not be available in electronic formats. For more information about Wiley products, visit our web site at www.wiley.com. Library of Congress Cataloging-in-Publication Data: Basu, Bikramjit. â•… Tribology of ceramics and composites : a materials science perspective / Bikramjit Basu and Mitjan Kalin. â•…â•…â•… p. cm. â•… Includes index. â•… ISBN 978-0-470-52263-9 (cloth) 1.╇ Ceramic materials–Mechanical properties.â•… 2.╇ Ceramic materials–Fatigue.â•… 3.╇ Mechanical wear.â•… 4.╇ Friction.â•… 5.╇ Tribology.â•… I.╇ Kalin, Mitjan.â•… II.╇ Title. â•… TA455.C43B38 2011 â•… 621.8'9–dc22 2010045250 oBook ISBN: 978-1-118-02166-8 ePDF ISBN: 978-1-118-02164-4 ePub ISBN: 978-1-118-02165-1 Printed in the United States of America. 10â•… 9â•… 8â•… 7â•… 6â•… 5â•… 4â•… 3â•… 2â•… 1
Bikramjit Basu dedicates this book with a great sense of gratitude to his parents, Mr. Manoj Mohan Basu and Mrs. Chitra Basu
Mitjan Kalin would like to dedicate this book to Matija, his inspiration, pride, and happiness; and to Janja, for her understanding and support
CONTENTS
PREFACE
xvii
FOREWORD BY PROF. IAN HUTCHINGS
xxi
FOREWORD BY PROF. KARL-HEINZ ZUM GAHR
xxiii
ABOUT THE AUTHORS SECTION Iâ•… CHAPTER 1â•…
2.1 2.2 2.3 2.4 2.5
3.4 3.5
INTRODUCTION
3
6
OVERVIEW: TRIBOLOGICAL MATERIALS
Introductionâ•… 7 Definition and Classification of Ceramicsâ•… Properties of Structural Ceramicsâ•… 9 Applications of Structural Ceramicsâ•… 11 Closing Remarksâ•… 14 Referencesâ•… 16
CHAPTER 3â•…
3.1 3.2 3.3
FUNDAMENTALS
Referencesâ•…
CHAPTER 2â•…
xxv
7
8
OVERVIEW: MECHANICAL PROPERTIES OF CERAMICS
Theory of Brittle Fractureâ•… 18 Cracking in Brittle Materialsâ•… 23 Definition and Measurement of Basic Mechanical Propertiesâ•… 3.3.1 Hardnessâ•… 24 3.3.2 Compressive Strengthâ•… 27 3.3.3 Flexural Strengthâ•… 28 3.3.4 Elastic Modulusâ•… 30 3.3.5 Fracture Toughnessâ•… 31 Toughening Mechanismsâ•… 33 Closing Remarksâ•… 37 Referencesâ•… 37
18
24
vii
viiiâ•…
CONTENTS
CHAPTER 4â•…
4.1 4.2 4.3 4.4 4.5
CHAPTER 5â•…
5.1 5.2 5.3 5.4 5.5
44
49
54
FRICTIONAL HEATING AND CONTACT TEMPERATURE
Tribological Process and Contact Temperatureâ•… 60 Concept of “Bulk” and “Flash” Temperatureâ•… 61 Importance and Relevance of Some Ready-to-Use Analytical Modelsâ•… Review of Some Frequently Employed Ready-to-Use Modelsâ•… 64 6.4.1 Assumptions in Various Modelsâ•… 65 6.4.2 Model Descriptions and Implicationsâ•… 65 6.4.2.1 Archard Modelâ•… 66 6.4.2.2 Kong–Ashby Modelâ•… 67 Referencesâ•… 68
CHAPTER 7â•…
7.1 7.2
39
FRICTION
Introductionâ•… 49 Laws of Frictionâ•… 49 Friction Mechanismsâ•… 51 Friction of Common Engineering Materialsâ•… Closing Remarksâ•… 58 Referencesâ•… 59
CHAPTER 6â•…
6.1 6.2 6.3 6.4
SURFACES AND CONTACTS
Surface Roughnessâ•… 39 Surface Topography and Asperitiesâ•… 41 Real Contact Areaâ•… 42 Contact Load Distribution and Hertzian Stressesâ•… Closing Remarksâ•… 47 Referencesâ•… 48
60
63
WEAR MECHANISMS
Introductionâ•… 70 Classification of Wear Mechanismsâ•… 72 7.2.1 Adhesive Wearâ•… 73 7.2.2 Abrasive Wearâ•… 75 7.2.2.1 Abrasion of Compositesâ•… 77 7.2.3 Fatigue Wearâ•… 78 7.2.4 Oxidation and Tribochemical Wearâ•… 80 7.2.5 Fretting Wearâ•… 81 7.2.5.1 Fretting Modesâ•… 82 7.2.5.2 Mechanics of Elastic Contacts under Fretting Conditionsâ•… 7.2.5.3 Mechanics of Elastic–Plastic Contacts under Fretting Conditionsâ•… 86 7.2.5.4 Fretting Regimesâ•… 86 7.2.5.5 Determination of Fretting Regimesâ•… 89 7.2.5.6 Fretting Mapsâ•… 89 7.2.5.7 Velocity Accommodation in Frettingâ•… 91 7.2.5.8 Friction Logsâ•… 92 7.2.6 Solid Particle Erosionâ•… 92 7.2.6.1 Erosion of Ductile Materialsâ•… 94 7.2.6.2 Erosion of Brittle Materialsâ•… 96
70
84
â•… ix
CONTENTS
7.3
Closing Remarksâ•… Referencesâ•… 99
CHAPTER 8â•…
8.1 8.2
LUBRICATION
Lubrication Regimesâ•… Stribeck Curveâ•… 107 Referencesâ•… 109
SECTION IIâ•… CHAPTER 9â•…
9.1 9.2 9.3 9.4 9.5
9.6 9.7 9.8 9.9 9.10
FRICTION AND WEAR OF STRUCTURAL CERAMICS
OVERVIEW: STRUCTURAL CERAMICS
113
CASE STUDY: TRANSFORMATION-TOUGHENED ZIRCONIA
Backgroundâ•… 142 Wear Resistanceâ•… 144 Morphological Characterization of the Worn Surfacesâ•… 146 Zirconia Phase Transformation and Wear Behaviorâ•… 149 Wear Mechanismsâ•… 152 Relationship among Microstructure, Toughness, and Wearâ•… 154 Influence of Humidity on Tribological Properties of Self-Mated Zirconiaâ•… Wear Mechanisms in Different Humidityâ•… 157 Tribochemical Wear in High Humidityâ•… 160 Closing Remarksâ•… 163 Referencesâ•… 164
CHAPTER 11â•…
11.1 11.2 11.3 11.4
101 101
Introductionâ•… 113 Zirconia Crystal Structures and Transformation Characteristics of Tetragonal Zirconiaâ•… 114 Transformation Tougheningâ•… 116 9.3.1 Micromechanical Modelingâ•… 116 Stabilization of Tetragonal Zirconiaâ•… 117 Different Factors Influencing Transformation Tougheningâ•… 118 9.5.1 Grain Sizeâ•… 119 9.5.2 Yttria Contentâ•… 121 9.5.3 Yttria Distributionâ•… 122 Stress-Induced Microcrackingâ•… 125 Development of SiAlON Ceramicsâ•… 126 Microstructure of S-sialon Ceramicsâ•… 127 Mechanical Properties and Crack Bridging of SiAlON Ceramicâ•… 129 Properties of Titanium Diboride Ceramicsâ•… 132 Referencesâ•… 138
CHAPTER 10â•…
10.1 10.2 10.3 10.4 10.5 10.6 10.7 10.8 10.9 10.10
98
142
156
CASE STUDY: SIALON CERAMICS
Introductionâ•… 167 Materials and Experimentsâ•… 168 Tribological Properties of Compositionally Tailored Sialon versus β-Sialonâ•… Tribological Properties of S-Sialon Ceramicâ•… 179
167
172
xâ•…
CONTENTS
11.5
Concluding Remarksâ•… Referencesâ•… 183
CHAPTER 12â•…
12.1 12.2 12.3 12.4 12.5 12.6
182
CASE STUDY: MAX PHASE—TI3SIC2
Backgroundâ•… 185 Frictional Behaviorâ•… 188 Wear Resistance and Wear Mechanismsâ•… 188 Raman Spectroscopy and Atomic Force Microscopy Analysisâ•… Transition in Wear Mechanismsâ•… 193 Summaryâ•… 194 Referencesâ•… 195
185
190
CHAPTER 13â•…
CASE STUDY: TITANIUM DIBORIDE CERAMICS AND COMPOSITES 13.1 13.2 13.3
13.4 13.5
197
Introductionâ•… 197 Materials and Experimentsâ•… 198 Tribological Properties of TiB2–MoSi2 Ceramicsâ•… 200 13.3.1 Friction and Wearâ•… 200 13.3.2 Wear and Dissipated Energyâ•… 202 13.3.3 Wear and Abrasion Parameterâ•… 203 13.3.4 Material Removal Mechanismsâ•… 204 Tribological Properties of TiB2–TiSi2 Ceramicsâ•… 204 Closing Remarksâ•… 206 Referencesâ•… 208
SECTION IIIâ•…
FRICTION AND WEAR OF BIOCERAMICS AND BIOCOMPOSITES CHAPTER 14â•…
14.1 14.2
14.3 14.4 14.5 14.6 14.7
14.8
213
Introductionâ•… 213 Some Useful Definitions and Their Implicationsâ•… 215 14.2.1â•… Biomaterialsâ•… 215 14.2.2 Biocompatibilityâ•… 216 14.2.3 Host Responseâ•… 216 Experimental Evaluation of Biocompatibilityâ•… 217 Wear of Implantsâ•… 221 Coating on Metalsâ•… 223 Glass-Ceramicsâ•… 224 Biocompatible Ceramicsâ•… 226 14.7.1 Bioinert Ceramicsâ•… 226 14.7.2 Calcium Phosphate-Based Biomaterialsâ•… 226 Outlookâ•… 228 Referencesâ•… 229
CHAPTER 15â•…
15.1 15.2
OVERVIEW: BIOCERAMICS AND BIOCOMPOSITES
CASE STUDY: POLYMER-CERAMIC BIOCOMPOSITES
Introductionâ•… 233 Materials and Experimentsâ•…
235
233
â•… xi
CONTENTS
15.3 15.4 15.5 15.6 15.7
Frictional Behaviorâ•… 237 Wear-Resistance Propertiesâ•… 240 Wear Mechanismsâ•… 242 Correlation among Wear Resistance, Wear Mechanisms, Material Properties, and Contact Pressureâ•… 247 Concluding Remarksâ•… 248 Referencesâ•… 249
CHAPTER 16â•…
CASE STUDY: NATURAL TOOTH AND DENTAL RESTORATIVE
MATERIALS
251
16.1 16.2
Introductionâ•… 251 Materials and Methodsâ•… 254 16.2.1 Preparation of Human Tooth Materialâ•… 254 16.3 Tribological Tests on Tooth Materialâ•… 255 16.4 Production and Characterization of Glass-Ceramicsâ•… 255 16.5 Wear Experiments on Glass-Ceramicsâ•… 256 16.6 Microstructure and Hardness of Human Tooth Materialâ•… 257 16.7 Tribological Properties of Human Tooth Materialâ•… 260 16.7.1 Friction Behaviorâ•… 260 16.7.2 Wear Mechanismsâ•… 262 16.8 Wear Properties of Glass-Ceramicsâ•… 262 16.9 Discussion of Wear Mechanisms of Glass-Ceramicsâ•… 266 16.10 Comparison with Existing Glass-Ceramic Materialsâ•… 271 16.11 Concluding Remarksâ•… 273 Referencesâ•… 274 CHAPTER 17â•…
17.1 17.2 17.3 17.4 17.5 17.6 17.7
276
Introductionâ•… 276 Materials and Experimentsâ•… 277 Frictional Propertiesâ•… 278 Wear Resistance and Wear Mechanismsâ•… 278 Wear Debris Analysis and Tribochemical Reactionsâ•… 282 Influence of Glass Infiltration on Wear Propertiesâ•… 283 Concluding Remarksâ•… 284 Referencesâ•… 285
CHAPTER 18â•…
18.1 18.2 18.3
CASE STUDY: GLASS-INFILTRATED ALUMINA
TRIBOLOGICAL PROPERTIES OF CERAMIC BIOCOMPOSITES
287
Backgroundâ•… 287 Tribological Properties of Mullite-Reinforced Hydroxyapatiteâ•… 288 Friction and Wear Rateâ•… 288 18.3.1 Effect of Simulated-Body-Fluid Medium on Wear of Mullite-Reinforced Hydroxyapatiteâ•… 289 18.3.2 Surface Topography of Mullite-Reinforced Hydroxyapatite after Fretting Wearâ•… 293 18.3.3 Frictional Behaviorâ•… 293 18.3.4 Wear Micromechanisms of Hydroxyapatite-Based Materials in Simulated Body Fluidâ•… 296
xiiâ•… 18.4
CONTENTS
Concluding Remarksâ•… Referencesâ•… 302
SECTION IVâ•… CHAPTER 19â•…
19.1 19.2 19.3
19.4 19.5
298
FRICTION AND WEAR OF NANOCERAMICS
OVERVIEW: NANOCERAMIC COMPOSITES
307
Introductionâ•… 307 Processing of Bulk Nanocrystalline Ceramicsâ•… 309 Overview of Developed Nanoceramics and Ceramic Nanocompositesâ•… 19.3.1 Monolithic Nanoceramicsâ•… 311 19.3.2 Alumina-Based Nanocompositesâ•… 313 19.3.3 Tungsten Carbide-Based Nanocompositesâ•… 314 19.3.4 Zirconia-Based Nanocompositesâ•… 317 Overview of Tribological Properties of Ceramic Nanocompositesâ•… 318 Concluding Remarksâ•… 320 Referencesâ•… 322
309
CHAPTER 20â•…
CASE STUDY: NANOCRYSTALLINE YTTRIA-STABILIZED TETRAGONAL ZIRCONIA POLYCRYSTALLINE CERAMICS 20.1 20.2 20.3 20.4 20.5 20.6
325
Introductionâ•… 325 Materials and Experimentsâ•… 327 Tribological Propertiesâ•… 329 Tribomechanical Wear of Yttria-Stabilized Zirconia Nanoceramic with Varying Yttria Dopantâ•… 330 Comparison with Other Stabilized Zirconia Ceramicsâ•… 335 Concluding Remarksâ•… 335 Referencesâ•… 336
CHAPTER 21â•…
CASE STUDY: NANOSTRUCTURED TUNGSTEN CARBIDE–ZIRCONIA NANOCOMPOSITES 21.1 21.2 21.3 21.4 21.5 21.6
Introductionâ•… 338 Materials and Experimentsâ•… 339 Friction and Wear Characteristicsâ•… 340 Wear Mechanismsâ•… 345 Explanation of High Wear Resistance of Ceramic Nanocompositesâ•… Concluding Remarksâ•… 349 Referencesâ•… 349
SECTION Vâ•… CHAPTER 22â•…
347
LIGHTWEIGHT COMPOSITES AND CERMETS
OVERVIEW: LIGHTWEIGHT METAL MATRIX COMPOSITES AND
CERMETS 22.1 22.2
338
Development of Metal Matrix Compositesâ•… Development of Cermetsâ•… 356 Referencesâ•… 358
353 353
CONTENTS
â•… xiii
CHAPTER 23â•…
CASE STUDY: MAGNESIUM–SILICON CARBIDE PARTICULATEREINFORCED COMPOSITES 23.1 23.2 23.3 23.4 23.5 23.6
362
Introductionâ•… 362 Materials and Experimentsâ•… 363 Load-Dependent Friction and Wear Propertiesâ•… 363 Fretting-Duration-Dependent Tribological Propertiesâ•… 366 Tribochemical Wear of Magnesium–Silicon Carbide Particulate-Reinforced Compositesâ•… 371 Concluding Remarksâ•… 375 Referencesâ•… 376
CHAPTER 24â•…
CASE STUDY: TITANIUM CARBONITRIDE–NICKELBASED CERMETS 24.1 24.2 24.3 24.4
24.5 24.6 24.7
24.8
Introductionâ•… 377 Materials and Experimentsâ•… 379 Energy Dissipation and Abrasion at Low Loadâ•… 381 Influence of Type of Secondary Carbides on Sliding Wear of Titanium Carbonitride–Nickel Cermetsâ•… 386 24.4.1 Wear Mechanismsâ•… 387 Tribochemical Wear of Titanium Carbonitride–Based Cermetsâ•… 387 24.5.1 Evolution of Tribochemistry and Contact Temperatureâ•… 387 Influence of Tungsten Carbide Content on Load-Dependent Sliding Wear Propertiesâ•… 393 High Temperature Wear of Titanium Carbonitride–Nickel Cermetsâ•… 397 24.7.1 Wear Mechanismsâ•… 398 24.7.2 Discussion of High-Temperature Oxidation and Its Relation to Material Removalâ•… 401 24.7.3 Thermal Oxidationâ•… 402 24.7.4 Influence of Different Secondary Carbide Additionâ•… 403 Summary of Key Resultsâ•… 403 Referencesâ•… 404
CHAPTER 25â•…
25.1 25.2 25.3 25.4 25.5 25.6
377
CASE STUDY: (W,Ti)C–CO CERMETS
407
Introductionâ•… 407 Materials and Experimentsâ•… 408 Microstructure and Mechanical Propertiesâ•… 409 Wear Propertiesâ•… 410 Correlation between Mechanical Properties and Wear Resistanceâ•… Concluding Remarksâ•… 418 Referencesâ•… 419
413
SECTION VIâ•…
FRICTION AND WEAR OF CERAMICS IN A CRYOGENIC ENVIRONMENT CHAPTER 26â•…
26.1 26.2
OVERVIEW: CRYOGENIC WEAR PROPERTIES OF MATERIALS
Backgroundâ•… 423 Designing a High-Speed Cryogenic Wear Testerâ•…
425
423
xivâ•… 26.3
26.4
CONTENTS
Summary of Results Obtained with Ductile Metalsâ•… 26.3.1 Self-Mated Steelâ•… 427 26.3.2 Titanium/Steel Coupleâ•… 430 26.3.3 Copper/Steel Sliding Systemâ•… 433 Summaryâ•… 437 Referencesâ•… 437
427
CHAPTER 27â•…
CASE STUDY: SLIDING WEAR OF ALUMINA IN A CRYOGENIC ENVIRONMENT 27.1 27.2 27.3 27.4
27.5
439
Backgroundâ•… 439 Materials and Experimentsâ•… 440 Tribological Properties of Self-Mated Aluminaâ•… 442 Genesis of Tribological Behavior in a Cryogenic Environmentâ•… 449 27.4.1 Friction of Self-Mated Alumina in LN2â•… 449 27.4.2 Brittle Fracture and Wear of Self-Mated Alumina in LN2â•… 450 Concluding Remarksâ•… 452 Referencesâ•… 452
CHAPTER 28â•…
CASE STUDY: SLIDING WEAR OF SELF-MATED TETRAGONAL ZIRCONIA CERAMICS IN LIQUID NITROGEN 28.1 28.2 28.3 28.4 28.5 28.6 28.7
Introductionâ•… 454 Materials and Experimentsâ•… 456 Friction of Self-Mated Y-TZP Material in LN2â•… 456 Cryogenic Wear of Zirconiaâ•… 459 Cryogenic Sliding-Induced Zirconia Phase Transformationâ•… Wear Mechanisms of Zirconia in LN2â•… 464 Concluding Remarksâ•… 466 Referencesâ•… 467
454
460
CHAPTER 29â•…
CASE STUDY: SLIDING WEAR OF SILICON CARBIDE IN A CRYOGENIC ENVIRONMENT 29.1 29.2 29.3 29.4 29.5 29.6 29.7
469
Introductionâ•… 469 Materials and Experimentsâ•… 470 Friction and Wear Propertiesâ•… 470 Thermal Aspect and Limited Tribochemical Wearâ•… 473 Tribomechanical Stress-Assisted Deformation and Damageâ•… 479 Comparison with Sliding Wear Properties of Oxide Ceramicsâ•… 481 Concluding Remarksâ•… 482 Referencesâ•… 483
SECTION VIIâ•…
WATER-LUBRICATED WEAR OF CERAMICS
CHAPTER 30â•…
FRICTION AND WEAR OF OXIDE CERAMICS IN AN AQUEOUS ENVIRONMENT 30.1 30.2
Backgroundâ•… 487 Tribological Behavior of Alumina in an Aqueous Solutionâ•…
488
487
â•… xv
CONTENTS
30.3
30.4
30.2.1 Electrochemical Properties and Wear Characterization of Self-Mated Aluminaâ•… 491 30.2.2 Surface Roughness and Frictional Behaviorâ•… 492 Tribological Behavior of Self-Mated Zirconia in an Aqueous Environmentâ•… 493 30.3.1 Zirconia Transformation and Wearâ•… 497 30.3.2 Electrochemical Aspect of Wearâ•… 498 Concluding Remarksâ•… 499 Referencesâ•… 500
SECTION VIIIâ•…
CLOSURE
CHAPTER 31â•…
PERSPECTIVE FOR DESIGNING MATERIALS FOR TRIBOLOGICAL APPLICATIONS
505
INDEX
509
PREFACE
Tribology, by definition, is the science and technology of interacting surfaces in relative motion. Such scientific understanding has significant technological relevance for various engineering industries. Broadly, tribology deals with the concepts of friction, wear, and lubrication. Over the last few decades, it has been widely recognized that tribology, being an interdisciplinary area, involves the interaction of concepts drawn from multiple disciplines, including mechanical engineering, materials science, physics, and chemistry. The development of new materials (bulk or coating) with better friction and wear resistance, as well as the progress in tribology research, clearly requires an improved understanding in multiple disciplines as well as the development of new design methodologies in order to obtain better properties in relation to tribological performance. Even though tribology is still not broadly known as a research field to many in industry or academia, we are all intrigued by the topic in everyday life, as well as in almost every engineering application. Across the world, very few universities teach this subject; however, the subject is gaining importance. There are many books on tribology written from different perspectives, such as materials science, mechanics, mechanical engineering, lubrication and additives, physics, and chemistry. This book is intended to cover mostly the materials science aspects applicable to tribology science. Researchers interested in automotive, aerospace, biomaterials, hardmetals, and related applications would look for a complete set of possible materials for those applications, as well as wear and friction mechanisms. On the other hand, people from the materials science community would look for details of mechanisms, effect of microstructure, working conditions, lubrication, environment, and so on, which again are covered here due to very broad materials selection. This book places the utmost importance on the microstructure–material-properties–tribological-properties relationship for the range of advanced materials that are covered herein. The description of the wear micromechanisms of the various materials will provide a strong background to readers on how to design and develop new tribological materials. From the aforementioned perspective, this book is structured into various thematic sections, and each section contains a number of chapters. This book was designed to motivate students and young researchers as well as to provide experts in the area with a healthy balance of topics for teaching and academic purposes, primarily for two disciplines: materials science/metallurgy and mechanical engineering. It is expected that this book, if used as a text, would strongly benefit senior undergraduate and postgraduate students. xvii
xviii
Preface
Section I of this book is designed to provide the readers with a background in the area of tribology and basic materials science. Characteristics of material surfaces in terms of surface roughness and various material properties are discussed, as well as the fundamentals of the friction, wear mechanism, and lubrication. This is followed by Section II, where the tribological properties of structural ceramics, which include zirconia, sialon, ternary carbides, and high-temperature ceramics, such as borides, are discussed. This selection of materials also represents a class of technologically important and emerging ceramics. It is shown how the microstructure and mechanical properties both determine the wear resistance of these materials. One area in which ceramics and polymers are increasingly important is biomedical applications. In Section III, the tribological properties of hydroxyapatitebased bioceramic composites are discussed first. Polymers are known for their poor wear resistance; it is shown how the development of hybrid polymer-ceramic biocomposites can lead to higher wear resistance while retaining good frictional properties of the polymer matrix. This is followed by a discussion on the wear properties of some of the stabilized zirconia ceramics. The two chapters in this section deal with the materials that are important in dental restoration. A rather recent development in the materials world is the synthesis of nanoÂ� ceramic composites. In view of this, Section IV discusses the friction and wear properties of zirconia and WC-based nanocomposites, which are processed using the advanced processing technique of spark plasma sintering. A summary of the literature on the tribological properties of various nanoceramics is also included. In the last decade, lightweight composites have been considered for use in automotive and other applications requiring good wear resistance. Similarly, new generation cermets, based on TiCN as well as mixed carbide cermets, are also being developed as a replacement for widely used WC-Co cemented carbides. Hence, Section V demonstrates how these new–generation materials will behave at tribological contacts. While our understanding of the dry, unlubricated tribological properties of various materials is extensive, such understanding in a cryogenic environment and under high speed sliding conditions is rather limited. In view of this, Section VI discusses the tribological properties of oxide and non-oxide ceramics in liquid nitrogen with reference to similar properties under ambient and room temperature sliding conditions. In Section VII, the tribological properties of various ceramics in aqueous environments are discussed, with reference to regimes and pH regions and their effect on performance. The book concludes with Section VIII, which covers the various issues to be investigated in the near future in the design and development of materials with better tribological properties. This section summarizes the information provided in the book and gives an insight into the broader knowledge of these materials and advice on how to use them in various applications. The above-described structure of this book as well as the succession of various sections and chapters is expected to be useful in helping both students and experts pursuing the area of tribology of advanced materials to gradually build up knowledge
xix
Preface
of the fundamentals and, subsequently, to understand the most recent advances. In particular, this book has the following major important features: (1) the fundamental science of tribology is presented, thus allowing the book to be used as a textbook for teaching, academic, or research purposes; (2) a broad range of materials is covered, such as advanced tough ceramics, high-temperature ceramics, biomaterials, and nanoceramics, to illustrate how the materials science aspect can be realized while analyzing the tribological results; and (3) the book will appeal to a large number of active researchers from various disciplines of metallurgy and materials science, ceramics, and mechanical engineering. This book is an outcome of several years of teaching undergraduate and postgraduate courses in the area of tribology of materials, advanced ceramics, composite materials, and biomaterials, and other related fundamental courses in materials science, which were offered to students of the Indian Institute of Technology (IIT) Kanpur, India, as well as at the Faculty of Mechanical Engineering at the University of Ljubljana, Slovenia. More important, the research results of many of the postgraduate students from our groups are also summarized in some chapters. B. Basu would like to specifically acknowledge some of his past and present students, B. V. Manoj Kumar, G. B. Raju, Shekhar Nath, P. Suresh Babu, Amartya Mukhopadhyay, K. Madhav Reddy, Animesh Choubey, M. Surender, S. Bajaj, N. Sinha, Tufan Kumar Guha, P. Maji, Rohit Khanna, Subhodip Bodhak, Srimanta Das Bakshi, D. Sarkar, Manisha Taneja, Ravi Kumar, A. Tewari, T. Venkateswaran, U. Raghunandan, Divya Jain, Nitish Kumar, Amit S. Sharma, Ashutosh K. Dubey, Alok Kumar, Sushma Kalmodia, Shilpee Jain, Neha Gupta, Indu Bajpai, Garima Tripathi, Prafulla K. Mallik, Anup Patel, Rajeev Kumar, and Atiar R. Molla. The dedication of these students to understanding the tribological properties of a range of ceramics and composites is reflected in the research work summarized in many of the chapters of this book. With gratitude, B. Basu appreciates the past and present research collaboration with a number of researchers and academicians, including Drs. Omer Van Der Biest, Jozef Vleugels, R. K. Bordia, G. Sundararajan, S. K. Mishra, A. K. Suri, R. Mitra, I. Manna, A. Basumallick, J. Ramkumar, B. Subramonian, Manoj Gupta, K. C. Hari Kumar, R. G. Vitchev, Hasan Mondal, Ferhat Kara, Nurcan Kalis Ackibas, P. Gilman, S. C. Koria, R. K. Dube, M. Karanjai, D. Roy, M. C. Chu, S. J. Cho, Doh-Yeon Kim, Jo Wook, and S. Kang. The encouragement and collaboration with two of his colleagues, the late Prof. R. Balasubramaniam and the late Prof. V. S. R. Murty, is also remembered. B. Basu also expresses sincere thanks to his long-term friend and mentor, Dr. Jaydeep Sarkar, for his constant encouragement during the writing of this book. B. Basu also remembers the constant inspiration of a number of colleagues and former teachers, including Profs./Drs. S. Ranganathan, K. Chattopadhyay, Sanjay K. Biswas, Ashutosh Sharma, N. K. Mukhopadhyay, Indranil Manna, D. Basu, Anoop K. Mukhopadhyay, Brian Lawn, M. V. Swain, M. Hoffman, Vikram Jayaram, Goutam Biswas, D. Mazumdar, Dipankar Banerjee, Atul Chokshi, and B. S. Murty. M. Kalin acknowledges cooperation related to tribology of ceramics from J. Vižintin and F. Kopacˇ from his group at the Faculty of Mechanical Engineering, and S. Novak and G. Dražicˇ from the Jožef Stefan Institute in Ljubljana, Slovenia. M. Kalin would like to acknowledge help from M. Polajnar and J. Kogovšek for
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Preface
assistance in final technical revisions of this book. He would also like to express particular thanks to S. Jahanmir (MiTi Heart Corp, USA) and K. Kato (Tohoku University, Japan) for their mentoring, scientific, and personal advice during their joint work at the Ceramics Department at the National Institute of Standards and Technology (NIST, Gaithersburg, MD), and Tohoku University (Japan), respectively, as well as for the remarkable support and kind friendship over many years. The authors would like to thank Drs. Ian Hutchings, Said Jahanmir, Koji Kato, and Karl-Heinz Zum Gahr for writing the comments on and forewords to this book. The authors would like to take this opportunity to acknowledge the financial support, of various governmental agencies of India, including the Indian Space Research Organisation (ISRO), Department of Atomic Energy (DAE), Department of Biotechnology (DBT), Defense Research and Development Organization (DRDO), Council of Scientific and Industrial Research (CSIR), Department of Science and Technology (DST), UK-India Education and Research Initiative (UKIERI), and Indo-US Science and Technology Forum (IUSSTF) in the last two decades, which facilitated research in the area of tribology of advanced materials at IIT Kanpur. B. Basu also expresses gratitude to Mr. N. M. Dube and his colleagues at DUCOM, Bangalore, for designing and fabricating custom-made fretting and high-speed cryogenic tribotesters. B. Basu expressed sincere thanks to Mr. Divakar Tiwari and his present students (Amit, Ashutosh, Anup, Neha, Indu, Shilpee, and Alok) for their untiring efforts and effective assistance during various stages of the manuscript preparation. We would also like to thank IIT Kanpur for extending financial and other support during the writing this book. The continuous financial support from the Ministry of Higher Education, Science and Technology of Slovenia, as well as the Slovenian Research Agency, over the years is also greatly appreciated. Finally, we would like to acknowledge the continuous support extended by our parents, in-laws, and family members, Pritha and Prithvijit and Janja and Matija, during the course of the writing of this book. IIT Kanpur, India, and IISc, Bangalore, India Ljubljana, Slovenia July 2011
Bikramjit Basu Mitjan Kalin
FOREWORD Engineering ceramics form a diverse and important class of materials, with a wide range of properties and applications, from rolling bearings to dental implants, and from high-performance cutting tools to artificial hip joints. In these applications and many more, the tribological behavior of the ceramic is paramount, but the properties by which the material is specified are often “standard” and easily measured ones such as density, hardness, Young’s modulus, modulus of rupture, and perhaps fracture toughness. As we now know from extensive research, these properties are often poor predictors of tribological performance. Better understanding of the behavior of ceramics in tribological applications, and of the detailed influence of microstructural features such as porosity, phase, and grain size distributions, as well as the tribochemical processes that occur at the material’s surface, will benefit all manufacturers and users of these materials and will enable their properties and value to be optimized. A deep appreciation of materials science and engineering, coupled with both the chemical and mechanical influences which act on the ceramic in use, is needed to understand the wear and friction of these materials. Fracture, plastic flow, and tribochemical processes can all play key roles in the wear and friction of ceramics. It can be argued that their tribological behavior is even more complex than that of metals. This book, which focuses on the subject from a materials science perspective, forms a valuable contribution to the literature on the tribology of engineering ceramics and their composites, and the authors are to be congratulated on its comprehensive scope. Prof. Ian Hutchings Cambridge, UK
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FOREWORD Microstructures of structural ceramics and their composites have been developed during the last decades mainly for applications under static and dynamic mechanical, thermal, or corrosive loads. Among others, the aim was to improve fracture toughness to overcome the inherent brittleness and increase the reliability of ceramic components in high-loaded applications by optimizing microstructural features, such as size and shape of grains or reinforcing phases, as well as processing technologies. However, to use the potential of ceramic materials in components under high tribological loading, materials microstructures have to be adjusted based on a competent knowledge of tribological mechanisms involved and the structure–property relationships. Using case studies, this book contributes to filling the gap in our understanding of the effects of structures of ceramic materials on tribological behavior. Beginning with fundamental aspects of structure and properties of ceramic materials as well as an introduction to tribology, it covers the tribological behavior of a wide range of materials from structural ceramics through bioceramics, biocomposites, and nanocomposites to cermets. This book can be very useful for newcomers, such as students, in the field of ceramics and tribology, as well as for readers with an interest in utilizing the high potential of ceramic materials in tribological applications. Prof. Karl-Heinz Zum Gahr Karlsruhe Institute of Technology (KIT), Germany
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ABOUT THE AUTHORS
Bikramjit Basu, PhD Associate Professor, Department of Materials Science and Engineering, Indian Institute of Technology, Kanpur, India; E-mail: bikram@ iitk.ac.in; currently at Materials Research Center, Indian Institute of Science, Bangalore, India Dr. Bikramjit Basu is currently an Associate Professor at the Indian Institute of Science, Bangalore, and is on leave from the Indian Institute of Technology (IIT), Kanpur, India. Bikramjit obtained his undergraduate and postgraduate degrees, both in Metallurgical Engineering, from the National Institute of Technology (NIT), Durgapur, and the Indian Institute of Science, Bangalore, in 1995 and 1997, respectively. He earned his PhD in Ceramics at Katholieke Universiteit Leuven, Belgium, in March 2001. He returned to India to join IIT Kanpur in November 2001 as Assistant Professor after a brief postdoctoral research experience at the University of California, Santa Barbara. He held visiting positions at the University of Warwick (U.K.), Seoul National University (South Korea), and University Polytechnic Catalonia (Spain). Dr. Basu has authored or co-authored more than 140 peer-reviewed research papers with 20 papers in the Journal of American Ceramic Society. He is the principal editor of the book Advanced Biomaterials: Fundamentals, Processing and Applications (John Wiley & Sons Inc., in association with American Ceramic Society), which was published in September 2009. He is on the editorial boards of five international journals and serves as a reviewer of more than 20 SCI journals in the area of ceramics and biomaterials. He has edited a number of special issues of various journals, including Journal of Materials Science, International Journal of Applied Ceramics Technology, and Journal of Biomedical Materials Research: Part B. At IIT Kanpur, Dr. Basu established a vibrant research program in the area of tribology, structural ceramics, and biomaterials. His research spans the interdisciplinary areas of ceramics, tribology, and biomaterials. In developing interdisciplinary research programs in tribology and ceramics, he has collaborated with the xxv
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About the Authors
materials scientists of the International Advanced Research Center for Powder Metallurgy and New Materials (ARCI), Defense Metallurgical Research Laboratory (DMRL), National Metallurgical Laboratory (NML), Central Glass and Ceramics Research Institute (CGCRI), Indian Space Research Organization (ISRO), and Bhabha Atomic Research Center (BARC). In the area of tribology, he has made significant contributions in establishing the correlation between wear micromechanisms and material properties for a large number of ceramics/composites, including toughened ceramics, such as yttriastabilized tetragonal zirconia polycrystals (Y-TZP), Ti3SiC2, sialon, and other materials, such as (W,Ti)C-Co, TiB2-based ceramics, and TiCN-Ni-XC(X=Nb/W/Ta/Hf). Using a self-designed high-speed cryo-tribometer, Dr. Basu and his co-workers performed a critical set of experiments to understand friction and wear mechanisms of high-purity metals and ceramic bearings. Such a study has relevance for space applications. His research mostly focused on developing microstructure-based understanding of wear mechanisms for various ultra-fine grained ceramics, nanocomposites, and biomaterials, as well as critically analyzing wear resistance properties in the light of the mechanics-based models. His fundamental contribution is the development of analytical models to predict the tribochemical and tribomechanical wear of ceramics. In recognition of his contributions to the field of ceramics, tribology, and biomaterials, Dr. Basu received noteworthy awards from the Indian Ceramic Society (2003), Indian National Academy of Engineering (2004), and Indian National Science Academy (2005), and was awarded the “Metallurgist of the Year” (2010), by the Ministry of Steels, Government of India. He is the first Indian from India to receive the prestigious “Coble Award for Young Scholars” from the American Ceramic Society in 2008. In 2010, he received the NASI (National Academy of Science, India)–SCOPUS Young Scientist award.
Dr. Mitjan Kalin Faculty of Mechanical Engineering, University of Ljubljana, Askerceva 6, 1000 Ljubljana, Slovenia; E-mail:
[email protected] Since obtaining his PhD in 1999 (University of Ljubljana, Slovenia), Dr. Kalin has focused primarily on the research of wear and friction mechanisms for advanced materials, such as ceramics and coatings, as well as on boundary lubrication, tribochemistry, and nanotribology. In the last 10 years, he has led about 15 single-investigator, bilateral, and multilateral projects, more than half of which were international. He has contributed to about 120 conference proceedings and 70 peer-reviewed papers. He has published eight chapters in interna-
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About the Authors
tional books, as well as eight full-text student-course scripta. He is a co-editor of the book Tribology of Mechanical Systems: A Guide to Present and Future Technologies (ASME Press, 2004). He has delivered over 30 invited lectures and talks at conferences, institutes, universities, and technology-driven companies worldwide. He has also presented about 80 reports and studies made for industrial partners in the area of maintenance, wear, and lubrication. He is author or co-author of eight patents, two of them U.S. patents. Dr. Kalin has received three awards from the Faculty of Mechanical Engineering for scientific and research work, and two awards from the Slovenian Society for Tribology for international recognition. He is also a recipient of the prestigious Burt L. Newkirk Award (ASME, 2006) and the Slovenian state Zois award (2006) for important scientific achievements. Dr. Kalin is a reviewer for over 30 international peer-reviewed journals in various fields of engineering, material science, physics, chemistry, and nanotechnology. Since 2006, he has been Associate Editor of the ASME Journal of Tribology. He is also a member of the editorial boards of the Journal of Industrial Lubrication and Tribology, Emerald (2004– ), Advances in Tribology, Hindawi (2009– ), and ISRN Mechanical Engineering, Hindawi (2010– ), and a member of the publishing council of the SV-JME Journal of Mechanical Engineering (2007– ). He has been Guest Editor of several special issues of Tribology International (Elsevier) and Lubrication Science (Wiley). He has also been recently appointed Editor of Lubrication Science (Wiley, 2012– ). He also reviews various proposals for Wiley, ASME Press, and Springer, over 30 SCI peer-reviewed journals in tribology and several related fields, as well as many proposals for national and international research agencies, such as the European Commission and the European Research Foundation. He is a secretary and member of the Executive Committee (1997– ) of the Slovenian Society for Tribology, and one of its founding members. He also serves as an executive board member (2006– ) of the Slovenian Society for Materials. He is also an active member of the Society of Tribologists and Lubrication Engineers (STLE); since 2001 he has served in various positions in the Ceramics and Composites Committee, and in 2004 he was elected as president of this committee. He has been a member of organizing, international, and/or advisory boards at many international conferences, and in 2009 acted as a chair of the Engineering Conferences International (ECI) conference “Advances in Boundary Lubrication and Surface Boundary Films” (Seville, Spain, 2009). He is currently full professor (2010– ) and Head of the Chair for Tribology, Technical Diagnostics, and Maintenance (2011– ) at the University of Ljubljana, Slovenia. He has also done postdoctoral work at NIST, the Catholic University (Leuven, Belgium), Tohoku University (Japan), and the University of Pisa (Italy). Currently, he holds the position of Vice-Dean for Research and International Affairs at the Faculty of Mechanical Engineering at the University of Ljubljana on his second term.
SECTION
I
FUNDAMENTALS
CHAPTER
1
INTRODUCTION It is imperative to define and introduce various fundamental concepts of tribology in this very first section. After a brief introduction, this chapter provides some general discussion on tribology with a particular emphasis on interdisciplinary aspects. This chapter is followed by a discussion in Chapter 2 on the typical nature and properties of metals, ceramics, and polymers that are used widely for various tribological applications. The material-intrinsic surface properties, such as hardness, strength, ductility, and work hardening, are very important factors for wear resistance. Since this book largely discusses the case studies of ceramics and composites, the mechanical properties of ceramics are reviewed in Chapter 3. However, other factors, such as load, relative speed, lubrication, temperature, and environment (ambient, inert atmosphere, relative humidity), are equally important. Importantly, the chemical nature and compatibility of mating materials with environment and lubricants as well as their interplay—which defines tribochemical reactions—will have a significant influence on friction and wear of ceramics and composites. Therefore, it is very important to understand the performance of engineering materials and to correlate them with their properties. Such correlation should be made in terms of microstructural, physical, electrical, mechanical, and (tribo)chemical properties under different tribological conditions. In the above perspective, Chapter 4 describes the physical characteristics of typical engineering surfaces, while Chapters 5 and 6 discuss the fundamentals of friction as well as origin/quantitative analysis of frictional heating, respectively. The phenomenological and mechanistic description of various wear mechanisms with a particular focus on fretting wear is provided in Chapter 7. The last chapter of this section briefly discusses the fundamentals of lubrication. Tribology is now widely accepted as “the science of interacting surfaces in relative motion and practices related there to.”1 Tribology embraces primarily the study of friction, wear, and lubrication and it is strongly an interdisciplinary field.2 As such, it is much broader in terms of areas that affect it, as well as its having a large effect on many other areas in engineering and sciences. The word “tribology” is derived from the Greek word tribos, which means rubbing.3 It is interesting to observe from a tribologist’s perspective that, although tribology seems to be a fresh field, having enormous potential for fundamental research, the history of this important branch of science has its roots in the early ages of humans.4 Using the friction
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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4â•…
CHAPTER 1â•… Introduction
Friction at tribological interface Wear of mating solids
Variables/system parameters
Fundamental disciplines
Interactions
Mating materials and surface properties
Metallurgy, materials science and engineering Physics and chemistry
Tribological environment, operating parameters, mechanics at interface
Mechanical engineering
Figure 1.1â•… Concept triangle illustrating the interaction of basic science and engineering disciplines and multiple parameters, as involved in the science of tribology.
between wood and/or stones in inventing fire is considered as the first utilization of the concept of tribology, and it belongs to Stone Age humans. The other important instances of tribology during early ages seem to be in making drills, bearings, sledges for transporting heavy loads, and so on. In particular, the Egyptian civilization has a record of understanding the significance of friction and wear as well as lubrication during construction of giant pyramids. As illustrated in Figure 1.1, the science of tribology can be explained based on synergistic interaction among the concepts and ideas drawn from fundamentals of physics and chemistry, as well as metallurgy, materials science, engineering, and mechanical engineering. It may be noted here that conventional metallurgical engineering principles are commonly used to tailor surface properties, for example, case hardening or nitriding of steels to improve wear resistance. On the other hand, various new materials (e.g., ceramics and polymers as well as their composites) have been developed in last few decades using the processing–structure–property correlation, a fundamental concept used in materials engineering. Many of these materials are now considered as potential replacements for traditional metallic materials. In mechanical engineering, multiple research groups actively contributed to the lubrication aspect of tribology. This has major relevance for ball bearings and various lubricated mechanical joints and bearings. A number of textbooks have dealt largely with the lubrication aspect. In this book, a great emphasis has therefore been placed on obtaining understanding based on the materials science aspect. Overall, the response of any tribocouple depends on the surface properties of two mating materials as well as on the tribological interaction with the environment; such interaction also critically depends on the mechanics at the tribological interface. This will result
â•… 5
CHAPTER 1â•… Introduction
in friction and wear damage of both the mating solids. Therefore, it is important to understand how two mating materials will respond mechanically at a loaded contact experiencing relative motion; equally important is the interaction of the environment and lubrication with the mating couple. These are explained in various chapters throughout this book. Friction, under nonlubricated conditions, is considered as the resistance to motion that arises from the solid surface interactions at the real area of contact.5–7 On the other hand, in the presence of lubricants, their viscous flow components and solid–liquid interactions play a major or a key role. Having low friction is important for certain applications such as hinges, rivets, bearings, and human hip joints, but applications such as brakes, clutches, and tires on roads in contrast require much higher or even “high” friction. In any case, the ability to control the optimal friction for a particular application is the goal of every designer; this ability depends on many of the aforementioned parameters and, in great part, on understanding and the performance of materials and their surface properties. The progressive loss of material due to the tribological interactions at contacting interfaces under relative motion is considered as the definition of the term “wear.” Wear may arise from the contact and relative motion of the solid body against a mating solid, but it often also includes a liquid or gaseous counterbody (water jet, air bubble implosions). Thus, the conditions of wear include several forms (sliding, rolling, erosion, impact, cavitation, etc.) in various atmospheric conditions. Wear is detrimental in many engineering applications, leading to failure of the various components and, finally, requiring repair or replacement. If we consider a technical system that drives a great part of modern economies, such as a car or any other vehicle, we find that most expenses for maintenance and replacement are due to tribological issues and wear. For example, wear of tires, brakes, wheel joints, bearings, and windshield wipers, scratches on paint and glass, and replacement of oil and oil filters are major concerns of end users, and consequently of manufacturers as well. On the other hand, sometimes wear is desirable: high wear rates are required for efficient production in some surface finishing processes such as polishing and grinding. However, like friction, not only the amount, but also, primarily, the control and prediction of wear of materials in every application is important for appropriate use and maintenance and is thus the key to success of every tribological application. It is understood that the extent of friction and wear depends on the system in which the component is utilized.8 Moreover, friction and wear are not intrinsic material properties, but the tribological features need to be considered as an engineeringsystem dependent property.9 Such control over friction and wear can be achieved by proper design, fabrication, and loading of the components. While friction is a direct and momentary response of the tribological components under the contact conditions, wear also includes the loading history of these components. Therefore, understanding and considering materials’ nature and responses are critical and very important for tailoring the properties of the materials and achieving the desirable conditions of friction and wear.10 In this regard, the aim of the material scientist is to control the properties of a specific component by incorporating suitable changes in its composition at the microstructural level during the processing stage, keeping
6â•…
CHAPTER 1â•… Introduction
tribological requirements in mind. Such microstructural modification should be correlated with performance at different tribological conditions in order to understand, design, manufacture, and control conventional and novel materials in a specific application, and to use all their capabilities and advantages. This is particularly important for advanced and heavily engineered materials, such as composites, ceramics, and nanomaterials. In this book, we tend to present the key material properties that are available to a material scientist for designing a material and that have critical influences on the material’s tribological behavior. We also discuss the fundamentals of tribology and related key parameters, finally describing and presenting examples of the microstructural control of advanced ceramic and composite materials for optimal tribological performance. We focus on a specific segment of high-tech materials that have a great potential for use in many applications and that, primarily, give much freedom and possibility for future development and innovative solutions. Although the book mainly focuses on the materials perspective, we would like to stress that response and success of every material, no matter how well designed, will finally depend on the tribological system, and this should therefore be considered in the early stages of materials and system development.
REFERENCES ╇ 1â•… B. Bhushan. Principles and Applications of Tribology. A Wiley-Interscience Publication, John Wiley & Sons, New York, 1999. ╇ 2â•… I. M. Hutchings. Tribology: Friction and Wear of Engineering Materials. Butterworth-Heinemann Publications, Guernsey, UK, 1992. ╇ 3â•… A. D. Sarkar. Friction and Wear. Academic Press, London, 1980. ╇ 4â•… H. Czichos. Tribology. Elsevier, Amsterdam, 1978. ╇ 5â•… J. F. Archard. Contact and rubbing of flat surfaces. J. Appl. Physics. 24 (1953), 981–988. ╇ 6â•… K. L. Johnson. Contact Mechanics, Vol. 26. Cambridge University Press, London, New York, Sydney, 1985, 230. ╇ 7â•… N. P. Suh. Tribophysics. Prentice-Hall, Englewood Cliffs, NJ, 1986. ╇ 8â•… G. W. Stachowaik and A. W. Batchelor. Engineering Tribology. Tribology Series, Vol. 24. Elsevier, Amsterdam, 1993. ╇ 9â•… K.-H. Zum Gahr. Microstructure and Wear of Materials. Elsevier, Oxford, 1987. 10â•… E. Rabinowicz. Friction and Wear of Materials, 2nd ed. Wiley, New York, 1995.
CHAPTER
2
OVERVIEW: TRIBOLOGICAL MATERIALS This chapter discusses the need for development of new materials in the context of tribological applications, followed by general classification of materials. Particular reference is made to the properties and applications of ceramics and their composites.
2.1 INTRODUCTION Materials have dominated technological development, and their importance has been recognized in all industrialized countries. The driving forces behind the development of “advanced materials” are various technological, socioeconomic, and environmental requirements, including the following: 1. improved performance, integrity, and reliability of engineering systems; 2. higher durability of products; 3. higher efficiency and lower energy consumption of engineering systems; 4. lightweight and high-strength structures. There is a general consensus that engineering materials can be classified into three primary classes: metals and alloys, ceramics and glasses, and polymers.1–6 Among these three primary classes, metals, metallic alloys, and polymers are, by far, more widely used than ceramics and glass for various structural and engineering applications. The widespread use of metallic materials is driven by their high tensile strength, high toughness (crack growth resistance), and ability to be manufactured in various sizes and shapes using reproducible fabrication techniques. Similarly, polymers have distinct advantages in terms of their low density, high flexibility, and availability in different shapes and sizes. Nevertheless, polymeric materials have low melting point (less than 400°C) as well as very low strength and elastic modulus. Compared with ceramics, metals have much lower hardness, and many commonly used metallic materials have much lower melting point. From this perspective, ceramics and glasses have advantageous properties, including refractoriness
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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CHAPTER 2â•… Overview: Tribological Materials
(capability to withstand high temperatures), strength retention at high temperature, high melting point, and good mechanical properties (hardness, elastic modulus, and compressive strength). In view of such an attractive combination of properties, ceramics are considered as potential materials for high-temperature structural applications and various tribological applications requiring high hardness and wear resistance. To combine various advantageous properties of the three primary material classes, a derived material class (i.e., composites) is being developed. The composites are generally defined as a class of materials that comprise at least two intimately bonded microstructural phases aimed to tailor properties (e.g., elastic modulus, hardness, strength) for specific applications; it is expected that a specific property of a composite should be higher than the simple addition of that property of the constituent phases. Depending on whether metals, ceramics, or polymers comprise more than 50% by volume of a composite, it can be further classified as a metal matrix composite (MMC), a ceramic matrix composite (CMC), or a polymer matrix composite (PMC). From the microstructural point of view, a composite contains a matrix (metal, ceramic, polymer) and a reinforcement phase. The crystalline matrix phase can have an equiaxed or elongated grain structure; the reinforcement phase can have different shapes, for example, particulates, whiskers, and fibers. The reinforcement shapes can be distinguished in terms of aspect ratio: particulates can be spheroidal; whiskers have an aspect ratio greater than 2; fibers have the largest aspect ratio. It is widely recognized now that the use of fibers or whiskers can lead to composites with anisotropic properties (different properties in different directions). As far as nomenclature is concerned, it is a common practice to designate a composite as M-Rp, M-Rw, or M-Rf, where M and R are the matrix and reinforcement, respectively, and the subscripts (p, w, f) essentially indicate the presence of reinforcement as particulates, whiskers, or fibers, respectively. One widely investigated MMC is AlSiCp composite. Also, Mg-SiCp is being developed as a lightweight composite. Several MMCs are used for automotive parts and structural components. Popular examples of CMCs include Al2O3–ZrO2â•›p, Al2O3–SiCw. These CMCs are typically used as wear parts and cutting tool inserts. Various resin-bonded PMCs are used for aerospace applications. While this book largely focuses on the tribological properties of ceramics and CMCs, a few chapters also deal with the friction and wear behavior of MMCs and PMCs that have ceramic reinforcements.
2.2 DEFINITION AND CLASSIFICATION OF CERAMICS A proper and exact definition of ceramics is very difficult. In general, ceramics can be defined as a class of inorganic nonmetallic materials that can be either processed or used at high temperatures and have ionic and/or covalent bonding.6 To the common person in society, the word “ceramic” means a coffee cup, sanitary tiles, and so on, which are traditional ceramic products. Although the major use of ceramics in last few decades was centered on applications such as construction materials, tableware, and sanitary ware, the advancement of ceramic science and technology
2.3 Properties of Structural Ceramics
â•… 9
since the early 1990s has enabled the application of this class of materials to extend from more traditional applications to such cutting-edge technologies as aerospace, nuclear, electronics, and biomedical, among others.1–6 In fact, ceramics are classified as traditional ceramics and engineering ceramics in many textbooks. Traditional ceramics are largely silica or clay based and typically involve low-cost fabrication processes. On the other hand, engineering ceramics are processed from high-purity ceramic powders and their properties can be tailored by varying process parameters and, thereby, microstructures. Also, engineering ceramics are, by far, more expensive than traditional ceramics. A large cross section of people in the developing world is slowly getting to know and realize the applications of engineering ceramics.7 To this end, this book sheds light on how various structural ceramics can be useful in tribological applications. Based on applications, engineering ceramics are usually categorized into two major classes: structural ceramics and functional ceramics. While the development of structural ceramics is mainly driven by the optimization of mechanical strength, hardness, and toughness,3 the performance of functional ceramics is determined by electric, magnetic, dielectric, optical, and other properties. In the ceramics community, structural ceramics can be further classified as (1) oxide ceramics (Al2O3, ZrO2, SiO2, etc.) and (2) non-oxide ceramics (SiC, TiC, B4C, TiB2, Si3N4, TiN, etc.). It needs to be categorically mentioned here that friction and wear properties determine the performance of structural ceramics in various engineering applications. This aspect is discussed in various chapters of this book.
2.3 PROPERTIES OF STRUCTURAL CERAMICS In general, ceramics have many useful properties, such as high hardness, stiffness, and elastic modulus, wear resistance, high strength retention at elevated temperatures, and corrosion resistance associated with chemical inertness.8 These are highlighted in Figure 2.1; the relative comparison among metals, ceramics, and polymers of various aspects in the context of tribological applications is summarized in Tables 2.1 and 2.2. Also, all the relevant mechanical properties are introduced in the next chapter. As is discussed in some of the subsequent chapters, although good hardness is required for resistance against abrasive and adhesive wear, higher elastic modulus is necessary for better resistance against Hertzian contact damage. From Table 2.1, it is clear that many engineering ceramics have better hardness and elastic modulus than steels, and, therefore, they are expected to have better abrasive and adhesive wear resistance. Also, it is known that ceramics can be used at very high temperature (more than 1000°C), where any other material class cannot be used. As far as maximum operating temperature is concerned, Ni-based superalloys are typically used at 1000°C. In contrast, some nitride and some oxide ceramics can be used at temperatures close to 1500°C. In addition to high melting point, ceramics have an advantage over metals in terms of high-temperature strength. High melting point can be very useful in terms of two aspects: (1) high contact temperatures, which are generated at a tribological contact, depending on the friction and operating parameters; and (2) high-temperature machining and tribological
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Excellent hardness and elastic modulus (better contact damage resistance)
High compressive strength (relevant for applications requiring high load at tribocontact)
Structural ceramics Low density (better specific properties than many metals)
High melting and mechanical property retention at elevated temperature (high temperature tribological applications)
Figure 2.1â•… Various properties of relevance to tribological applications of structural ceramics. TABLE 2.1â•… Physical, Mechanical, and Thermal Properties of Some Important Metallic, Polymer, and Ceramic Materials, Which Are Relevant for Various Tribological Applications8
Material
Density (ρ), gm/cm3
Elastic modulus (E), GPa
Fracture toughness (KIc), MPam1/2
Vickers hardness (HV), GPa
Thermal conductivity (K), W/m/K
Steel Cast iron Al-alloy Al2O3 ZrO2 Si3N4 SiC Polyamide (PA) Polyimide (PI) Polytetrafluoroethylene (PTFE) High-density polyethylene (HDPE)
7.8–7.9 7.1–7.4 2.6–2.9 3.9 5.6–6.25 3.2 3.2 1.1–1.14 1.3 2.1–2.3
210 64–181 60–80 210–390 140–210 170 450 2–4 3–5 0.4
50–214 6–20 23–45 3–5 8–10 4–7 4.5 3 —
1–9 1–8.5 0.25–1.40 14–19 12 16–18 25 0.8–1 — 0.12
30–60 30–60 121–237 25–35 2 25–50 90–125 0.25–0.35 0.37–0.52 0.25
0.92
0.2
0.13
0.33–0.57
1–2
TABLE 2.2â•… A Comparison among Three Primary Material Classes in Terms of Relevant Aspects in Relation to Tribological Applications10
Mass forces (F) Hertzian pressures (P) Friction-induced temperature increase (T ) Adhesion energy (Ad) Abrasion (Ab) Tribochemical reactivity (R)
Fpolymer╯<╛Fceramics╯<╯Fmetal Ppolymer╯<╛Pmetal╯<╯Pceramics Tmetal╯<╯Tpolymer╯<╯Tceramics Adpolymer╯<╯Admetal╯<╯Adceramics Abceramics╯<╯Abmetal╯<╯Abpolymer Rpolymer, Rceramics╯<╯Rmetal
2.4 Applications of Structural Ceramics
â•… 11
applications. As far as frictional heating is concerned, many ceramics, such as borides, have good thermal conductivity, and this enables easier heat dissipation through ceramics at tribological contacts. In addition, ceramics have lower density than many metals (see Table 2.1). Although polymers have the lowest density, many ceramics (alumina, SiC) have half the density of steel-based materials. Thus, highspeed turning or cutting operations are possible with ceramic-based tool inserts. The low density of ceramics can therefore be used in tribological applications requiring lightweight components as well as high specific properties (specific modulus, specific strength). It may be mentioned here that “specific properties” are properties normalized with respect to density. It is also known that, although ceramics are inferior in tensile properties, the compressive strength of structural ceramics is eight times larger than the tensile strength. The tribological contact experiences large compressive stresses at both surface and subsurface regions (see Chapter 4). Therefore, the high compressive strength of ceramics, mostly greater than 1â•›GPa, would be useful in applications involving high load at the tribological contact. At this point, it may also be mentioned that the compressive strength of ceramics is largely superior to that of metals. Many of the advantageous properties of ceramics, as explained previously, are utilized in many of the key technological applications of ceramics, which include engine parts, rocket nozzles, bioceramics for medical implants, heat-resistant tiles for the space shuttle, nuclear materials, and storage and renewable-energy devices. Despite having many attractive properties, as mentioned previously, the major limitation of ceramics for structural and some nonstructural applications is poor fracture toughness (see Table 2.1). Over the years, it has been realized that an optimum combination of high toughness with high hardness and strength is required for the majority of the current and future applications of structural ceramics. To address this need, the development of ceramic composites with optimal combination of mechanical properties has been the major focus in the ceramics community.6
2.4 APPLICATIONS OF STRUCTURAL CERAMICS There is a tremendous industrial need for new tribological materials. This need is realized in metal-forming industries, bearings, gears, valve guides and tappets in engines, seals and bearings involving fluid and gas transport often under corrosive conditions, and so on. The majority of these applications are currently served by hardened steels and WC-based hard metals with or without surface coatings.9 However, new materials or improved existing materials are needed to meet the increasing demand in the tribology world. It can be reiterated that ceramics, because of their ionic and/or covalent bonding, have a useful combination of physicomechanical properties (elastic modulus, hardness, and strength) and corrosion resistance. In many structural and tribological applications, ceramics are recognized as having great potential to replace existing materials for a series of rubbing pairs, such as seal rings, valve seats, extrusion dies, cutting tools, bearings, and cylinder liners.10 The materials of interest will have to combine high hardness, toughness, strength, elastic modulus, and wear resistance coupled with relatively low density, resulting
12â•…
CHAPTER 2â•… Overview: Tribological Materials
in low inertia under reciprocating stresses. Furthermore, the fundamental understanding of the relationships among composition, microstructure, processing route, mechanical properties, wear behavior, and performance should be clarified for optimal use of engineered materials in tribological applications. The development of new tribological materials is directed along two avenues: (1) coatings on conventional metallic substrates and (2) monolithic ceramics and ceramic composites. Of course, for particular low-load and low-temperature conditions, polymers and polymer composites have found some niche applications nicely suited for these materials. Coatings are mostly fabricated from hard carbides, nitrides, or borides, with development of diamond or diamond-like (C–H) films at the higher end of the hardness–cost scale.11 Coating thickness is normally between 1 and 50â•›µm, depending on the deposition process (physical vapor deposition [PVD], chemical vapor deposition [CVD], or electrolytic). This presents limitations in terms of lifetime or property influence of the relatively soft substrate. Diamond-like carbon (DLC) coatings, in particular, have been improved tremendously in processing technology and performance, so now they can provide low wear and low friction for many applications, including lubricated and nonlubricated conditions.12–19 Thicker coatings may be applied by thermal spraying (in the millimeter range) but are limited in chemistry, compatibility with substrate properties (thermal expansion, etc.), and cohesion. Monolithic ceramics, especially those with improved strength and toughness, have been a focus of development in different research labs and industries since the 1970s.20 However, monolithic ceramics are not optimum for all engineering applications. Ceramic composites such as metal matrix and PMCs are now the established approaches to designing structural materials.21,22 Ceramic reinforcements are commercially available in different forms such as whiskers, platelets, particulates, and fibers. Two major classes of ceramic composites are fiber-reinforced and particle/ whisker-reinforced ceramic composites. A popular example of the first class is silicon carbide fiber-reinforced glass–ceramics. Four major drawbacks normally restrict the widespread use of this material class for structural applications: high cost of ceramic fibers, expensive composite production route, chemical compatibility of the fiber with the matrix, and oxidation of SiC fibers at high temperatures. To this end, particle-reinforced CMCs offer a viable and relatively cost-effective option for development of materials with improved and optimal combinations of mechanical properties (hardness, toughness, and strength). Many of the chapters in this book deal with the tribological properties of particulate-reinforced CMCs. In the world of ceramic materials, yttria-doped zirconia ceramics, in particular yttria-stabilized tetragonal zirconia polycrystalline (Y-TZP) ceramics, are regarded as a strong candidate for structural applications due to the excellent combination of strength (≈700–1200â•›MPa) and fracture toughness (≈2–20â•›MPaâ•›m1/2) in combination with good chemical inertness.23,24 The high toughness of the zirconia monoliths stems from the stress-induced transformation of the tetragonal (t) phase to monoclinic (m) in the stress field of propagating cracks, a concept widely known as transformation toughening.25 Basic microstructural requirements for the effective contribution from transformation toughening is the maximum retention of the tetragonal phase at the application temperature with sufficient transformability to m-ZrO2
2.4 Applications of Structural Ceramics
â•… 13
in the crack tip stress field. Since the discovery of the transformation toughening concept some two decades ago,26 this approach has been successfully utilized to toughen several intermetallic,27 glass,28 and ceramic29 microstructures. Extensive efforts have also been made to increase the toughness of alumina by adding zirconia, a class of materials known as zirconia-toughened alumina (ZTA). Among various non-oxide ceramics, the borides of transition metals, such as TiB2, because of their high melting point, high hardness, electrical and thermal conductivity, and high wear resistance, are used for a variety of technological applications.30 Monolithic TiB2, that is, without any second phase addition, has excellent hardness (≈25â•›GPa at room temperature), good thermal conductivity (≈64â•›W/m/°C), high electrical conductivity (electrical resistivityâ•›≈â•›13╯×╯10−8â•›Ωm), and considerable chemical stability.31 Some of these attractive properties are ideally suited to be exploited for tribological applications. However, the relatively low fracture toughness (≈5â•›MPaâ•›m1/2) and modest bending strength (≈500â•›MPa) coupled with poor sinterability of monolithic TiB2 limits its use in many engineering applications.32 In the materials world, TiB2 is often used as the reinforcement phase not only for ceramics, but also for metallic alloys such as Al-alloys33,34 to develop composites with improved abrasive wear resistance. Furthermore, TiB2 as well as TiCN-, TiN-, or TiC-containing materials are known as electroconductive materials and can be shaped by electrodischarge machining (EDM) to manufacture complex components, greatly increasing the number of industrial applications of these ceramic materials.35,36 In Figures 2.2 and 2.3, the use of ceramics in various engineering applications requiring wear resistance is shown. One of the major uses of ceramics and cermets is cutting tool inserts (see Fig. 2.2a). Another application of ceramics, which attracted wider attention, is ball bearings (see Fig. 2.2b). Ceramic balls enclosed in a steel race, that is, hybrid bearings, are now used in turbopumps of the space shuttle main engine. The friction and wear properties of alumina, zirconia, and SiC in cryogenic environments are being investigated as such studies are relevant to cryotribological applications.37–39 The lower density of ceramics, compared with metals, is useful for high speed bearing applications. Ceramic ball bearings are commercially available with diameter from 4â•›mm to as large as 20–30â•›mm, and they are made from silicon nitride or sialon (Si6−zAlzOzN8−z, with z being the substitution level). Some non-oxide ceramics, such as SiC, are used as seal materials in automotive applications, particularly in engines requiring high heat resistance (see Fig. 2.2c). Ceramics are also used in many metal-forming applications. For example, Si3N4-based materials are used as dies or other tools for wire-drawing applications (see Fig. 2.2d). In addition, ceramics because of their good wear resistance are also used in paper-cutting industries. Another emerging area where ceramics are slowly penetrating is biomedical applications. In total hip joint replacement (THR), ceramic balls (Al2O3, ZrO2) are used as femoral ball heads, which make tribological contact with the acetabular cup, made from polymeric materials (see Fig. 2.3a). Also, hydroxyapatite (HAp)-based bioceramics are used as bone spacers. In endoscopic surgery, various ceramic tools are used (see Fig. 2.3b). Apart from biocompatibility, wear resistance is an important property in all such applications.
14â•…
CHAPTER 2â•… Overview: Tribological Materials
(a)
(c)
(b)
(d)
Figure 2.2â•… Some representative applications requiring good tribological properties: (a) ceramic-based cutting tool insert materials and metal holders used in various conventional machining applications; (b) Si3N4-based ball bearings and wear-resistant components; (c) SiC-based ceramic seal materials; and (d) Si3N4-based tools for wiredrawing applications.7
2.5 CLOSING REMARKS It must be borne in mind that, in tribological applications, materials-related influences must be seen in a broader context, because these materials are components of “tribological systems.” Materials for tribological applications span all material classes, including metals, ceramics, and polymers and their composites. Importantly, these materials differ considerably in their physicomechanical properties, as can be seen in Table 2.1.8 For example, metallic materials are characterized by high values
2.5 Closing Remarks
â•… 15
(a)
(b)
Figure 2.3â•… Some illustrative examples to show the use of ceramics in biomedical applications: (a) total hip joint replacement (THR) and total knee joint replacement (TKR), bone spacers; (b) various ceramic parts used in medical surgeries (endoscopic surgery, etc.).7
of tensile strength, fracture toughness, and thermal conductivity. Outstanding properties of ceramics include high elastic moduli, compressive strength, and hardness, which decrease a little with increasing temperature. However, their disadvantage is low fracture toughness. A positive aspect for polymers is their low density, but negative factors are their low thermal durability, hardness, and strength. Friction and wear properties cannot be directly correlated with the bulk properties of a material, because the tribological properties are determined by surface composition and properties. Moreover, interfacial properties affected by environment and tribochemical reactions are often the key parameters. Being system dependent, these properties, including both mating materials, influence the tribological properties of any contact and component considerably (Table 2.1). Among all
16â•…
CHAPTER 2â•… Overview: Tribological Materials
material classes, ceramic materials have an added advantage, when compared with metals (such as steels), with reference to their low mass forces, better abrasion resistance, and moderate tribochemical behavior. Less positive aspects of ceramics include higher modulus-dependent contact pressures, the shift of the Hertzian shear stress maximum from bulk to the surface, and the high friction-induced temperature increase, because of relatively low thermal conductivity (alumina and zirconia). Polymers, compared with metals, are beneficial due to their low interfacial adhesion energy, and such a property leads to lower friction in polytetrafluoroethylene (PTFE) and polyethylene (PE) systems. Also, polymers can only be used under low Hertzian contact pressures (a few orders of megapascal), due to their premature failure by viscoelastic and plastic deformations at higher loads.4 Some of these aspects will be discussed in the following chapters. It is known that the influence of materials on the behavior of tribological systems is most pronounced or severe in unlubricated, that is, “dry” operating conditions; it is therefore important to review the unlubricated tribological properties of different classes of materials. Additionally, the tribological properties, as recorded under different environments (water, cryogenic, etc.), are also of interest for real-life applications. All these aspects, for a variety of ceramics and composites, are discussed in several chapters of this book.
REFERENCES ╇ 1â•… H. Yanagida, K. Koumoto, and M. Miyayama. The Chemistry of Ceramics. John Wiley & Sons, New York, 1995. ╇ 2â•… M. W. Barsoum. Fundamentals of Ceramics. Taylor & Francis, 2003. ╇ 3â•… C. B. Carter and M. G. Norton. Ceramic Materials. Springer, New York, 2007. ╇ 4â•… Y. M. Chiang, D. P. Birnie, and W. D. Kingery. Physical Ceramics. John Wiley & Sons, New York, 1997. ╇ 5â•… K. H. Zum Gahr. Sliding wear of ceramic-ceramic, ceramic-steel and steel-steel pairs in lubricated and unlubricated contact. Wear 133 (1989), 1–22. ╇ 6â•… W. D. Kingery, H. K. Bowen, and D. R. Uhlmann. Introduction to Ceramics. John Wiley & Sons, 2nd ed. New York, 1976. ╇ 7â•… D. W. Richerson. Magic of Ceramics. Wiley-American Ceramic Society, 2000. ╇ 8â•… H. Czichos and K.-H. Habig. In Tribology Handbook Friction and Wear. Vieweg Verlag, Braunschweig, 1992, p. 560. ╇ 9â•… A. Mukhopadhyay and B. Basu. Recent development of WC-based cermets and nanocomposites. J. Mater. Sci. 46 (2011), 571–589. 10â•… H. Czichos, D. Klaffke, E. Santner, and M. Woydt. Advances in tribology: The materials point of view. Wear 190(2) (1995), 155–161. 11â•… B. Subramonian, K. Kato, K. Adachi, and B. Basu. Experimental evaluation of friction and wear properties of solid lubricant coatings on SUS440C steel in liquid nitrogen. Tribology Lett. 20(3–4) (2005), 263–272. 12â•… A. Matthews, A. Leyland, K. Holmberg, and H. Ronkainen. Design aspects for advanced tribological surface coatings. Surf. Coat. Technol. 100–101 (1998), 1–6. 13â•… A. Grill. Diamond-like carbon: State of the art. Diam. Relat. Mater. 8 (1999), 428–434. 14â•… J. Fontaine, C. Donnet, A. Grill, and T. L. Mogne. Tribochemistry between hydrogen and diamondlike carbon films. Surf. Coat. Technol. 146–147 (2001), 286. 15â•… A. Erdemir and C. Donnet. Tribology of diamond-like carbon films: Recent progress and future prospects. J. Phys. D Appl. Phys. 39(18) (2006), 311–327.
REFERENCES
â•… 17
16â•… A. Neville, A. Morina, T. Haque, and M. Voong. Compatibility between tribological surfaces and lubricant additives—How friction and wear reduction can be controlled by surface/lubes synergies. Tribology Int. 40(10–12) (2007), 1680–1695. 17â•… M. Kalin, I. Velkavrh, and J. Vizintin. Review of boundary lubrication mechanisms of DLC coatings used in mechanical applications. Meccanica 43 (2008), 623–637. 18â•… I. Velkavrh, M. Kalin, and J. Vizintin. The performance and mechanisms of DLC-coated surfaces in contact with steel in boundary-lubrication conditions—A review. Strojniski Vestn.—J. Mech. Eng. 54(3) (2008), 189–206. 19â•… M. Kalin, I. Velkavrh, and J. Vizintin. The Stribeck curve and lubrication design for non-fully wetted surfaces. Wear 267 (2009), 1232–1240. 20â•… J. D. Cawley and W. E. Lee. Oxide ceramics, in Structure and Properties of Ceramics, Materials Science and Technology, Vol. 11, R. W. Cahn, P. Haasen, and E. J. Kramer (Eds.). VCH, Weinheim, Germany, 1994, 101–114. 21â•… M. Rühle and A. G. Evans. High toughness ceramics and ceramic composites. Prog. Mater. Sci. 33 (1989), 85. 22â•… B. Basu. Zirconia–titanium boride composites for tribological applications. PhD thesis, Katholieke Universiteit Leuven, Belgium, March, 2001. 23â•… B. Basu. Toughening of Y-stabilized tetragonal zirconia ceramics. Int. Mater. Rev. 50(4) (2005), 239–256. 24â•… P. F. Becher and M. V. Swain. Grain size dependent transformation behavior in polycrystalline tetragonal zirconia ceramics. J. Am. Ceram. Soc. 75 (1992), 493. (b) J. B. Wachtman. Mechanical Properties of Ceramics. John Wiley & Sons, New York, 1996, 391–408. 25â•… D. J. Green, R. H. J. Hannink, and M. V. Swain. Microstructure-mechanical behavior of partially stabilised zirconia (PSZ) materials, Chapter 5, in Transformation Toughening of Ceramics, CRC Press, Boca Raton, FL, 1989, 157–197. 26â•… R. C. Garvie, R. H. J. Hannink, and R. T. Pascoe. Ceramic steel? Nature 258 (1975), 703–704. 27â•… B. Basu, J. Vleugels, and O. Van Der Biest. Processing and mechanical properties of ZrO2-TiB2 composites. J. Eur. Ceram. Soc. 25 (2005), 3629–3637. 28â•… T. Höche, M. Deckwerth, and C. Rüssel. Partial stabilisation of tetragonal zirconia in oxynitride glass-ceramics. J. Am. Ceram. Soc. 81(8) (1998), 2029–2036. 29â•… B. Basu, J. Vleugels, and O. Van Der Biest. Toughness tailoring of yttria-doped zirconia ceramics. Mater. Sci. Eng. A 380 (2004), 215–221. 30â•… R. A. Cutler. Engineering Properties of Borides, in Engineered Materials Handbook, Vol. 4, Ceramics and Glasses. ASM International, The Materials Information Society, Materials Park, OH, 1991. 31â•… J. Vleugels, B. Basu, K. C. H. Kumar, R. G. Vitchev, and O. Van Der Biest. Unlubricated fretting wear of TiB2 containing composites against bearing steel. Metallurgical Mater. Trans. A 33(12) (2002), 3847–3859. 32â•… G. Brahma Raju and B. Basu. Densification, sintering reactions, and properties of titanium diboride with titanium disilicide as a sintering aid. J. Am. Ceram. Soc. 90(11) (2007), 3415–3423. 33â•… D. Roy, S. S. Singh, B. Basu, W. Lojkowski, R. Mitra, and I. Manna. Studies on wear behavior of nano-intermetallic reinforced Al-base amorphous/nanocrystalline matrix in-situ composite. Wear 266 (2009), 1113–1118. 34â•… D. Roy, S. Ghosh, A. Basumallick, and B. Basu. Preparation of Fe-aluminide reinforced in situ metal matrix composites by reactive hot pressing. Mater. Sci. Eng. A 415 (2006), 202. 35â•… B. V. Manoj Kumar, J. Ramkumar, B. Basu, and S. Kang. Electro-discharge machining performance of TiCN-based cermets. Int. J. Refract. Metals Hard Mater. 25 (2007), 293. 36â•… A. Bellosi, G. De Portu, and S. Guicciardi. Preparation and properties of electroconductive Al2O3based composites. J. Eur. Ceram. Soc. 10 (1992), 307–315. 37â•… T. Kumar Guha and B. Basu. Microfracture and limited tribochemical wear of silicon carbide during high speed sliding in cryogenic environment. J. Am. Ceram. Soc. 93(6) (2010), 1764–1773. 38â•… R. Khanna and B. Basu. Sliding wear properties of self-mated yttria-stabilised tetragonal zirconia ceramics in cryogenic environment. J. Am. Ceram. Soc. 90(8) (2007), 2525–2534. 39â•… R. Khanna and B. Basu. Low friction and severe wear of alumina in cryogenic environment: A first report. J. Mater. Res. 21(4) (2006), 832–843.
CH A P T E R
3
OVERVIEW: MECHANICAL PROPERTIES OF CERAMICS Ceramics are generally known for having a good combination of high hardness, high compressive strength, high elastic modulus, and low fracture toughness. Along with the description of these properties, the theory of brittle fracture as well as an understanding of indentation-induced cracking is required to realize tribological properties of ceramics. From this perspective, this chapter discusses the physics of brittle fracture as well as the principles of various mechanical properties of ceramics, with particular reference to toughness measurement techniques. Toward this end, various toughening mechanisms are also mentioned.
3.1 THEORY OF BRITTLE FRACTURE There is a general consensus that ceramics are possibly best known for their brittleness rather than for some of their outstanding properties, which include the strength retention of some ceramics (SiC, Si3N4) at high temperature and the extremely good biocompatibility and bioactivity of others (hydroxyapatite). In this section, some of the theories of brittle fracture of ceramics and glasses are reviewed, and the concepts of toughening mechanisms are discussed. During the early stage of the development of understanding of brittle fracture, Inglis considered the aspect of the stress field at a crack tip and proposed a theory based on stress concentration at the crack tip.1 For an elliptical hole with major axis c and minor axis b in a solid continuum under externally applied tensile stress σ, the stress at both edges of the cavity can be expressed as
σ max = σ(1 + 2c / b).
(3.1)
It is important to note that, at the tip or edge of a cavity, a tensile stress field is realized, whereas on the upper, relatively flat surface of a crack, compressive stress is realized. For câ•›>>â•›b, a hole can be realistically imagined as a crack with crack tip radius of curvature ρ at the crack tip (see Fig. 3.1) given by ρ╯=╯b2/c. Accordingly, Equation 3.1 can be rewritten as Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
18
3.1 Theory of Brittle Fracture
â•… 19
σ
c
σ
(
)
σmax = 1 + 2 ρc σ ≈ 2 ρc σ σmax Radius of crack tip
for c >> ρ
ρ σmax
Figure 3.1â•… Schematic illustration showing the stress concentration at the crack tip edge.
c σ max = σ 1 + 2 . ρ
(3.2a)
From geometrical considerations, it is evident that aâ•›>>>â•›ρ, and, therefore, the maximum stress at the crack tip edge can be expressed as
σ max = 2σ
c . ρ
(3.2b)
From Equation 3.2b, the maximum stress at the crack tip will be much higher than the nominal stress (σmaxâ•›>>â•›σ). According to Inglis, fracture can only occur when the stress at the crack tip is just sufficient to break interatomic bonds ahead of the crack tip. From first-principle considerations,2 the theoretical cohesive strength can be obtained from the maximum in total interatomic force between two atoms, that is, the stress required to rupture interatomic bonds,
Eγ σ th = a0
1/ 2
,
(3.3)
where E is the elastic modulus, γ is the surface energy, and a0 is the equilibrium interatomic distance under unstrained conditions.
20â•…
CHAPTER 3â•… Overview: Mechanical Properties of Ceramics
Combining Equations 3.2b with 3.3, one gets σ max = σ th , 2σ
c Eγ = ρ a0
1/ 2
,
Eγρ σc = 4 a0 c
(3.4)
1/ 2
.
Subsequently, Orowan observed that the crack tip radius of curvature is of approximately the same magnitude as the interatomic distance,3 that is, ρ╯=╯ao. This results in the following expression for the critical fracture stress: Eγ σ inglis = 4c
1/ 2
.
(3.5)
For ceramics and glasses, the most widely accepted theory of brittle fracture is Griffith’s theory,4 which is proposed on the basis of the total change in potential energy during crack propagation under external tensile loading. Considering a rectangular solid plate with a through-thickness central hole being loaded in tension, the total energy change for the system can be expressed as ∆u = − ∆uel + ∆us, where Δuel is the elastic strain energy release around the elliptical hole, and Δus is the change in surface energy as the cavity extends perpendicular to the tensile stress direction. Considering that the plate has a thickness t that is largely smaller than the width w, that is, tâ•›<<â•›w, Griffith proposed that elastic strain energy can be assumed to be released over an elliptical volume around the crack, with the major axis being twice the crack length and the minor axis being the crack length (see Fig. 3.2a). Using Hooke’s law of elasticity to estimate the change in elastic strain energy (negative quantity) and an additional positive contribution from the change in surface energy due to creation of two crack surfaces, the total change in potential energy of the cracked body can be expressed as
∆u =
−a2 −σ 2 π × 2c × c × t + 2c × t × 2 γ π × 2c 2 + 4cγ . 2E 2E
(3.6)
The variation of different energy terms in Equation 3.6 is schematically shown in Figure 3.2b. At the maximum of the Δu–c curve, the critical fracture stress or corresponding critical crack length (c*) can be defined as σc =
2 γE πc*
c* =
2 γE . πσ 2
or
3.1 Theory of Brittle Fracture
â•… 21
σ
2c
(a)
∆u (Energy)
∆us
∆u c* 0
∆uel c (Semicrack length) (b)
Figure 3.2â•… (a) Schematic illustrating a rectangular plate containing a through-thickness elliptically shaped cavity and (b) energetics involved in Griffith’s theory of brittle fracture.
22â•…
CHAPTER 3â•… Overview: Mechanical Properties of Ceramics
On the basis of Griffith’s theory, the fracture will occur if one, if not both, of the following conditions is satisfied: 1. external stress σâ•›≥â•›σc, 2. intrinsic flaw size câ•›≥â•›c*. Therefore, cracks of size less than c* cannot grow at a given external stress, as any infinitesimal crack growth will lead to an overall increase in energy of the system. When câ•›≥â•›c*, any infinitesimally small increase in crack length can lower the total energy along the downhill of the Δu–c curve. With an increase in external stress, the Δu–c curve will be shifted toward the left, and, therefore, c* will be correspondingly lowered. This implies that critical cracks of finer sizes can grow at higher stress, leading to early fracture. At this juncture, one aspect of crack growth behavior needs to be mentioned. In glasses and some ceramics (e.g., Si3N4), the cracks with sizes smaller than critical crack size (c╯<╯c*) can grow in an unstable manner at a given stress level in a moist or humid environment, leading to fracture.5,6 This is called subcritical crack growth and is attributed to environmental interaction, leading to chemical attack at the crack tip, leading to easy bond breakage. From the preceding discussion, it is evident that the fracture of brittle solids depends on a combination of σc and c*. The stress intensity factor is therefore defined as K = Yσ πc , where Y is a factor dependent on the location or orientation of the crack and on the loading condition. In classical fracture mechanics theory, three different modes of loading of the crack faces (see Fig. 3.3) are considered: tensile or crack opening mode (mode I), shear mode (mode II), and tearing mode (mode III). Correspondingly, the stress intensity factor can be defined as KI, KII, and KIII. Among three modes, the failure of brittle solids is largely due to mode I failure, and accordingly, the critical mode I crack tip stress intensity factor is defined as K Ic = Yσ c πac . The value of Y varies from around 1 to 1.1; KIc is considered as a parameter to describe fracture toughness.
Mode I Crack opening mode
Mode II Sliding mode
Mode III Tearing mode
Figure 3.3â•… Three modes of loading of crack faces. Mode I (tensile or crack opening mode) is the most widely observed fracture mode.6
3.2 Cracking in Brittle Materials
â•… 23
3.2 CRACKING IN BRITTLE MATERIALS As mentioned earlier, the growth of preexisting flaws limits the strength of brittle solids. It is therefore important to know the characteristics of cracking and how cracks develop during external loading. In the case of distributed loading (e.g., under spherical indents), cone cracks develop, as shown in Figure 3.4. During continuous loading with a blunt indenter, the cracks initially form at the indent edges and, thereafter, propagate at an angle downward. The entire crack pattern evolves to a conical geometry during unloading, as observed clearly in the case of (i)
+
+
(iv) P r
ct 2a
(ii)
(iii)
(v)
+
+
+
– (vi)
a
2ba c
(a)
15
0
5
10
15
20
Indenter load , P (kN)
Cone crack 10
5
0
0
5
10 15 3/2 (mm) Crack size, c3/2 I (b)
20
Figure 3.4â•… Schematic showing the development of cone crack during loading and unloading of a spherical blunt indenter against a flat surface of a brittle material and the geometrical parameter associated with cone crack configuration (a). The cone crack formation in soda lime glass is also shown in (a). The load-dependent crack growth is shown in (b).6
24â•…
CHAPTER 3â•… Overview: Mechanical Properties of Ceramics
transparent soda lime glass (see Fig. 3.4). A different cracking pattern (radial– median and lateral cracks) in brittle solids is reported under concentrated loading via a sharp indenter (e.g., Vickers indentation). During initial loading, the radial– median cracks develop from the highly deformed zone beneath the sharp indenter. These cracks grow with time during loading; during unloading, a lateral crack pattern develops along the transverse direction. With complete unloading, the fully developed lateral cracks intersect the free surface of the brittle solid (see Fig. 3.5). During an abrasion process, characterized by a harder solid sliding on a relatively softer surface, the mechanical interaction between two solids can be considered as equivalent to multiple overlapping indentation processes under concentrated loading. Therefore, multiple lateral cracks generated from overlapping indented regions will intersect each other before meeting free surfaces. Such physical phenomena can cause material removal, that is, wear, and this has been analytically modeled by Tiwari et al.7 Based on the fundamental fracture mechanics theory,6 Lawn observed the following relationship: P ∝ cl3 / 2 or P ∝ cc3 / 2 , where P is the indentation load and cl or cc is the lateral or cone crack length, respectively. Such a relationship was established after several experimental measurements of cracking in transparent soda lime glass in an inert environment, and it clearly indicates that the severity of mechanical wear damage will increase with increase in indent load. A detailed fracture-mechanics-based treatment of cracking in brittle solids is reported elsewhere.6
3.3 DEFINITION AND MEASUREMENT OF BASIC MECHANICAL PROPERTIES 3.3.1â•… Hardness Conventionally, the hardness of a material is defined as the resistance to permanent deformation. This property is measured by indenting flat polished surfaces. For most ceramics, the hardness value is obtained with the Vickers indentation technique using the following expression:
P H v = 1.854 2 , d
(3.7)
where Hv is Vickers hardness, P is the applied load, and d is the average value of two diagonals. Two aspects need to be considered when experimentally measuring hardness of engineering ceramics: 1. The indent load should be such that it does not cause cracking from indent corners or edges as well as a stable or well-developed indentation developed without any spalling or damage around the indentation.
3.3 Definition and Measurement of Basic Mechanical Properties
(i)
+
–
(iv)
(ii)
+
–
(v)
â•… 25
P
2c 2a
(iii)
(vi)
–
+
L
R (a)
0.6
Indenter load , P (kN)
Radial crack
0.4
0.2
0
0
0.5
1.0 1.5 3/2 (mm) Crack size, c3/2 I
2.0
(b)
Figure 3.5â•… (a) Schematic showing the development of radial–median (marked as R) and lateral cracks (marked as L) during loading and unloading of a sharp indenter against a flat surface of a brittle material, respectively (left) and the geometrical parameter illustrating the dimension of highly deformed zone and crack length (right). (b) Load-dependent crack growth.6
26â•…
CHAPTER 3â•… Overview: Mechanical Properties of Ceramics
2. The hardness of a new ceramic composition or of a ceramic processed using a new synthesis (sintering) route should be measured using varying indent loads. This can reveal any “indentation size effect” and a conservative estimate of “true hardness” can be obtained. To obtain a reliable measurement for hard ceramics, it is suggested to use electron microscopy to measure the indent diagonal (length scale in “micrometer order”) as any small error in measuring the diagonal length will lead to a large error in hardness (see also Eq. 3.7). To illustrate the extremely high hardness of ceramic materials, Figure 3.6 presents a comparison of hardness properties of various materials. Many of the ceramics have much higher hardness than all the refractory metals. In general, the hardness of ceramics varies in the range of 10–40â•›GPa. Just to realize how high such
800
WC CBN
Young’s modulus (GPa)
600
TiN TaN
TaC
400
W
CrB2
ZrN
TiC NbC SiC HfC ZrC B4C Cr3C2
TiB2
Al2O3
ZrB2
Mo
WSi2 TiSi2
Si3N4 TaB2
200
MoSi2
ZrSi2
ZrO2 TiO2
Ta
Nb
Oxides
Silicides
Borides
Nitrides
Carbides
Refractory metals
0
Ni- and Coalloys
SiO2
(a)
Figure 3.6â•… Comparison of various materials in terms of basic mechanical properties: (a) elastic modulus and (b) hardness.23
3.3 Definition and Measurement of Basic Mechanical Properties
â•… 27
CBN
4000
B 4C
Microhardness (kg/mm2)
3000
TiB2 TiC SiC H1C ZrC NbC TaC WC Cr3C2 W2 C Cr3C2
2000
Mo2C Cr23C6
H1B2
Si3N4 TiN H1N ZrN
ZrB2 CrB2 TaB2
TaN
1000
Al2O3
MoSi2 TaSi2 Cr3Si2
ZrO2 Cr2O3
WSi2
TiO2
TiSi2 SiO2
Oxides
Silicides
Borides
Nitrides
Carbides
Refractory metals
Ni- and Coalloys
0
W Mo Nb Ta
(b)
Figure 3.6â•… (Continued )
values are, the hardness of fully hardened martensitic steel is around 7â•›GPa. Most ceramics, such as zirconia, alumina, and SiC, are around two to three times harder than fully hardened steel. As is explained later, the high hardness of ceramics imparts good wear resistance in various engineering applications.
3.3.2â•… Compressive Strength Although ceramics have low tensile strength, they have a superior compression property. The difference can be ascribed to microstructural response to crack growth and the nature of crack propagation under two different loadings. The stress–strain response under tension and compression is compared in Figure 3.7. While ceramics behave like a perfectly linear elastic material up to fracture, they exhibit a nonlinear response after reaching peak load (much higher than that in tension) in compression. As opposed to tensile crack growth, the cracks tend to extend vertically along the
28â•…
CHAPTER 3â•… Overview: Mechanical Properties of Ceramics
Serrations during compressive failure
σc
Stress
Spall σt
Strain
Figure 3.7â•… Stress–strain behavior of a brittle ceramic during compression with σc indicating the compressive strength. For comparison, the tensile stress–strain plot is superimposed to illustrate significantly higher compressive strength than tensile strength (σt) of a brittle ceramic. The compression failure mechanisms are also shown.
loading direction in compression. The serration in compression stress–strain response is primarily due to spalling of a small test volume, as the growing cracks either coalesce with each other or meet an unconstrained surface. Clearly, a delayed fracture behavior is expected in compression, and typically, compressive strength is around eight times higher than tensile strength (see also Fig. 3.7). Compressive strength can be measured using a universal testing machine (UTM), which can be used for tensile, flexural, and compression testing. For measurement of compressive strength, the cylindrical test sample with height-to-diameter ratio of 1.0 or larger is normally used. The samples are placed between two parallel plates of the machine, and the sample is loaded with a constant crosshead speed (typically around 0.05â•›mm/s). During the entire compression test, the load– displacement response can be obtained using a computer attached to the UTM. The compressive strength (σcs) can be calculated from the fracture load using the simple formula
σ cs =
P , A
(3.8)
where P is the maximum load (fracture load) and A is the cross-sectional area.
3.3.3â•… Flexural Strength In view of the inherent difficulty in obtaining tensile test samples, the strength of the ceramics is alternatively measured under flexure mode either by three-point or by four-point bend configuration. For this purpose, either rectangular or circular cross-section samples are used in a bend fixture. Either a concentrated load is applied in a three-point configuration or a distributed load is applied at two different places in a four-point configuration (see Fig. 3.8). The flexural strength thereafter is calcu-
3.3 Definition and Measurement of Basic Mechanical Properties
â•… 29
P
L/2
L/2
P/2
P/2 (a) P/2
P/2
L/4
L/2
L/4 P/2
P/2 (b)
Figure 3.8â•… Schematic illustration of three-point (a) and four-point (b) flexural strength measurement of ceramic bar of rectangular cross section.
lated on the basis of measured fracture load, following the fundamental principle of mechanics of solids. For three-point loading, the fracture strength can be obtained using the following expression:
σf =
3PL , 2bd 2
(3.9)
where σf is the flexural strength of the material, P is the fracture load, L is the span length, b is the width of the sample, and d is the thickness of the specimen. Similarly, the flexural strength for a four-point bend configuration can be estimated using the expression
30â•…
CHAPTER 3â•… Overview: Mechanical Properties of Ceramics
σf =
3PL . 4bd 2
It is important to mention that, during flexural testing, the loading surface is placed in compression, while the opposite surface is placed under tension. Also, the stress linearly decreases along the thickness (z direction). For four-point loading, maximum tensile stress is distributed over a larger area of the sample as opposed to the three-point flexural mode; hence, a lower and conservative estimate of the strength is obtained in four-point flexural tests.
3.3.4â•… Elastic Modulus The elastic modulus is another important property in the context of tribological property evaluation, and this can be obtained using a dynamic elastic property analysis instrument, which measures the elastic modulus by impulse excitation of vibration. The rectangular samples are impacted upon by a steel ball near the highfrequency sensor of the instrument. The values of Young’s modulus can be estimated using the following formula8:
mf 2 L3 E = 0.9465 f 3 T , b t
(3.10)
t T = 1 + 6.858 . L
(3.11)
where
E is Young’s modulus, m is mass, ff is the natural frequency in the flexure dimension, b is width, t is thickness, and L is length. This technique is widely used in measuring elastic modulus properties of various structural ceramics.9,10 The elastic modulus (E) and Poisson’s ratio (ν) can also be determined using an ultrasonic method using lithium niobate crystals for reusing transmitting and receiving signals, which are typically generated at 10â•›MHz resonant frequency. The velocities of the longitudinal and shear waves can be calculated from the thickness of the specimen and the travel time of the waves across the thickness or height of specimen. The following relationships can be used to determine the modulus properties:
(1 + ν)(1 − 2 ν) (ρCL ), 1− ν 1 / 2(C1 / Cs )2 − 1 v= , (CL / Cs )2 − 1
E=
(3.12) (3.13)
where ν╯=╯Poisson’s ratio, ρ╯=╯density of specimen (g/cm3), CL╯=╯velocity of longitudinal wave (m/s), Cs╯=╯velocity of shear wave (m/s), and E╯=╯elastic modulus (GPa).
3.3 Definition and Measurement of Basic Mechanical Properties
â•… 31
In Figure 3.6a, the elastic modulus of ceramics is compared with various other metals. The modulus values of several engineering ceramics vary in the range of 200–600â•›GPa. While the Young’s modulus (often reported as elastic modulus) of some oxide ceramics (e.g., Al2O3) is comparable with that of refractory metal (e.g., W), considerably higher modulus can be obtained with many non-oxide ceramics. For example, Young’s modulus of more than 600â•›GPa can be obtained with CBN, WC, and others. It can be noted here that the elastic modulus of steel is around 210â•›GPa, and many of the ceramics can have modulus two or more than two times higher. As is explained later, the modulus of materials plays an important role in determining the contact damage in elastic contacts when a ceramic is one of the mating materials. Also, high modulus of materials is beneficial to ensuring good wear resistance, and therefore, ceramics have an edge over metals in such a scenario.
3.3.5â•… Fracture Toughness It has been reported that the toughness of brittle materials is dependent on the test techniques,11–16 which are widely classified into long crack and short crack methods. Long crack methods include the single-edge notched beam (SENB) and single-edge V-notched beam (SEVNB) techniques. Short crack techniques involve measurement of the crack lengths (radial–median) around hardness indentations, from which the toughness data can be approximated using various reported models.11,13,16–20 In the context of tribological properties, the indentation technique can be used and an illustrative example is shown in Figure 3.9. This is particularly relevant as far as the tribological applications of ceramics, in which case it is the short crack fracture toughness which will have a larger influence on wear behavior. This is particularly true as microcracks and subsequent spalling are observed on worn surfaces of many ceramics. It is also recommended that the toughness of ceramics with new compositions be measured at different loads to check whether toughness increases with crack length, leading to “R-curve” behavior. It should be mentioned here that the indentation method is now routinely used for convenience to compute the fracture toughness of small and relatively brittle specimens, which are otherwise hard to machine into standard test samples (e.g., SENB, SEVNB). Small surface cracks of controlled size and sharpness can be induced in brittle materials such as ceramics by hardness indenters (Fig. 3.10). The median cracks (indicated by mc in Fig. 3.10), emanating from the corners of a Vickers indentation, are arrested when the residual stress driving force (Kres) at the crack tip is in equilibrium with the fracture toughness6:
K Ic = A( E/H )n × ( P/c3 / 2 ) = χ( P/c3 / 2 ) = K res, 1/2
(3.14)
where KIc is the indentation fracture toughness (Pa m ), E is the elastic modulus (GPa), H is the Vickers hardness (GPa), P is the indentation load (N), χ is the residual stress factor, and c is the crack length measured from the center of the indent impression (m). Furthermore, fracture toughness (KIc) values at different indent loads can be estimated from the measurement of radial cracks (indicated by rc in Fig. 3.10) around the Vickers indentations, according to the formula proposed by Anstis et al.17:
CHAPTER 3â•… Overview: Mechanical Properties of Ceramics
2c
32â•…
indent
l
2a cracks
99 µm
(a)
149
µm
50 µm (b)
Figure 3.9â•… (a) Schematic of indentation-induced cracking and (b) scanning electron microscopy (SEM) observation of indentation-induced cracking in a ceramic.
K Ic = 0.016( E/H )1/ 2 P/c 3 / 2.
(3.15)
The expression for toughness estimation as proposed by Anstis et al.17 was modified by Kaliszewski and coworkers13 to account for the effect of the compressive stresses due to the surrounding transformation zone:
K Ic = 0.019( E/H )1/ 2 P/c3 / 2.
(3.16)
3.4 Toughening Mechanisms
c
â•… 33
2a l
2a
rc
mc
lc
lc
(a)
(b)
Figure 3.10â•… Cross-sectional views of material surface around Vickers indentations showing (a) the formation of lateral cracks (lc) and median cracks (mc) in high-load regime and (b) the radial cracks (rc) and lateral cracks in low-load regime (b). (Reproduced with permission from Reference 19.)
Evans and Wilshaw18 reported that for a number of brittle materials, Palmqvist cracks were formed in the low-load regime. The dimensions of the Palmqvist and median cracks are related by the following expression: l/a = c/a −1.
(3.17)
For Palmqvist cracks (0.25╯<╯l/a╯<╯2.5), the fracture toughness (KIc) can be obtained using the expression of Niihara et al.19: ( K Ic ϕ/H a1/ 2 ) × ( H/Eϕ)2 / 5 = 0.035(l/a)−1 / 2 ,
(3.18)
where φ╯=╯3 and l╯=╯Palmqvist crack length for median cracks. For median cracks (c/aâ•›≥â•›2.5), the corresponding expression is ( K Ic ϕ/H a1/ 2 ) × ( H/Eϕ )2 / 5 = 0.129(l/a )−3 / 2.
(3.19)
20
Shetty et al. modified the equation of Niihara et al. and proposed the following expression for toughness estimation:
K Ic = 0.025( E/H )0.4 ( HW )1/ 2 ,
(3.20)
where W╯=╯P/4a (P is the indentation load, 2a is the Vickers diagonal).
3.4 TOUGHENING MECHANISMS Prior to the discussion on various possible toughening mechanisms in ceramics, it is important to mention that the brittleness of ceramics is due to a multitude of factors6: 1. In ceramics with predominantly ionic binding, the dislocations can only glide on specific slip planes, as schematically shown in Figure 3.11a.
34â•…
+ – + – + – + –
CHAPTER 3â•… Overview: Mechanical Properties of Ceramics
– + – + – + – +
+ – + – + – + –
– + – + – + – +
+ – + – + – + –
– + – + – + – +
+ – + – + – + –
– + – + – + – +
+ – + – + – + –
– + – + – + – +
(a)
(b)
Figure 3.11â•… Schematic illustration showing (a) the possibility of dislocation glide at specific planes in ionic ceramic and (b) the difficulty in dislocation glide due to rigid bond network in covalent ceramic.
2. For ceramics with predominantly covalent bonds, the dislocation movement is quite difficult due to directional properties and inherently rigid bond network, as such movement necessitates bonds to be broken and remade as well as bond angles to be distorted (see Fig. 3.11b). 3. Many ceramics do not have five independent active slip systems, and therefore, any homogeneous deformation without localized fracture is not feasible. 4. The dislocation core width of ceramics is shorter than that in metals, leading to the requirement of high Pierls–Nabarro stress for dislocation glide.2 Once cracking starts, it is therefore rather easy for cracks to grow in ceramics. In the case of metals, yielding occurs due to dislocation movement in the crack tip stress field. This absorbs a fraction of the available energy at the crack tip, thereby reducing the total driving force for further crack propagation. In view of the lack of
3.4 Toughening Mechanisms
Process zone
â•… 35
Bridging zone
Microcracking Fiber reinforcement
Phase transformation
Whisker reinforcement
Crack deflection
Ductile metal bridging
Figure 3.12â•… Summary of various toughening mechanisms in ceramic-based materials.6
dislocation glide in ceramics, such a phenomenon is ruled out; therefore, once a crack attains critical size, it can grow in an unstable manner, leading to fracture of the ceramic component. From fundamental aspects, if the interaction between a growing crack and the microstructure can absorb a fraction of the energy available at the crack tip stress field, then the driving force for crack propagation will be lowered. In other words, the crack opening displacement will be consequently reduced, and as a result, the crack tip will be blunted. Such mechanisms to improve crack growth resistance, that is, toughening mechanisms, can be broadly classified into two generic types (see Fig. 3.12): 1. Process zone mechanisms:╇ The enhanced crack growth resistance in transformation-toughened ceramics is realized due to phase-transformationinduced volume expansion or microcracking or crack deflection in the process zone around the crack tip. In case of transformation toughening, for example, transition from tetragonal to monoclinic zirconia in the crack tip stress field involves volume expansion. In a constrained microstructure, this results in compressive stress on the crack faces, leading to closure of the crack tip. Crack deflection is mostly realized in composites containing particulates, whiskers, or fibers as the second phase. Essentially, the crack path tortuosity is increased as the crack bypasses the hard or rigid second phase, and the model of Faber and Evans21,22 predicts a maximum increment in crack-deflection-induced toughening in composites with around 30╛vol╛% particulate reinforcement. The particle size and shape as well as the distribution of second phase particles also influence the achievable toughness increment. Later chapters of this book show how different toughening mechanisms influence the wear properties of ceramics and composites. 2. Bridging zone mechanisms:╇ This type of toughening is realized in fiberreinforced or whisker-reinforced composites. For example, both SiC matrix and SiC fibers are inherently brittle. However, the toughness of SiC/SiCf
36â•…
CHAPTER 3â•… Overview: Mechanical Properties of Ceramics
Ultimate tensile strength, su
Matrix microcracking stress, sm
Stress
Fiber pullout
Monolithic materials
Strain Figure 3.13â•… Schematic illustration of tensile stress–strain response of ceramic-fiber-reinforced ceramic composites vis-à-vis response of monolithic ceramic. (Adapted from Reference 23.)
composites can be very high and the underlying mechanisms involve crack bridging due to fiber–crack interaction. Such interactions in composites essentially involve matrix–fiber debonding or matrix microcracking in the crack wake of the matrix crack or mode I crack, crack deflection at the interface, fiber pullout, and, finally, crack bridging. The combination of these mechanisms essentially leads to nonlinear stress–strain tensile response of the composite, leading to damage-tolerant behavior (see Fig. 3.13).23 The extent of fiber pullout as well as properties of matrix–fiber interface strongly determines the toughness increment. The fiber toughening is also accompanied by modulus enhancement, and this is governed by the rule of mixtures. Fiber-reinforced composites, however, have their own limitations, which include an expensive production cycle and anisotropic properties with higher values along the fiber direction. Among various reinforcements, fibers have the largest aspect ratio, followed by whiskers, and then particulates. In whisker-reinforced composites with random whisker alignment, crack bridging also occurs but to a lower extent than fiber toughening. Crack bridging involves whisker pullout, and the toughness increment depends on the volume fraction of whiskers. A third type of toughening generally takes place in another important class of materials, known as cermets, which are characterized by dispersion of ceramic phase in a metallic matrix, for example, WC–Co. In the crack tip stress field, the ductile flow of metallic particles leading to bridging of crack faces reduces the available driving force for crack propagation. The toughness
REFERENCES
Material Composition - non-oxide ceramic (tribochemical wear) - oxide ceramic (brittle fracture) - glass (subcritical crack growth)
â•… 37
Elastic Modulus - contact damage resistance
Wear Resistance
Hardness - abrasive/adhesive
Fracture Toughness - short crack toughness - lateral cracking
Figure 3.14â•… Schematic illustrating the influence of various mechanical properties in the context of the tribological applications of ceramics and composites.
of WC–Co cermet therefore decreases with lowering of Co content. From the material development point of view, a critical amount of Co is required to facilitate liquid-phase sintering, and the balance of hardness and toughness demands the tailoring of Co content.
3.5 CLOSING REMARKS As a concluding note, it is important to mention that a combination of mechanical properties as well as material composition, in terms of type of matrix and reinforcement, influences the wear resistance properties of monolithic ceramics and ceramic composites, as shown in Figure 3.14. The development of new ceramics should ideally aim to obtain the best combination of higher E-modulus, hardness, and fracture toughness, so that such materials would have a better resistance against tribomechanical wear, as is discussed in various chapters of this book.
REFERENCES ╇ 1â•… C. E. Inglis. Stresses in a plate due to the presence of cracks and sharp corners. Trans. Inst. Nav. Arch. 55 (1913), 219–210. ╇ 2â•… G. E. Dieter and D. Bacon. Mechanical Metallurgy. McGraw Hill, London, 1988. ╇ 3â•… E. Orowan. Fatigue and Fracture of Metals, Symposium at Massachusetts Institute of Technology, USA. John Wiley & Sons, New York, 1952. ╇ 4â•… A. A. Griffith. Philos, Vol. 221A. Transactions of Royal Society, London, 1920, 163–198. ╇ 5â•… S. M. Wiederhorn. Moisture assisted crack growth in ceramics. Int. J. Fract. Mech. 4(2) (1968), 171–177. ╇ 6â•… B. R. Lawn. Fracture of Brittle Solids. Cambridge University Press, Cambridge, 1993. ╇ 7â•… A. Tiwari, B. Basu, and R. Bordia. Model for fretting wear of brittle ceramic. Acta Mater. 57 (2009), 2080–2087. ╇ 8â•… Standard Test Method for Dynamic Young’s Modulus, Shear Modulus and Poisson’s Ratio for Advanced Ceramics by Impulse Excitation of Vibration ASTM Designation, June, 1995, 375. C1259-95.
38â•…
CHAPTER 3â•… Overview: Mechanical Properties of Ceramics
╇ 9â•… G. Roebben, B. Basu, J. Vleugels, J. Van Humbeeck, and O. Van Der Biest. The innovative impulse excitation technique for high temperature mechanical spectroscopy. J. Alloys Comp. 310(1–2) (2000), 284–287. 10â•… B. Basu, L. Donzel, J. Van Humbeeck, J. Vleugels, R. Schaller, and O. Van Der Biest. Thermal expansion and damping characteristics of Y-TZP. Scr. Mater. 40(7) (1999), 759–765. 11â•… G. D. Quinn and R. C. Bradt. On the Vickers Indentation Fracture Toughness Test. J. Am. Ceram. Soc. 90(3) (2007), 673–680. 12â•… B. R. Lawn. Indentation of ceramics with spheres: A century after Hertz. J. Am. Ceram. Soc. 81(8) (1998), 1977–1994. 13â•… M. S. Kaliszewski, G. Behrens, A. H. Heuer, M. C. Shaw, D. B. Marshall, G. W. Dransmann, and R. W. Steinbrech. Indentation studies on Y2O3-stabilized ZrO2: I. Development of indentationinduced cracks. J. Am. Ceram. Soc. 77(5) (1994), 1185–1193. 14â•… B. R. Lawn, N. P. Padture, H. Cait, and F. Guiberteau. Making ceramics “ductile”. Science 263 (1994), 1114–1116. 15â•… J. B. Quinn and G. D. Quinn. Indentation brittleness of ceramics: A fresh approach. J. Mater. Sci. 32 (1997), 4331–4346. 16â•… P. Chantikul, G. R. Anstis, B. R. Lawn, and D. B. Marshall. A critical evaluation of indentation techniques for measuring fracture toughness. J. Am. Ceram. Soc. 64(9) (1981), 539–543. 17â•… G. R. Anstis, P. Chantikul, B. R. Lawn, and D. B. Marshall. A critical evaluation of indentation techniques for measuring fracture toughness: I, direct crack measurements. J. Am. Cer. Soc. 64 (1981), 533–538. 18â•… A. G. Evans and T. R. Wilshaw. Quasi-static solid particle damage in brittle solids—I. Observations, analysis and implications. Acta Metallurgica 24(10) (1976), 939–956. 19â•… K. Niihara, R. Morena, and D. P. H. Hasselman. Evaluation of KIc of brittle solids by the indentation method with low crack-to-indent ratios. J. Mater. Sci. Lett. 1 (1982) 13–16. 20â•… D. K. Shetty, I. G. Wright, P. N. Mincer, and A. H. Heuer. Indentation fracture toughness of WC–Co ceramics. J. Mater. Sci. 20 (1985), 1873–1882. 21â•… K. T. Faber and A. G. Evans. Crack deflection processes—I. Theory. Acta Metallurgica 31(4) (1983), 565–576. 22â•… K. T. Faber and A. G. Evans. Crack deflection processes—II. Experiment. Acta Metallurgica 31(4) (1983), 577–584. 23â•… (a) K. K. Chawla. Ceramic Matrix Composites. Chapman and Hall, New York, 2003. (b) K. K. Chawla. Composite Materials: Science and Engineering. Springer, New York, 1998.
CHAPTER
4
SURFACES AND CONTACTS This chapter discusses the characteristics of engineering surfaces in the context of tribological applications. Quantification in terms of surface roughness parameters as well as the implication of such parameters in terms of distinguishing various surfaces is discussed. Finally, a description of surface and subsurface stress fields for ballon-flat tribocontact is provided, based on Hertzian contact mechanics.
4.1 SURFACE ROUGHNESS Tribological surfaces are never as ideal as we can draw them on a piece of paper, but are made using real machines, from real materials, measured with devices having finite accuracy, and so on.1 Therefore, they always contain irregularities, defects, and discrepancies from what is considered to be how they should look (see Fig. 4.1). Basically, we can divide tribological surfaces into a few general types: smooth and even, smooth and wavy, rough and even, and rough and wavy. These “errors” can be described as microgeometrical or macrogeometrical. We can consider microgeometrical errors as roughness, which is mainly caused by the direct contact of the tool at the surface. On the other hand, macrogeometrical errors are of larger degree and are typically caused by the dynamic and/or kinematic characteristics of the tools or machines that manufacture or finish engineering surfaces. For example, while using a grinding wheel, the eccentricity will cause waves on the finished surface (macrogeometrical error), while defects on the grinding wheel surfaces (e.g., rounded or missing hard particles) will affect the roughness of the surface (microgeometrical error).2 Of course, it is questionable what the absolute values are for the aforementioned definitions—what is wavy and what is rough (at least on an atomic level, the surfaces are uneven and not smooth). First, there must be ways to measure and evaluate them, but also the same values will not have the same meaning for every application. For example, asperities of a micrometer or a few hundred nanometers at the gear surface are acceptable, while these are completely intolerable for a hard disk drive surface, where a gap between the disk and disk head can be as low as a few nanometers. Therefore, standard roughness parameters should be used to define the values of roughness and also to eliminate effects of measuring waviness, which
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
39
40â•…
CHAPTER 4â•… Surfaces and Contacts
Figure 4.1â•… Schematic illustration of topography of nominal engineering surface (bold line) and real surface to show the microscale roughness and waviness on the surface (dashed line).
Rq = 0.58a Ra = 0.25a
Rq = 0.37a Ra = 0.25a
+a
–a
good
bad Y
Rq
Ra
X mean line
Figure 4.2â•… Schematic illustration showing surface roughness parameters (bottom) as well as roughness profiles of typical bad or rough surfaces in the context of tribological applications (top).
is usually a subject of design and manufacturing rather than of tribological optimization of the surfaces. In this way, two-dimensional (2D) parameters such as Ra and Rq are used, and lately also three-dimensional (3D) parameters (S)—similarly defined but considering several surface profiles (i.e., topographic characteristics). With reference to a typical 2D surface profile (Fig. 4.2), the arithmetical mean deviation or average surface roughness (Ra) for a given sample length can be defined as the mean height of the surface profile (peaks and inverted valleys). Therefore, Ra is the arithmetical mean of all profile values of the roughness profile, that is, average value of the departure of the profile from the mean line:
Ra =
1 l
∫
1
0
y( x ) dx,
(4.1)
where y(x) is the profile height from the mean line at a given distance x on the profile length, and l is the total profile length. The parameter Rq is called the root mean square value of the roughness profile, which is mathematically defined as follows:
4.2 Surface Topography and Asperities
Rq =
1 l
∫
1
0
y 2 ( x )dx .
â•… 41 (4.2)
Everyone who uses these parameters should also be aware of their exact meaning, because they are mathematical and/or statistical functions of certain measuring device readings. Therefore, errors are caused by measuring as well as by the calculation itself. Some parameters are simple average values of all asperities, which cannot distinguish peaks and valleys, but some are more sophisticated and consider these qualitative differences. Sometimes these are irrelevant, for example, in standard well-known and controlled surface preparations, and used just to control the quality of the process. On the contrary, with bearing surfaces that are usually lubricated (seals, bearings, piston rings, etc.), some peaks or valleys can cause various unexpected and uncontrolled negative (leakage) or positive (lift-up) effects. So, it must be decided which surface roughness parameters should be used and measured to represent a surface by the most appropriate parameter(s), as this will affect the cost and complexity of measurement. Finally, while there are well-defined standard roughness parameters, there exist several methods that can be used for measuring the surfaces. They can be contact types (stylus-tip, atomic force microscopy [AFM]) or noncontact types (optical interferometry, electronic acquisition in scanning electron microscopy [SEM]), each having positive and negative aspects with respect to simplicity, cost, dimensional limitations, optical reflectivity, modification or deformation of measured matter, and so on. However, it is also important to critically judge the measurements on different scales, which is often the case when comparing data of AFM and stylus-tip instruments. Namely, as mentioned previously, any data are a consequence of measurement and mathematical operation, and sometimes these can be misleading when not properly understood and evaluated.
4.2 SURFACE TOPOGRAPHY AND ASPERITIES Surface topography describes 3D characteristics of the surface. Although surface roughness is measured using 2D surface profiles, topographic information consists of several successive profiles. The distance between these profiles should be as small as possible to represent the surface as closely as possible to its actual form. However, there is always some lateral resolution that limits the measuring device and analytical models used to represent the actual surface realistically. Nevertheless, the topographical visualization of the surface can provide a significant improvement over 2D profiles, because it shows much better the statistical nature of the surface, orientation of scratches, real size of holes or asperities, and so on. Two almost equal 2D surface profiles (identical Ra value) can belong to two completely different types of surfaces, and therefore, additional roughness parameters (Rq) can be used to distinguish such surfaces, as shown in Figure 4.3. Three-dimensional laser profilers have become handy and useful tools for excellent visualization of surfaces and are used to measure all the 2D and sometimes 3D surface characteristics according to standard roughness evaluations.
42â•…
CHAPTER 4â•… Surfaces and Contacts
Ra = 2.4 µm
Ra = 2.5 µm
Ra = 2.4 µm
Figure 4.3â•… Schematic illustration of three engineering surfaces with nearly identical Ra values but with surface asperities of different heights and slopes.
One feature that becomes more obvious in 3D visualization of a surface is the slope of the surface asperities. Namely, when measuring roughness parameters, the asperities always look very sharp and their slopes are steep (see Fig. 4.3). Typical engineering surfaces are characterized by asperities of slope of less than 20°, typically less than 5–10°. It is often neglected or forgotten that the asperity heights of engineering surfaces are on the order of micrometers, and more typically nanometers. On the other hand, the measuring lengths, which are standardized, are in millimeters or at a minimum hundreds of micrometers. Accordingly, the two measurements cannot be presented in the same scale, because they are different orders of magnitude (see, e.g., Fig. 4.4). In light of such reasons, heights are usually presented in a scale that is enlarged 20 or 25 times compared with profile length. Because of this, the heights appear larger than they are and slopes steeper. Also, as the surfaces are smoother, the asperities become flatter. For example, various locations of an apparently single asperity, as marked by ‘A,’ ‘B,’ ‘C,’ and ‘D’ on the bottom part of Figure 4.4, are clearly discernible in the top part of Figure 4.4 due to change in x-y scale.
4.3 REAL CONTACT AREA However, it is certainly important to summarize from the preceding two sections that the surfaces are not “flat”! They always have some heights above the surface
â•… 43
4.3 Real Contact Area
D
C
B
A
500x 5000x B A
C
D
5000x 100x
Figure 4.4â•… Schematic illustration of the scale of difference between asperity heights along the length of a surface profile.
V
2a
A
V
V Ai
Figure 4.5â•… Schematic illustration of the contact of a single asperity with a nominally flat surface.
mean line, and there are some valleys below it. Accordingly, when two surfaces approach each other and eventually come in contact, this will never occur on the whole surface, but only on the asperities that touch first. The area of these junctions is called the real contact area. In reality, it depends on the roughness, material mechanical properties (hardness, elastic modulus, etc.), load, type of material deformation (brittle, plastic), and so on. A schematic illustration of single-asperity contact against a nominally smooth flat surface is shown in Figure 4.5. As real engineering surfaces contain n such asperities, the real contact area (Ar) can be expressed as the sum of individual contact areas (Ai) of such asperities:
44â•…
CHAPTER 4â•… Surfaces and Contacts n
Ar =
∑ A.
(4.3)
i
i =1
There is an everlasting question—what is the size of the real contact area in certain contacts and applications? Although there have been many attempts in the past to model or measure it, there would be almost as many answers to this question as there would be people responding. Accordingly, it is not our intention to go deeper into discussion of this area. However, it is believed that values below 20% or even 5% or less of the nominal contact area are common. Finally, there is an important consequence that should be noted in this discussion. That is to say, that the contact pressures in tribological contacts are not the nominal ones calculated based on macrogeometrical data, but the pressures build up on the real contact area. Moreover, because of this, the materials are exposed to deformation and/or fracture. Furthermore, the heat that is generated through frictional losses is distributed into materials mainly via conduction, rather than convection or radiation, so the surface temperatures can also reach much higher values than normally expected.
4.4 CONTACT LOAD DISTRIBUTION AND HERTZIAN STRESSES Tribological contacts are usually small, especially when dealing with concentrated nonconformal contacts, such as those in bearings, gears, and various tools. A common contact-mechanics treatment is to use well-established Hertzian theory, which can be employed for a variety of contact geometries, if elastic contact conditions can be assumed.3 In particular, for tribology or materials research, it is very convenient to use test samples with well-defined geometries, materials, and loads; in these cases, Hertzian theory works very well, and a detailed description is available in the book by Johnson.1 Most often, a sphere-on-flat contact geometry is used (Fig. 4.6), and
Stationary (rigid body) p0 q0
U
X
2a z
Elastic (perfectly elastic)
Figure 4.6â•… Schematic illustration of Hertzian contact geometry and contact pressure distribution (p0) over contact diameter (2a).
â•… 45
4.4 Contact Load Distribution and Hertzian Stresses
in such a contact design, the contact parameters can be calculated as presented next as an example. Hertzian theory with its many adaptations for specific cases, conditions, and parameters is published in a variety of literature, so it is not presented in detail here, apart from the simple ball-on-flat case. As shown in Figure 4.6, the contact pressure distribution, p(r), has a parabolic shape and its distribution can be calculated as
p(r ) = p0 1 −
r2 , a2
(4.4)
where r2╯=╯x2╯+╯y2. Contact pressure increases from its zero value at the contact edges to a maximum p0 at the center of the contact. Based on this distribution a generally valid solution for the contact area and maximum contact pressure can be derived as
p0 =
3W , 2πa 2
(4.5)
where the radius of the contact area is a=
3
3WR* . 4E*
In these expressions, W is normal contact load, and E* and R* are equivalent elastic modulus and radius of contact bodies, respectively, which can be calculated using the following expressions:
1 (1 − ν12 ) (1 − ν22 ) = + , E* E1 E2 1 1 1 = + . R* R1 R2
(4.6) (4.7)
In Equations 4.6 and 4.7, suffixes 1 and 2 refer to two mating solids and ν is Poisson’s ratio. For the resistance of a given material against motion at a tribological contact under load, it is important to calculate the stresses at the surface and subsurface. Two very important cases for static Hertzian contacts should be noted (see Figs. 4.7 and 4.8). One important fact that comes from this calculation and affects the contact limiting loads significantly is that the radial normal stress changes from compressive to tensile at the edge of the contact, which presents a danger of cracking (Fig. 4.7). Another important conclusion, relevant primarily for long-term loading and described in detail in the delamination theory of wear, relates to subsurface shear stresses (see Fig. 4.8). Namely, the shear stress has a maximum value somewhere below the surface—in this particular case of point static contact for steel (i.e., ν╯=╯0.3), at a depth of 0.48a. This implies that cracks, which initiate in ductile materials at a value only half the tensile strength, will grow under the surface at relatively low loads and
46â•…
CHAPTER 4â•… Surfaces and Contacts
1
0.5 σθ p0
compressive 0
tensile 0
0.5
σr p0
1
1.5
r a
Figure 4.7â•… Variation of radial normal stress (surface) around a static Hertzian contact.1
0
0.31 τ1 p0
0.5
1 σr p0
0.48
σz p0 1
1.5
2 z a
Figure 4.8â•… Variation of subsurface normal and shear stress field of a static Hertzian contact.1
4.5 Closing Remarks
â•… 47
the surface may not experience any changes or “warning” until catastrophic failure, in the form of delaminated debris causing pitting, appears. Of course, by adding a tangential component to the load, that is, when considering friction force, the loading situation and stress distribution will change. Under static loading, stresses and deformations tend to be symmetric; however, during sliding, the stresses depend on the direction of sliding (tangential friction force) and become asymmetric.3 An obvious and frequently relevant case is that the normal stress in front of the contact will only be compressive, whereas tensile stresses will appear in the tail of the contact. Because of this, crack opening and/or surface spallation will always occur after the slider has passed a certain point. Moreover, tangential load also brings the shear stresses closer to the surface. The position of maximum shear stress is important and should be calculated in the case of coated or graded materials, where any critical position (borders of different phases) should avoid stress concentrations. Accordingly, it is advisable to calculate the stresses for any application or experimental design to avoid maximum values at critical positions, as well as to limit the loads for a long-term suitable contact performance.
4.5 CLOSING REMARKS As a concluding note, the influence of surface roughness parameters on various aspects of asperity–asperity contact behavior as well as friction is illustrated in Figure 4.9. It should be clear from the discussion in later chapters that the surface roughness parameters influence the coefficient of friction (COF) during the runningin period; however, the steady-state COF will depend on dynamic asperity–asperity interactions during the post-running-in period, and the influence of initial roughness would be minimal in the steady-state period. The adhesion or deformation of the asperity contact, as determined by the roughness parameters, would also influence the COF during the running-in period. Also, the initial roughness parameters, in combination with the physical properties of the mating materials, determine whether the contact behaves more like an elastic or a plastic contact, as described elsewhere.2
Surface roughness of mating materials (R a, R q, etc.)
Elastic-plastic contact
Asperity-asperity contact adhesion or deformation
COF during running-in period
Figure 4.9â•… Schematic illustration showing how the surface roughness parameters can influence friction and properties and behavior of asperity contacts. COF, coefficient of friction.
48â•…
CHAPTER 4â•… Surfaces and Contacts
REFERENCES 1â•… K. L. Johnson. Contact Mechanics. Cambridge University, Cambridge, UK, 1985. 2â•… B. Bhushan. Principles and Applications of Tribology. John Wiley & Sons, New York, 1999. 3â•… K. L. Johnson. Surface interaction between elastically loaded bodies under tangential forces. Proc. R. Soc. Lond. A230 (1955), 531.
CHAPTER
5
FRICTION This chapter briefly discusses the laws of friction and factors contributing to friction at tribological interfaces experiencing sliding motion. A summary of the friction values of some common engineering metals, ceramics, and polymers is provided, and the difference in frictional response is discussed in light of the differences in structure and bonding characteristics.
5.1 INTRODUCTION Friction is the resistance to motion during sliding or rolling that is experienced when one solid body moves tangentially over another contacting surface1 (see Fig. 5.1). The resistive tangential force, which acts in a direction directly opposite to the direction of sliding motion, is called the friction force. To overcome this force, a certain amount of work is required, which is in fact an energy loss. This can be observed as frictional heat generation, noise, vibration, or the energy required to deform materials in contact, as well as causing tribochemical reactions. The next chapter discusses the aspect of frictional heat generation and contact temperature. In general, friction force is usually considered as an undesired, or “bad,” phenomenon; however, in certain applications, it can also be a “good” one. For example, without friction it would be impossible to walk, use automobile tires on a roadway, or pick up objects. Even in some machine applications, such as vehicle brakes, clutches, and frictional transmission of power (such as belt drives), friction is maximized. However, in most other sliding and rotating components, such as bearings, seals, and piston rings, friction is undesirable because it is a source of relevant energy losses and wear of moving surfaces in contact. In these cases, friction needs to be minimized.
5.2 LAWS OF FRICTION Three basic laws of intrinsic (or conventional) friction are generally obeyed as a first approximation over a wide range of applications. Two of these laws are often referred as Amontons’ laws,2 after the first person to present them in a clear and
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
49
50â•…
CHAPTER 5â•… Friction
W F
Direction of motion
W
W F F
F W W
F
(a)
(b)
Figure 5.1â•… Schematic illustrations of (a) a body sliding on a surface (free-body diagram), and (b) a body rolling on a horizontal surface. W is the normal force and F is the friction force.1
straightforward definition, although it is known today that the same phenomena were recognized and studied earlier by DaVinci, whose notes were, however, discovered later. The first Amontons’ law states that the frictional force is directly proportional to the normal load, which is mathematically described as
F = µW ,
(5.1)
where F is the friction force, W is the normal force, and μ is a constant defining this linear proportion, known as the coefficient of friction (COF). This actually means that by multiplying the normal load by a certain factor, the frictional force will change by the same multiplication factor. This is also the most simple and commonly used definition for friction. Typically, the rolling COF, as experienced by ball bearings against a raceway is much lower than the COF under sliding motion. This chapter is primarily focused on sliding friction. The second Amontons’ law states that the friction force is independent of the apparent area of contact between the contacting bodies. Thus, the same body, regardless of the physical size or shape of the contacting surface, will result in the same frictional force when moving over a certain counterbody. It is easy to interpret this law when we imagine moving a rectangular box over a flat on sides with different sizes: the required force will always be the same. To these two laws, a third law is often added, known as Coulomb’s3 law. It states that the friction force (or COF) is independent of sliding velocity once motion has started. With this, Coulomb made a clear distinction between static friction and kinetic friction. Namely, it is well known that the frictional force required to start the motion of a body is larger than that required to maintain it. The coefficient of static friction may be as much as 20–30% higher than the coefficient of kinetic fric-
â•… 51
5.3 Friction Mechanisms
Tang gential fforce
Typically 10–100 ms F static
Relative motion initiated
F kinetic
Time
Figure 5.2â•… Tangential force as a function of time or displacement; Fstatic is the static friction force required to initiate motion and Fkinetic is the kinetic friction force required to sustain motion.1
tion in some cases. A schematic illustration showing the frictional response revealing the distinction between static and kinetic friction is shown in Figure 5.2. However, anybody who has ever measured friction would know that the COF will not remain (exactly) constant with change in velocity. The factors affecting this change include heating and consequently materials modifications (softening), formation of boundary layers, running-in, and different asperity impacts. Moreover, the friction force does not remain constant as a function of distance or time and sometimes even produces a form of oscillation, based on a so-called stick–slip phenomenon. During the stick phase, the friction force builds up to a certain value and a large enough force has to be applied to overcome the static friction force, and then slip occurs at the interface. These three laws are entirely empirical; situations in which these laws are not followed do not imply violation of more fundamental laws of nature. Actually, friction seldom exactly follows these laws, but typically depends on many other parameters: surface morphology, material properties, and environmental conditions having the most influence. To some extent, it is surprising how difficult it is to predict the friction by knowing only the materials and basic operating conditions. However, we must consider that friction is not a material property, but rather a system property that depends on both materials in contact, as well as on the operating conditions and environment. With this in mind, it is clear that these parameters and their interrelations change all the time, and so does the friction.
5.3 FRICTION MECHANISMS Generally, friction involves energy dissipation mechanisms during relative motion of two bodies. As two engineering surfaces are brought into contact, contact occurs at the tips of asperities and the load is supported by the discrete contact spots that
52â•…
CHAPTER 5â•… Friction
are formed. Therefore, under nonlubricated conditions, the frictional force will be the total sum of the tangential frictional forces generated at all these spots. The conditions at the asperities can vary, and this affects the mechanisms of friction (and wear) at these junctions. Typically, several mechanisms can act simultaneously, but it is likely that one mechanism will be predominantly responsible for the overall friction behavior, contributing the larger part to the total friction. It is thus the goal to recognize the most relevant friction mechanism(s) that need to be controlled under specific contact conditions. Friction under nonlubricated conditions arises due to several possible mechanisms: adhesion, deformation, and plowing/cutting being the most relevant. The adhesion term constitutes the force required to shear the adhesive bonds formed at the interface in the regions of real area of contact. In the case of metals, ceramics, and other hard materials, the deformation term constitutes the force needed for deformation of asperities (microscale), whereas plowing and/or cutting is the contributing term when asperities of one material, typically harder, groove or cut the surface of the mating material (macroscale). Plowing of one or both surfaces can also occur by wear particles trapped between them. In the case of viscoelastic (rubbery) materials, the contribution to friction is largely attributed to the deformation term due to elastic deformation hysteresis. In case of friction behavior, the effect of adhesion, deformation of asperities, and plowing can be discussed as follows. During sliding, changes in the conditions of mating surfaces occur, which affect friction and wear properties. Initially, the friction force is typically very high, which coincides with the static friction before the motion of the bodies and with the interlocking of asperities in the very early period of sliding. However, the length of this period can vary and can be sometimes very short, only a few passes, and thus, in terms of the component lifetime, negligible. On the other hand, sometimes it can last much longer and even never stop, if the mating combination is inappropriate for the so-called running-in. During the running-in period, high asperities may be deformed or fractured, wear occurs, and surfaces may mate better; initial surface films (oxides, contaminants) may be worn, new steady films may be formed, or structural changes may occur. These changes result in friction either going up or coming down from the initial value. The run-in period is critical for long interface life as incorrect run-in can result in serious damage and early failure. Accordingly, proper run-in can be extremely useful for “smooth” compliance of the engineering surfaces without any subsurface damage, which can lead to long-lasting beneficial surface conditions. In many cases, special surface treatments can be used to facilitate the beneficial running-in, such as phosphating and functionally graded multilayer surface treatments. After some period of contact under operation (i.e., after run-in), the friction force generally stabilizes into what is known as steady-state sliding. Depending on many parameters, including run-in, several scenarios are possible for further development of the COF, leading to low or high friction. Often, after sliding for a period of time in steady state, friction increases again and reaches another plateau, as shown by the S-shaped curve in Figure 5.3a. Namely, during or after the steady-state period, changes in the interface may further occur, such as roughening and trapped particles, which lead to an increase in friction. This process can repeat, and friction can reach
â•… 53
5.3 Friction Mechanisms
Coefficient of friction
Transition period Steady-state sliding S1 Run-in period
Steady-state sliding S2 Average Range
II
I
Distance (time) (a)
IV
Coefficient of friction
III
II I
Distance (time) (b)
Figure 5.3â•… Coefficient of friction as a function of sliding distance with (a) a typical S-shaped curve showing run-in period and (b) four hypothetical cases.1
several plateaus. After a useful interface lifetime, the interface fails and friction may become very high. The shape of friction curves can be affected by the interface materials as well as by operating conditions. Friction may increase in different patterns, such as the following: (I) the friction may remain at its initial value for some time and slowly increase to another steady-state value; (II) after being at the initial value for some time, it may first increase to a high value then level off at a lower value (but higher than the initial value); (III) it may increase to a high value, level off to this value, drop to a lower value, and increase again to a high value; (IV) it may change in a nonrepeatable manner, as shown in Figure 5.3b. In almost all cases, the COF would reach a high
54â•…
CHAPTER 5â•… Friction
value after some period of sliding. Identical metals in contact exhibit the behavior shown in case (I); the increase is associated with plowing because of roughening and trapped wear particles. In smooth surfaces involving elastic deformation with adhesive friction, the dominant component, the increase in COF is associated with smoothing of the surfaces, leading to a larger component of adhesive friction.4 In Figure 5.3b, the drop in the COF in case (II) is associated with smoothening of the two hard surfaces experiencing plastic deformation, which results in a drop in the plowing component of the friction. For elastic contacts where the adhesive component is dominant, roughening and/or trapped wear particles reduce the real area of contact, which in turn reduces the adhesive component of the friction. The drop in the friction in case (III) in plastic contacts can be associated with ejection of wear particles, and a subsequent increase is associated with generation and entrapment of wear particles. A significant increase in friction to an unacceptably high value in a short period in case (IV) is associated with a poorly selected material pair in which several sources contribute to friction.
5.4 FRICTION OF COMMON ENGINEERING MATERIALS In this section, it needs to be clearly pointed out that the friction is not a material property, but a tribological system property. This means that friction cannot be predicted based only on known contacting material(s), but it will depend on the mating material, contact (component) operating conditions, stiffness of the system and contacts, and surface conditions (roughness, topography, etc.) as well as on the environment and any tribochemical reaction that forms in situ, often only monolayer thin, boundary films. Nevertheless, some material combinations develop typical forms of boundary films and exhibit typical behavior when in contact, because of their predominant properties that affect friction behavior. For example, some materials are always hard; some are brittle and remain rough due to brittle fracture; some smear and deform, and tend to become smooth causing high real contact area; some are chemically unstable; and some oxidize, some do not. Based on these material characteristics, and in the absence of any very specific conditions, such as extreme operating conditions, (particular) lubrication, and extreme environments, we may find a relatively good approximation of the values of expected COF. These values are sometimes found in the literature as friction between certain materials. Although not fully correct, they are valuable in terms of understanding the basic principles of mating these materials and can be used for comparison between different combinations. They also give us rough estimation of friction losses. Some of these typical friction values of metals, alloys, ceramics, polymers, and solid lubricants are presented in Tables 5.1–5.3. In Table 5.1, the unlubricated COF values of metals and alloys as well as various metallic materials both self-mated and against mild steel, under ambient conditions, are provided. The variation of COF for a given friction couple depends on variation in normal load, sliding velocity, and so on. As can be noted, COFs of lower than 0.3 are not achievable with known
TABLE 5.1â•… Typical Values of the Coefficient of Kinetic Friction of an Unlubricated Metal or Alloy Sliding on Itself or on Mild Steel at Room Temperature in Air1
Material
Coefficient of friction Self-mated
On mild steel
Pure metals Precious Metals Au, Pt, Ag
1–1.5
0.3–0.5
Soft Metals In, Pb, Sn
0.8–2
0.5–0.8
0.8–1.2 0.5–0.6 0.8–1.2 0.8–1.5 0.5–0.6 0.7–0.9 0.5–0.6 0.7–0.9
0.5–0.6 0.4–0.5 0.6–0.7 0.8–1.5 0.4–0.6 0.6–0.9 0.4–0.6 —
— 0.8–1 0.7–0.9
0.2–0.4 0.3–0.5 —
Metals Al Co Cu Fe Mo Ni Ti W Alloys Leaded brass (Cu, Zn, Pb) Gray cast iron Mild steel
Intermetallic alloys Co-based alloy Ni-based alloys
0.3–0.5 0.6–0.9
TABLE 5.2â•… Typical Values of the Coefficient of Kinetic Friction of an Unlubricated Ceramic on Itself at Room Temperature in Air1
Material
Coefficient of friction
Al2O3 BN Cr2O3 SiC Si3N4 TiC WC TiN Diamond
0.3–0.6 0.25 0.25 0.3 0.25 0.3 0.3 0.25 0.1
— —
56â•…
CHAPTER 5â•… Friction
TABLE 5.3â•… Typical Values of the Coefficient of Kinetic Friction of Polymers and Solid Lubricants on a Hard Surface at Room Temperature in Air1
Material
Coefficient of friction Polymer–plastics
Acetal Polyamide (nylon) High-density polyethylene (HDPE) Polytetrafluoroethylene (PTFE) (Teflon)
0.2–0.3 0.15–0.3 0.15–0.3 0.05–0.10
Polymer–elastomers Natural and synthetic rubber Silicone rubber
0.3–0.6 0.2–0.6 Solid lubricants
Layer-lattice solids MoS2 Graphite
0.05–0.10 0.05–0.15 Nonlayer-lattice solids
Fullerenes (C60)
0.05–0.10
engineering metallic materials. A comparison of Table 5.2 with Table 5.1 reveals that a greater number of the self-mated ceramics can exhibit low COF (0.3 or lower). Again, COFs of ceramics vary over a window of 0.3–0.7; interestingly, low COFs of 0.1–0.2 can be achieved with diamond-based materials and such low friction cannot be achieved with known engineering metals or alloys. Usually, clean metal and alloy surfaces exhibit high adhesion and, consequently, high friction and wear under self-mating conditions. The steady-state COF of contacting metallic surfaces in the absence of contaminant layers such as oxides or lubricants (e.g., in extreme high vacuum), can be very high, that is, much higher than values reported in Table 5.1. Strong metallic bonds are formed across the interface, and significant transfer films of metal from one body to another, or as loose wear debris, occur during sliding. Materials that can plastically deform at a frictional contact further increase adhesion due to junction growth and increase in the contact area, which needs to be sheared. However, the slightest contamination mitigates contact or forms chemical films, which reduce adhesion, resulting in reduced friction.5 The mechanical behavior of ceramics differs from that of metals and alloys because of the different nature of the interatomic forces of the covalent or ionic bonding in ceramics compared with that of the metallic bonding in metals and alloys. Ceramic materials of either bond type show only very limited plastic flow at room
â•… 57
5.4 Friction of Common Engineering Materials
temperature and much less ductility than metals. Although adhesive forces, of covalent, ionic, or van der Waals origin, are present between ceramic materials in contact, the low real area of contact results in relatively low values of the adhesive component of friction compared with metallic couples. Moreover, under clean environments, COFs of ceramic pairs do not reach the very high values observed in the case of clean metals, especially in ultrahigh vacuum or in the absence of oxygen.6 Nevertheless, ceramic pairs could result in high friction due to “mechanical” components of friction, such as asperity fracture and asperity interlocking. Surface preparation and finishing is essential to produce appropriate surface conditions for low friction. Moreover, it is known that environment (humidity) is critically important for ceramic friction (and wear). In certain cases, it could result in very smooth and low-friction tribochemical layers, while in other cases, it may cause hydrothermal transformation and surface fracture with high friction. An example of such behavior, as observed in the case of yttria-stabilized tetragonal ZrO2 ceramics is discussed in Chapter 10 of this book. Compared with metals and some ceramics, most polymers can exhibit much better frictional response in terms of lower steady-state COF. As shown in Table 5.3, many polymers can exhibit COF of 0.2 or lower under unlubricated conditions. From a structure–bonding point of view, polymers are characterized by a macromolecular chainlike structure. Each chain is made up of C, H, O, N elements, with covalent bonding along the chain and van der Waals bonding across the chain. Under stress, rearrangement or sliding of neighboring chains over each other is possible. Such a phenomenon contributes to time-dependent viscoelastic deformation behavior of polymers. Polymers can be broadly classified into three major groups: (1) thermoplastics (high-density polyethylene [HDPE], polytetrafluoroethylene [PTFE], etc.), which can be deformed to various shapes when heated; (2) thermosets (polyoxymethylene [acetal], polyamide [nylon], etc.), which cannot be deformed once shaped at high temperature; and (3) elastomers (rubber), which can exhibit large amounts of deformation (up to 900%) without fracture. Many polymers exhibit low friction and are thus often used as “solid” lubricants in engineering applications, sometimes also as “self-lubricating” materials, when addition of conventional lubricants is used in the polymer matrix. Polymers, in particular PTFE, polyamide, and HDPE are among the most commonly used materials due to their low friction and reasonably high load-bearing capacity. Among polymers, PTFE exhibits the lowest friction, which can be as low as 0.05. PTFE also performs well in different environments. The plastic flow of polymers at already low loads and temperatures certainly influences the friction behavior. Another important factor that influences the friction of polymers is their strong time dependency, implying that their mechanical properties change with time, in particular when associated with increased temperatures and deformation. Accordingly, the laws of friction (Amontons’, Coulomb’s) are not very useful for polymers due to their dependency on loads, speed, temperature, and change of mechanical properties with time. Therefore, all the data listed in Table 5.3 for various friction values should be evaluated in this respect. Although very low COF is achievable with the use of polymers, mainly the high wear rate and low melting or glass transition temperature (∼300–400°C) limits their wider application. In this context, polymer–ceramic
58â•…
CHAPTER 5â•… Friction
composites with improved physical properties can exhibit better wear resistance, which is discussed in Chapter 15 of this book. Some materials that have layered structure, such as graphite and MoS2, can exhibit friction as low as 0.05 and are the most commonly used solid lubricants. Their low friction originates from the fact that their basal planes are strongly covalently bonded, while these layers are weakly bonded to each other via van der Waals forces, which allows easy shear among them. Although graphite and MoS2 have very similar structure, their performance can be very different, depending on the environment. MoS2 performs best in terms of friction and wear at low humidity and ultrahigh vacuum, whereas graphite works best at high humidity. These lubricants, that is, MoS2 and graphite, have found many applications due to their form modality in which they can be used. Namely, they could be made in the form of a powder, a thin film, or as an additive in conventional lubricants or coatings. Most soft solid lubricants, which lubricate effectively, form an adherent transfer film on the surface, so that, after a short running-in period, such transfer film is formed. Thereafter, the real contact is established between the solid lubricant and transfer film. MoS2 can be synthesized also as nanoparticles in the form of fullerenes, nanowires, or nanotubes, and several potential applications are being explored. Mainly, these nanoparticles can be used in oils and greases or various coatings as additives. Another material, similar to MoS2, that is, WS2 is also becoming a promising nanoparticle additive material, also typically reducing friction. Although the mechanisms of friction reduction are still under discussion, the exfoliation into sheetlike low-shear layers and ball-bearing properties of nanoparticles are the most typical suggestions.
5.5 CLOSING REMARKS From the discussion in this chapter, it is clear that friction is a system-dependent property; that is, friction of an identical mating material combination would be different in a different environment or under a different combination of operating parameters (load, sliding velocity, etc.). Similarly, the friction of a given material under similar operating and environmental conditions would be different if the mating material were replaced by a different material. Therefore, a general statement, such as “the friction of steel is 0.5,” is meaningless, unless the operating parameters and mating material are specified. Since the frictional response is influenced only by the surface properties, an intelligent approach of putting a solid lubricant coating on an otherwise high-friction surface can provide low friction at a tribological contact. Typically, the frictional response of a mating material combination can be divided into two stages: running-in period and steady-state period. Since the transition from running-in period to steady-state period takes place over a shorter time scale from the start of asperity–asperity interactions, it is the steady-state frictional response that would be more important in the context of tribological applications. Also, it should be clear from the discussion in this chapter, as well as from later chapters, that it is not appropriate to correlate friction and wear properties for a given
â•… 59
REFERENCES
tribosystem; that is, a tribosystem can experience low friction but high wear rate or, alternatively, a tribosystem while exhibiting high friction can experience low wear rate.
REFERENCES 1â•… B. Bhushan. Principles and Applications of Tribology. John Wiley & Sons, New York, 1999. 2â•… G. Amontons. De la résitance causée dans les Machines. Memoires l’Académie Royale A (1699), 257–282. 3â•… C. A. Coulomb. Théorie des Machines Simples, en ayant regard au Frottement de leurs parties, et à la Roideur des Cordages. Mem. Math. Phys. X Paris (1785), 161–342. 4â•… B. Bhushan and A. V. Kulkarni. Effect of normal load on microscale friction measurements. Thin Solid Films 278 (1996), 49–56, 293, 333. 5â•… D. H. Buckley. Surface Effects in Adhesion, Friction, Wear and Lubrication. Elsevier, Amsterdam, 1981. 6â•… I. M. Hutchings. Tribology: Friction and Wear of Engineering Materials. CRC Press, Boca Raton, FL, 1992.
CH A P T E R
6
FRICTIONAL HEATING AND CONTACT TEMPERATURE During the last few decades, the tribological properties of various engineering materials have been extensively investigated with a major aim of evaluating their tribological potential as well as understanding the mechanism of material removal from contact interfaces. However, an in-depth understanding requires a detailed study of physicochemical and thermomechanical interaction occurring at the contacting interfaces, which has been studied to a noticeable extent.1–14 In this chapter, the interface (contact) temperature of engineering surfaces (multiple asperity model) exposed to sliding movement is discussed.
6.1 TRIBOLOGICAL PROCESS AND CONTACT TEMPERATURE When two bodies in contact are in relative motion to each other, work is done against friction. This work is mostly dissipated in the form of heat at the contact. The contact surface temperature strongly depends on the size and shape of the real contact area, along with the friction coefficient, normal load, sliding velocity, and thermal properties of the contacting bodies. Frictional heating, which occurs only in the real contact regions, also results in relatively steep temperature gradients in the subsurface layer. It has been shown in the past that these gradients can be of the order of few hundred degrees Celsius per micrometer and can go in depth only for several microns, as schematically shown in Figure 6.1. The nature of the subsurface temperature distribution depends on the thermal properties, such as diffusivity, heat capacity, and conductivity, and also on sliding velocity and loads. Friction and wear during sliding can be considered as broadly dependent on the dissipated frictional energy at the contact. The friction between two sliding bodies increases the contact temperature and generates energy that is subsequently spent in the form of deformation, cracking, or tribochemical reactions. The dissipated energy can be calculated using the following relation:
Ed = µWvt ,
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
60
(6.1)
â•… 61
6.2 Concept of “Bulk” and “Flash” Temperature
Surface
r ~2–5 µm
~10 µm Near-surface zone
Figure 6.1â•… A schematic of phase-transformed material formation at the asperity contacts due to extreme surface temperatures and extreme thermal gradients, according to experimental evidence presented in References 15 and 16. Normal load
Velocity
Sliding solid E2,n2
Extremely high temperature
Stationary solid, E1,n 1
Figure 6.2â•… Schematic of a sphere making relative movement on a nominally flat surface. E and v are elastic modulus and Poisson’s ratio, respectively, of the contacting solids.
where μ is the average coefficient of friction (COF), W is the normal load, v is the fretting velocity, and t is the total testing duration. The dissipated energy for any sliding wear experiment can be calculated using Equation 6.1. The dissipated energy can activate different processes involved in wear. This energy will participate in mechanical transformations, tribochemical phenomena, third-body transformation, and thermal phenomena.16,17
6.2 CONCEPT OF “BULK” AND “FLASH” TEMPERATURE In any tribological contact, forces are usually transmitted from a moving body to a stationary one or vice versa (see Fig. 6.2). At the sliding contact surfaces, transverse force acts only across asperities and the contacting surfaces experience either elastic or plastic deformation under tribological load depending on material parameters (E, H) or surface roughness parameters (asperity size and shape). The contact stress and distribution also depend on friction and the contact configuration (circular for sphere
62â•…
CHAPTER 6â•… Frictional Heating and Contact Temperature
vs. spherical or rectangular for crossed cylinders, etc.). In much of the published literature, it has been critically pointed out that contact temperature significantly influences oxidative wear in the case of metals,17,18 and tribochemical wear for conventional materials, as well as ceramics19–24 and composite materials.22,25 This is because, on the atomic scale, mechanical and chemical phenomena are generally dependent on the thermal energy available to assist or activate these processes. The temperature dependence of the microstructure and the mechanical and physical properties of the contacting solids affect the friction and the wear process. Mechanical properties of many materials as well as their lubricating properties can exhibit degradation due to rise in interface temperature. The temperature at the interface is defined by two quantitative parameters: bulk temperature and flash temperature. Although there is still some misunderstanding in these definitions, in the vast majority of the literature, the temperature of the contact area, which does not depend on the instantaneous frictional heat generation, is called the bulk temperature, whereas the rise in temperature at the contact due to frictional heat is called the flash temperature. One key problem here is the definition of the contact area. Namely, most of the frictional heat is dissipated from the contact into the bulk through the contact area. However, as we know, the contact area is not as large as the apparent contact but is smaller, that is, equals the real contact area. To illustrate this, the tribological surface interaction at microscale is shown in Figure 6.3. Such interaction should therefore be described by the multiple asperity interaction of two rough surfaces. So, the major question is the size of the real contact area during sliding. This is widely unknown but, unfortunately, affects the results of contact temperature calculations significantly.16 Namely, the models for determining real contact area differ and there exists no single accepted validated model. In addition to this, some contact temperature calculation models use their own built-in contact area model. So, it is difficult to compare several contact temperature calculation models with the same input parameters. Moreover, it is very questionable how the “contact” temperature differs at the asperities and at the other points of contact. Therefore, in many reports, the temperature calculations can differ significantly also
Asperity curvature Sliding solid B
Asperity height
Stationary solid A
Reference line
Figure 6.3â•… Schematic illustration of multiple asperity contacts at a tribological interface.
6.3 Importance and Relevance of Some Ready-to-Use Analytical Models
â•… 63
because some researchers use the apparent contacts, while others use the real contact area for their temperature calculations. Furthermore, at real contact points, the true value of contact pressure can be obtained by dividing the nominal load by the ratio of real to apparent contact area. This is represented by the shaded area in Figure 6.2. Each two-dimensional contact point serves as an asperity that can be loaded or not, depending on the distribution of the asperities, which, however, is very difficult to properly calculate. These problems have been discussed in detail, and presented numerically and experimentally, by Kalin and coworkers.13,16,26 It will be shown later that, for a stationary body in contact with a moving body, maximum flash temperature depends on the combination of different factors: normal load, sliding speed, COF, and thermal properties.13,16,26 Contact temperatures are therefore expressed in terms of the rate of supply of heat, the size and speed of the heat source, and the thermal properties of the materials. Bhushan27 proposed another model, including specific consideration of both transient and steady-state heat flow of specimen geometry, and of the thermal resistance of the contact between the slider and its clamping system. Based on the proposed expression, they introduced the concept of temperature maps in order to easily visualize the model. Temperature maps are diagrams with the normal load and velocity as axes, on which temperature contours are plotted. The temperature can be found for each load and velocity combination through this map. They have calculated the contact temperatures for several contacts (material pairs), considering normalized velocity and pressure. This is useful for a quick estimation of the contact temperatures; however, one should be aware of the assumptions of the model, and consider them in full; otherwise, significant differences and misunderstanding of the results can appear.13,16,26 A typical temperature map for ceramic–ceramic contact is shown in Figure 6.4. It can be observed that the temperature increases with increasing velocity as well as normal load.
6.3 IMPORTANCE AND RELEVANCE OF SOME READY-TO-USE ANALYTICAL MODELS It is well recognized that high temperature (more than ambient environment temperature) is produced at multiple asperity contacts (see Fig. 6.3) for short durations (∼10−3 seconds or less, depending on sliding speed) and occur at small distances (∼10−4â•›m or so, depending on load and asperity size distribution). However, it is difficult to measure contact temperature with an experimental tool with any reasonable accuracy,16 and, in the interpretation of the friction and wear results, recourse is generally made to estimating their magnitude using the theory originally formulated by Blok.1 In the early days of developing these models, many attempts were made to come to a reasonable conclusion, but most of the models varied significantly. Some concepts considered mostly the physics of the contacts and heat transfer, while some emphasized the material aspects and experimental evidence. Many experimental and theoretical studies have been made and many publications are available. However, as said initially, the major problem in properly determining the contact
64â•… 1.00
CHAPTER 6â•… Frictional Heating and Contact Temperature
Log (Velocity, m/s)
–3.00 Ball ALUMINA Ro = Ra = To =
on on 1.00 20.0 20
Flat ALUMINA mm µm °C
1.98
Ho = k= a= Uo = L1b = n=
Ball 12.0 30.0 9.80 0.80 10.0 100.0
Flat 12.0 30.0 9.80 c1 = L2b = Vc =
5.08
GPa W/m/K mm2/s 250.0 arctan x 4.50 m/s
–5.00
Log (F/An, MPa)
Log (F/HoAn)
BULK TEMPERATURE 2350 CONTOURS (°C) T1 = 50 T2 = 150 7 T3 = 300 T4 = 450 50 T5 = 650 150 9 T6 = 950 300 8 T7 = 1300 450 T8 = 1750 650 1300 T9 = 2350 950 1750 FLASH TEMPERATURE 2350 C –1.00
Log (Normalized Velocity)
300 150 50 C
4.00
–0.92
Figure 6.4â•… Typical temperature map illustrating the operating parameter dependent evolution of flash temperature.28
temperature arises from the unknown real contact area, real contact “geometry,” actual material properties at different temperatures, and even the distribution of shear stress (friction) at different contacting asperities.13,16,26
6.4 REVIEW OF SOME FREQUENTLY EMPLOYED READY-TO-USE MODELS All the analytical models are developed on the basis of a set of assumptions as well as considering the ideal contact behavior. At the tribocontacts, an asperity of solid B is sliding with harmonic oscillations on the stationary surface of solid A (see, e.g., Figs. 6.2 and 6.3). A normal force is transmitted under low contact stress from the oscillating to the stationary solid across one asperity. For derivation of the mathematical model, it is physically the same if solid A oscillates and B is stationary or vice versa. Based on such considerations, theoretical calculation of flash temperatures was attempted by several researchers. A few significant efforts among them are Archard’s and Holm’s average and maximum flash temperature model,2,4,13,16 Tian and Kennedy’s average and maximum flash temperature model,29 Greenwood and Alliston-Greener’s average flash temperature model for fretting,30 and Ashby, Abulawi, and Kong’s bulk and flash temperature model.28 These models differ in their physical, dynamic, and geometrical assumptions.13,16 For example, the geometry that was considered for calculating contact temperature by Ashby et al. is shown in Figure 6.5. Normally, one contacting body is considered to be at rest and the other to be moving.
6.4 Review of Some Frequently Employed Ready-to-Use Models
â•… 65
F HEAT SINK AT T9 PIN
l1
HEAT FLUX
SINK CONTACT 1 PROPERTIES Ac1, hc1 PIN: PROPERTIES H1, K1, a1 DISK: PROPERTIES H2, K2, a2
DISK 2ro An = πro2
MOTION, v
Figure 6.5â•… Schematic of frictional heat generation in the pin-on-disk configuration.28
6.4.1â•… Assumptions in Various Models In various analytical models for computing temperatures at tribological interfaces, the following assumptions are considered: 1. Elastic Hertzian contact between solids is considered for all stress–strain analysis. 2. The heat generated at contact is only due to the dissipation of frictional energy. 3. Sliding of two rough surfaces is considered as sliding of one rough surface, having equivalent or composite roughness, over a smooth surface. 4. A finite temperature gradient perpendicular to the contacting surface is considered. 5. Maximum temperature attained at a contact point is assumed to be the same for both the mating solids. 6. For a fretting contact, heat dissipation to the surroundings is negligible. 7. Flash temperature is the average of maximum temperature attained at an asperity–asperity contact. 8. Steady-state condition for heat generation is assumed.
6.4.2â•… Model Descriptions and Implications Sliding between two surfaces leads to heat generation at the contacting surface. During the tribological test, heat is extracted from the contact zone by thermal conduction, convection, and radiation. The dominant mode of heat dissipation from the contact will depend on the tribological environment. For dry sliding, heat will be dissipated by conduction to both mating solids. In contrast, heat dissipation can be via convection or evaporation in a cryogenic environment (such as liquid nitrogen).31,32 For a ball-on-flat contact configuration, the frictional energy dissipation can be computed by the following formula27:
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CHAPTER 6â•… Frictional Heating and Contact Temperature
q = µPa V ( Aa /Ar ),
(6.2)
where q ╯=╯heat generated per unit area (J/m2), μ ╯=╯COF, Pa ╯=╯load per unit apparent area (P/Aa) (N/m2), P ╯=╯load (N), Aa╯=╯apparent area of contact (m2), Ar ╯=╯real area of contact (m2), V ╯=╯Sliding speed (m/s). From Equation 6.2, it is evident that the higher the COF or sliding velocity, the more will be the heat generated at the sliding interface and, accordingly, the higher will be the flash temperature. 6.4.2.1 Archard Model. Archard2 was the first one to propose a model for contact temperature calculation, based on the assumption that the heat is generated at the area of true contact and that this heat is conducted away into the bulk of the rubbing solids (see Fig. 6.6). The calculation of the flash temperature in his model was done assuming that maximum temperature occurs when the maximum possible load is concentrated at the smallest possible area; that is, when the total applied load is borne by plastic deformation at a single area. The temperatures are calculated on the assumption that the heat is generated at the area of true contact and that this heat is conducted away into the bulk of the sliding bodies. The sliding system is approximated to a finite long band heat source with uniform distribution of heat intensity (see Fig. 6.6). Using the Peclet number (L╯=╯Va/2χ), the speed criterion for heat sources moving within different velocity ranges can be defined and the following expressions for the flash temperature can be arrived at:
Time t ago x'
Vt dx'
V t
l
x Heat source at origin at t=0
z'
z
Figure 6.6â•… Schematic of a moving band heat source (length 2/) on a semi-infinite solid body (stationary), illustrating the concept used in determining the contact temperature.27
6.4 Review of Some Frequently Employed Ready-to-Use Models
θm θm θm θm
â•… 67
= 0.25 NL at low speeds; L < 0.1; = 0.25βNL at moderately low speeds; 0.1 < L < 5; = 0.345 NL1 / 2 at high speeds; L > 100; = 0.435γNL1 / 2 at moderately high speeds; 5 < L < 100;
(6.3) (6.4) (6.5) (6.6)
µg πpm P1 / 2 v and L = , 1/ 2 J ρs 2χ (πpm )
(6.7)
where
N=
and where μ is the COF, g is gravitational acceleration, J is the mechanical equivalent of heat, pm is flow or yield pressure, s is specific heat, ρ is density, P is normal load, v is the speed of sliding, χ is thermal diffusivity, β is a constant that ranges from a value of about 0.95 at Lâ•›=â•›0.1 to about 0.5 at Lâ•›=â•›5, and γ is a constant that ranges from a value of about 0.72 at L╯=╯5 to about 0.92 at Lâ•›=â•›100. From Equations 6.3–6.7, it can be observed that flash temperature increases linearly with sliding speed and is proportional to (load)0.5 at low sliding speeds, whereas it is proportional to (velocity)0.5 and (load)0.25 at higher sliding speeds. 6.4.2.2 Kong–Ashby Model. From a phenomenological point of view, high sliding speed causes an increase in the contact temperature (flash temperature), which often helps oxidation reaction to take place. Kong and Ashby33 proposed a different model for flash temperature (Tf) calculations:
µPv 1 Tf − Tb′ = , Ar k1 k2 + l1f l2 f
(6.8)
where
Ar (Tb − To ) , An µFv 1 Tb = To + , An k1 k2 + l1b l2 b Tb′ = Tb −
Ar P = , An Ps
l1f =
rj n2πχ1 tan −1 , π1 / 2 rj v
l2 f =
rj n2 πχ2 tan −1 , 1/ 2 π rj v
(6.9)
(6.10)
(6.11) 1/ 2
(6.12)
1/ 2
(6.13)
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CHAPTER 6â•… Frictional Heating and Contact Temperature
2 Pr rj = r0 1 − 0 + 1 Ps ra An H 0 Ps = , (1 + 12µ 2 )1 / 2 H0 = least of ( H1, H2 ).
−1 / 2
,
(6.14) (6.15) (6.16)
In the preceding expressions, Tf is the flash temperature (K); An is the nominal contact area; Ar is the real contact area; χ1 and χ2 are thermal diffusivities of 1 and 2, respectively; P is the normal load; Ps is the seizure load; H is hardness; k1 and k2 are thermal conductivities of 1 and 2; l1f and l2f are equivalent linear heat-diffusion distances for surfaces 1 and 2 for flash heating; l1b, l2b are equivalent linear heatdiffusion distances for two mating surfaces, bulk heating (m); n is a measure of the lifetime of a contacting asperity; ra is the radius of a single isolated asperity junction; rj is the radius of a contact junction that can be made up of many unit asperities; r0 is the radius of nominal contact area; Tb′ is the effective bulk temperature for bulk heating; v is the sliding velocity, and μ is the COF. In the Kong–Ashby model, flash temperature is proportional to the (load)0.5, but dependence on sliding speed is complex in nature. In principle, the Archard and Kong–Ashby models predict similar functional dependence on load at low sliding speed, but they differ at higher sliding speeds and with respect to sliding dependence of flash temperature.
REFERENCES ╇ 1â•… H. Blok. The flash temperature concept. Wear 6 (1963), 483–494. ╇ 2â•… J. F. Archard. The temperature of rubbing surfaces. Wear 2 (1958), 438–455. ╇ 3â•… B. J. Griffiths. White layer formations at machined surfaces and their relationships to white layer formation at worn surfaces. ASME J. Tribol. 107 (1985), 165–171. ╇ 4â•… J. F. Archard. The temperature of rubbing surfaces: Part 2, the distribution of temperatures. Wear 128 (1988), 1–17. ╇ 5â•… J. Pezdirnik, J. Vizintin, and B. Podgornik. Temperature at the surface and inside an oscillatory sliding microcontact—Theoretical part. Tribol. Int. 32 (1999), 481–489. ╇ 6â•… B. Vick, M. J. Furey, and K. Iskandar. Theoretical surface temperatures generated from sliding contact of pure metallic elements. Tribol. Int. 33 (2000), 265–271. ╇ 7â•… A. Cameron, A. N. Gordon, and G. T. Symm. Contact temperatures in rolling/sliding surfaces. Proc. R. Soc. Lond. A 286 (1964), 45–61. ╇ 8â•… F. F. Ling. Surface Mechanics. Wiley Interscience, New York, 1973. ╇ 9â•… F. F. Ling and S. Pu. Probable interface temperature of solids in sliding contact. Wear 7(9) (1964), 23–34. 10â•… H. A. Francis. Interfacial temperature distribution within a sliding Hertzian contact. ASLE Trans. 14 (1970), 41–54. 11â•… J. A. Greenwood. An interpolation formula for flash temperatures. Wear 150 (1991), 153–158. 12â•… G. Liu, Q. Wang, and Y. Ao. Convenient formulas for modeling three-dimensional thermo-mechanical asperity contacts. Tribol. Int. 35 (2002), 411–423. 13â•… M. Kalin and J. Vizintin. Comparison of different theoretical models for flash temperature calculation under fretting conditions. Tribol. Int. 34 (2001), 831–839. 14â•… B. Vick and M. J. Furey. A basic theoretical study of the temperature rise in sliding contact with multiple contacts. Tribol. Int. 34 (2001), 823–829.
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15â•… F. Vodopivec, J. Vižintin, and B. Šuštaršicˇ. Effect of fretting amplitude on microstructure of IC-1.5 Cr Steel. Mater. Sci. Technol. 12 (1996), 355–360. 16â•… M. Kalin. Influence of flash temperatures on the tribological behaviour in low-speed sliding: A review. Mater. Sci. Eng. A 374 (2004), 390–397. 17â•… M. Surender, B. Basu, and R. Balasubramaniam. Wear characterization of electrodeposited Ni-WC composite coatings. Tribol. Int. 37 (2004), 743–749. 18â•… A. Choubey, B. Basu, and R. Balasubramaniam. Tribological behaviour of Ti-based alloys in simulated body fluid solution at fretting contacts. Mater. Sci. Eng. A 379 (2004), 234–239. 19â•… B. Basu, R. G. Vitchev, J. Vleugels, J. P. Celis, and O. Van Der Biest. Influence of humidity on the fretting wear of self-mated tetragonal zirconia ceramics. Acta Mater. 48 (2000), 2461–2471. 20â•… B. Basu, J. Vleugels, and O. Van Der Biest. Fretting wear behaviour of TiB2-based materials against bearing steel under water and oil lubrication. Wear 250 (2001), 631–641. 21â•… J. Vleugels, B. Basu, K. C. H. Kumar, R. G. Vitchev, and O. Van Der Biest. Unlubricated fretting wear of TiB2 containing composites against bearing steel. Metall. Mater. Trans. A 33(12) (2002), 3847–3859. 22â•… B. Basu, J. Vleugels, and O. Van Der Biest. Fretting wear behaviour of advanced ceramics and cermet against alumina. J. Mater. Res. 18(6) (2003), 1314–1324. 23â•… B. Basu, J. Vleugels, M. Kalin, and O. Van Der Biest. Friction and wear mechanism of SiAlON ceramics under fretting contacts. Mater. Sci. Eng. A 359 (2003), 228–236. 24â•… D. Sarkar, S. Ahn, S. Kang, and B. Basu. Fretting wear of TiCN-Ni cermet: Influence of secondary carbide content. P/M Sci. Technol. Briefs 5(2) (2003), 5–11. 25â•… B. V. Manoj Kumar, B. Basu, V. S. R. Murthy, and M. Gupta. The role of tribochemistry on fretting wear of Mg-SiC particulate composites. Composites A 36 (2005), 13–23. 26â•… M. Kalin. High temperature phase transformations under fretting conditions. Wear 249 (2001), 172–181. 27â•… B. Bhushan. Principles and Applications of Tribology. Wiley-Interscience, John Wiley & Sons, New York, 1999, 447. 28â•… M. F. Ashby, J. Abulawi, and H. S. Kong. Temperature maps for frictional heating in dry sliding. Tribol. Trans. 34 (1991), 577–587. 29â•… X. Tian and F. E. Kennedy Jr. Maximum and average flash temperatures in sliding contacts. ASME J. Tribol. 116 (1994), 167–174. 30â•… J. A. Greenwood and A. F. Alliston-Greiner. Surface temperatures in a fretting contact. Wear 155 (1992), 269–275. 31â•… R. Khanna and B. Basu. Sliding wear properties of self-mated yttria-stabilised tetragonal zirconia ceramics in cryogenic environment. J. Am. Ceram. Soc. 90(8) (2007), 2525–2534. 32â•… R. Khanna and B. Basu. Low friction and severe wear of alumina in cryogenic environment: A first report. J. Mater. Res. 21(4) (2006), 832–843. 33â•… H. S. Kong and M. F. Ashby. Friction-heating maps and their applications. Mater. Res. Soc. Bull. 16(10) (1991), 41–48.
CH A P T E R
7
WEAR MECHANISMS This chapter discusses some of the important wear mechanisms, including abra sion, adhesion, tribochemical–oxidation, fretting, and erosion. While a major part of this chapter focuses on the physics of material removal mechanisms in terms of the physicochemical changes at tribocontacts, efforts have also been made to discuss the quantified correlation of material removal with material parameters (hardness, toughness, elastic modulus) as well as operating parameters (load, sliding velocity, etc.).
7.1 INTRODUCTION In selecting and designing materials for various engineering applications involving relative motion, resistance against material damage or removal of material from contacting surfaces is an important—and sometimes the key—factor. Conventionally, wear can be defined as the phenomenon of progressive damage resulting in a loss of material.1 Very often high wear and high friction coincide, but this is not always the case and neither is it physically justified. The two phenomena should therefore be evaluated separately; however, one still should simultaneously pay attention to both of them, because they provide complementary information on tribological systems. Nevertheless, worn surfaces often cause high friction and so increase power losses, which is another consequence of wear. Moreover, a negative consequence of wear is not just the need to replace some mechanical part and the cost for it; one of the most important negative effects of wear is the costs associated with the mainte nance and production process. Thus, wear is closely related to other technical fields, such as maintenance and technical diagnostics; it is this relation that brings particu larly relevant importance to the field of wear—because of its significant economic impact. Furthermore, the performance of any tribological part that is fully or partially worn will deteriorate, often causing vibrations and heating and thus leading to lower product quality, surface damage, and decreased machine efficiency. Thus, wear is certainly a phenomenon that must be taken care of, when observed to any detectable extent. A practical solution, where possible, is to reduce wear (and friction) with lubrication. Lubricants tend to separate the surfaces and relative movement at
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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7.1 Introduction
â•… 71
contacting mating solids through the low-shear layers. The next chapter of this book deals with some fundamentals of lubrication, as well. However, it is important to note at this point that wear behavior completely changes when surfaces are lubri cated. Sometimes, effects that are negative for dry wear become positive (useful) under lubricated conditions. In addition, in the absence of a lubricant, such as oil or grease, tribological conditions are often called “dry.” However, practically every surface adsorbs layers of different species from the environment or oxide layers on top of the bulk material that can potentially act as lubricants, and so their presence changes the properties of bulk materials. When the wear behavior of a given “mate rial” is analyzed, this should certainly be considered. It is worth pointing out that wear is not always an undesirable process. Namely, all manufacturing processes that involve removal of material, such as cutting, drill ing, grinding, and polishing, are essentially wear processes. An improved under standing of wear mechanisms and material properties, together with technological parameters, can lead to substantial improvements and tailoring of such machining processes. Wear can also be classified or diversified based on the type of relative motion of contacting solid surfaces. The two most common motions are sliding and rolling, although spinning and fretting (as a special case of sliding) can have quite different effects on wear. In practical applications, however, most of the contact kinematics is not so clear and, in the majority of cases, simultaneous sliding and rolling occurs. For example, the most commonly used and studied tribological machine systems include bearings, gears, and cam-follower contact. It should be noted that pure rolling is a very rare situation and can occur only at some points (or a line) within the contact, even if the motion is macroscopically under rolling conditions. This is because of the differences in the radius of rotation (at the same angular velocity) of the contacting points of a certain body, causing a difference in the translatory veloc ity of these points. Therefore, most of the contacting junctions between two bodies will not have the same translatory velocity, and they are thus in a microsliding situ ation. This is important, because it suggests that sliding, even though it is sometimes microsliding, is the predominant type of motion between tribological parts. Some further exemptions and specifics of sliding are discussed in this chapter. Spinning, which is a third independent kinematic motion, is a rather rare situation in tribologi cal applications. There exist several mechanisms of wear, which cause damage to surfaces and subsurface regions to various extents. A survey of tribological properties of different materials indicate that some forms of wear cause faster material removal and more severe consequences, so in this respect are more “dangerous.” For example, abrasion and adhesion are typical high wear rate mechanisms that should be eliminated when present; otherwise, they can lead to catastrophic failure in rather short times. Some other forms, although considered less dangerous and indeed occurring at lower rates, can in some particular and nonconventional cases also result in extreme wear rates, for example, diffusion or tribochemical wear. On the other hand, other mechanisms usually occur with lower rates, such as cavitation, corrosion, erosion, or fretting. These are sometimes also called minor mechanisms of wear. Nevertheless, in some applications, failure will always occur “only” due to these forms. So, in these cases
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we have to take precautions for these particular “minor” mechanisms. For example, turbines often and predominately suffer from cavitation, valves from erosion, and wire ropes (cable cars, elevators), bearings, and gear seats from fretting. Thus, these wear mechanisms will most probably determine the useful life of these components. Therefore, even if the wear rate is lower, it does not mean it could be neglected. The mechanisms of wear are highly complex and, very often, more than one mechanism is acting at the same time. So it is important to understand the properties of individual mechanisms to be able to isolate them and recognize their interdepen dencies. Only after postulating the proper “diagnosis” of wear mechanisms is the proper “action” for this mechanism proposed. Consequently, the wear can be properly “cured”—removed or prevented. Therefore, information about mechanisms is crucial, and it is thus instructive to briefly review the fundamental wear mecha nisms. We discuss just the most common wear mechanisms here and in later chap ters. However, the reader can refer to a large number of books and articles that give a more comprehensive presentation and focus on many specifics of wear mechanisms. Moreover, in the past there were many attempts to model wear through dif ferent equations that would predict the expected amount of wear. From a practical point of view, this would be very useful because it would allow prediction, as well as better selection and optimization of specific tribological systems. However, the problem is the complexity of the wear process and the simultaneous action of several wear mechanisms in specific situations, which is hard to incorporate in any theoreti cal model. Several models are based on material properties but somehow underes timate environmental and operating conditions. Other models consider more working parameters but include fewer material properties.2 Finally, there is in situ material and contact modification, as well as wear debris formation, which make the situation more complex. At least, these models show the phenomenological behavior and effects of materials or conditions, so they are a handy tool for better design of tri bological contact. Hopefully, based on the information on materials, their wear mechanisms, and wear models provided in this book, one can answer the following questions with greater confidence: 1. From the microstructural aspects, how can one identify operative wear mechanisms? 2. Once the dominant wear mechanism is identified, how can one correlate it with material properties and operating parameters? 3. From the design perspective, what should be the guiding parameters to develop materials for better wear resistance?
7.2 CLASSIFICATION OF WEAR MECHANISMS Different materials as well as experimental parameters constitute various factors that influence the wear mechanisms at a sliding contact, such as adhesion, surface fatigue, tribochemical reactions, and abrasion.1 It is known that sliding of metallic materials is often characterized by plastic deformation and oxidative wear, whereas that of
7.2 Classification of Wear Mechanisms
â•… 73
ceramics and their composites is influenced by brittle fracture and tribochemical effects.3 However, the simultaneous occurrence of several mechanisms often com plicates the wear process of different materials, requiring systematic and selective tribological experiments and subsequent analyses. In general, wear loss as a function of time or of sliding distance is influenced by different operating conditions (sliding speed, load, temperature, and environment) and material parameters, characterized by surface finish, hardness, toughness, microstructure, geometry, interfacial ele ments, and so on. All these parameters make the progress of wear complex, as a number of mechanisms act simultaneously for a given tribocouple.
7.2.1â•… Adhesive Wear Adhesion is caused by the surface interactions and welding of asperity junctions at the sliding contact. In particular, interfacial adhesion may be due to ionic, covalent, metallic, hydrogen, and/or van der Waals bonds. If the junction is stronger than the bonds in the bulk, a part of the weaker material will detach and remain attached to the counterbody. This material is called the transfer film. In this way, one surface loses some of its material, resulting in irregular holes or pits, while the countersur face gains some material in the form of transfer film. It is quite understandable that sliding of such two bodies cannot be smooth anymore and that the next collision of transfer film will only be more severe. So, the sliding interface becomes rough and the component loses its dimensions and tolerances. The transfer film may grow in time and can be eventually removed as wear particles; it is often also work-hardened and, in this way, results in abrasive actions as well. The adhesion theory of friction and wear during sliding can be applied to the formation of wear particles in the fol lowing sequence: (1) loaded contact of single asperities on a pair of rubbing surfaces; (2) the formation, growth, and failure of adhesive junctions; (3) transfer of the mate rial to the countersurface; and (4) detaching of transfer material as loose wear particles. According to Archard’s wear theory,4 the volume per unit sliding distance (V) in a plastically deformed sliding contact is related to the normal load (W) and hard ness (H) of the softer body as V = KW / H ,
(7.1)
where K is a dimensionless wear coefficient indicating the probability of forming wear particles. To compare the wear behavior of different material systems, Lancaster3 proposed a wear coefficient, usually known as specific wear rate k, k = V / W.
4
(7.2)
Archard proposed that shearing will occur at the interface or in the weakest region in one of the sliding bodies. In general, interfacial adhesion strength is expected to be small compared with the breaking strength of surrounding local regions. Thus, rupture during shearing occurs at the interface (designated as path 1 in Fig. 7.1) in most of the contacts and no wear occurs in that sliding cycle. In a small fraction of contacts, rupture may occur in one of the two bodies (path 2 in Fig. 7.1) and a small fragment or lump (shaded region in Fig. 7.1) may become attached to the other surface.5
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2
1
Figure 7.1â•… Schematic illustrating the possibilities for the rupture of adhesive bonds (paths 1 and 2) during shearing of a tribological interface.5
B
A
B
A
B
A
C
D
C
C
C
B C C E
A D
D
E
F
Figure 7.2â•… Schematic showing the stages (A–F in order) in the detachment of fragment of a material due to plastic shearing of successive layers of asperity contacts.5
Kayaba and Kato6 presented another mechanism, which involves plastic flow at an asperity tip resulting in the detachment of wear particles. This is presented in Figure 7.2. According to this theory, plastic shearing of successive layers based on a slip line field (along the plane AC) occurs in conjunction with the propagation of a shear crack (AD), along which the fragment detaches (see Fig. 7.2). Sasada7 further explained that the formation of debris fragments is followed by adhesive transfer to the countersurface. The wear fragments may remain adhered to a surface, transfer to the mating surface, or become detached as large loose wear particles. These particles may be of roughly equal size in each dimension or may be flattened or elongated in the sliding direction.7 Further, during this process, surface asperities undergo plastic deformation and strain hardening. This leads to a factor of 2 increase in hardness over the bulk hardness. Adhesion is strong in contacts of chemically similar materials, those having large mutual solubility, where the atoms of the two bodies can easily interact. Metals, especially those with weak contaminant layers (oxides) or those that do not even form oxides, experience the strongest adhesive wear. Another observed effect that strongly relates to adhesion is the material’s toughness. Materials that tend to deform will easily attach to each other at the larger contact area, and, in this way, a larger contacting surface will be generated, where the bonds that need to be broken during sliding will form. A phenomenon that describes this effect of material deformation that increases the contact area and thus the amount of adhesive bonds is called junc tion growth. This is much more pronounced with nonreactive tough materials that can plastically deform than with brittle materials. So, the way to prevent adhesion is quite straightforward: that is, use chemically dissimilar materials that do not plasti cally deform (provide small real contact area) or use any form of surface “contami nants” (i.e., lubricants, oxide layers, low-surface-energy coatings, etc.).
7.2 Classification of Wear Mechanisms
â•… 75
7.2.2â•… Abrasive Wear Abrasion is a form of wear that usually produces high wear rates. Abrasion itself is not a unique wear mechanism; rather, it represents a common type of damage based on different physical mechanisms or destructive actions to materials. This depends on the materials and their properties as well as the dynamics of the contacts. Typically, abrasion is associated with the multiple indenting of hard asperities into a softer material. The most common type of abrasion is when hard asperities or particles slide over a softer surface, causing damage at the interface by plastic deformation. The hard particles may be in and/or between one or both surfaces in relative motion, or they may be hard protuberances on one or both of the relatively moving surfaces. Further, the hard particles may be the product of processing or of hard inclusion, may be reaction products formed at interactions of two surfaces, or may be the reac tion product of the debris and the atmosphere formed during sliding. In any case, the harder asperity deforms the softer material and causes a scratch or groove. When wear debris is formed, usually in the form of chips, this physical mechanism can be referred to as cutting mode. Sometimes, the hard asperity does not form wear debris, but only leaves a groove in the softer material, and this is known as plowing. During plowing, some material can be left at the edges of the groove, which can become new wear debris later, after more sliding passes. Very often, this type of debris undergoes plastic deformation and strain hardening occurs, with an increase in hard ness due to repeated sliding, and thus it becomes another source of abrasive debris. These two forms of abrasion are very typical for materials that behave in a nonbrittle manner, such as most metals. Another abrasive mechanism occurs when a hard asperity slides over a brittle material. Since such materials cannot sustain larger plastic deformation, they form cracks. However, depending on the severity of the contact, the cracks can go deep into the subsurface (central crack at 90°), form lateral cracks under 30°, and form small local cracks oriented in various directions. Such abrasive action can cause quite large wear losses and can damage extensive surface areas. Abrasion can also be considered for multiphase materials with soft matrix embedding a hard phase, or other materials that consist of large grains. A hard asper ity that slides over such a surface can interlock with one of these hard and/or large grains or phases. If the strength of the hard asperity is large enough, it can pull out the grain, resulting in a large pit at the surface. These areas also then act as stress concentrators and are potential sites for progression of surface damage during repeated sliding. In general, abrasive wear can be classified as two-body and three-body wear.5 When there are only two bodies in relative motion and one is substantially harder, such abrasive damage is referred to as two-body abrasion. This occurs in mechanical operations, such as grinding, cutting, and machining. In three-body abrasion, the hard abrasive particles act as interfacial third-body elements between the two moving primary bodies and are responsible for wear on either or both of the surfaces, pri marily depending on their hardness. This occurs, for example, in abrasive-free lapping and polishing. Wear is about one to two orders of magnitude smaller in
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Hard, rough surface
Soft surface Abrasive grits mounted on the top surface Soft surface (a)
Free abrasive grits Soft surface (b)
Figure 7.3â•… Schematic of a rough, hard surface mounted with abrasive grits sliding on a softer surface in the case of two-body wear (a), and free abrasive grits caught between the surfaces with at least one of the surfaces softer than the abrasive grits in the case of three-body wear (b).5
three-body abrasion than in two-body abrasion, because the third bodies act on the minimum energy principle and tend to find the “easier” way within the contact (see Fig. 7.3). Critical effects for abrasive wear also include the geometric properties of hard asperities. A simple model can show the effect of the asperity contact angle. Depending on the hard asperity angle, the deformation can actually change to three different modes: cutting, wedge formation, or plowing. The total wear volume Wv of the material displaced by all asperities in a distance x is given as9
Wv =
2Wx tan θ , πH
(7.3)
where W is the applied load, θ is the attack angle of a conical asperity, and H is the hardness of the softer surface. Of course, the toughness or brittleness of the material will also substantially influence the extent of deformation of the material and the shape of the scratch or groove. However, in addition to the shape and sharpness of the asperities, hardness appears as another key parameter in abrasion. It is clear that one material must be substantially harder than the other to cause remarkable abrasive damage. It is pro posed that at least 20% hardness difference is required for “effective” abrasion. On the other hand, if the difference is too large, the abrasive wear mechanism will change into complete plastic deformation and no wear debris will form. Thus, no loss will occur, eliminating the wear—of course, damage to the surface can still be substantial.
7.2 Classification of Wear Mechanisms
â•… 77
7.2.2.1 Abrasion of Composites. In our later chapters, we discuss the wear behavior of various composites, and it is thus appropriate to present here some specific models that can be used for these materials. The following phenomenologi cal concept will be utilized to analyze wear mechanisms in subsequent chapters. Usually, brittle solids have sharp corners and are subjected to high stress and abrasion. In high-stress abrasion condition, the sharp asperity of brittle solids follows a plastic deformation in a low load regime. Above the threshold load, brittle fracture occurs, and wear occurs by lateral cracking at a sharply increased rate.10 The thresh old, or critical load is proportional to (KIc/H)3 KIc. The ratio H/KIc is known as the index of brittleness, where H is hardness (resistance to deformation) and KIc is frac ture toughness (resistance to fracture). The lateral crack length c for a given sliding asperity contact can be expressed as11 ( E / H )3 / 5 5 / 8 c = α m1 W , (7.4) 1/ 8 K Ic H where αm1 is a material-independent constant that depends on the asperity shape. The depth d of the lateral crack is given by
E d = α m2 H
2/5
W H
1/ 2
,
(7.5)
where αm2 is another material-independent constant. The maximum volume of mate rial removed per asperity encounter per unit sliding distance is 2dc. If N asperities contact the surface with each carrying load W, then from Equations 7.4 and 7.5, the volume V of wear per unit sliding distance of the interface is given by
V = α3 N
(E / H )W 9 / 8 K Ic1 / 2 H 5 / 8
,
(7.6)
where α3 is a material-independent constant. The ratio (E/H) does not vary much for different hard brittle solids. Therefore, wear rate is inversely proportional to (fracture toughness)1/2 and (hardness)5/8. Importantly, the wear rate is proportional to (normal load)9/8, which implies that wear rate by lateral fracture increases more rapidly than linearly with the applied load as in plastic deformation. Due to this, wear debris particles are produced. The size and shape of debris particles are among the important parameters for assessing the transition from mild to severe wear. Usually, at low load regime, submicron debris particles are generated. The concept of indentation fracture mechanics can be adopted to calculate the minimum load required for abrasion-induced fracture12: 3
P* =
54.47β K Ic K Ic , πη2 g 4 H
(7.7)
where P* is the minimum load required to produce fracture from a point contact (N), β is the constant relating hardness to diagonal (2.16 for Vickers indentation), g is the geometrical constant (≈0.2), KIc is the fracture toughness of the material indented (MPaâ•›m½), and η is a constant. From Equation 7.7, it is clear that the transition load increases linearly with the fourth power of fracture toughness. Therefore, the transition to severe wear will take place at higher load, if the toughness of the material is increased.
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7.2.3â•… Fatigue Wear Surface fatigue can be roughly divided into low-cycle fatigue with the overloading in a short period of time, and more conventional fatigue caused by long-term repeated loading of the surface and subsurface regions. A general consequence of fatigue processes is the formation of cracks that subsequently lead to wear of the materials and loss of surface functionality. Fatigue of materials proceeds in the sequence of elastic and plastic deformation, work hardening and/or work softening, crack initiation, and crack propagation.13–15 A generally accepted physical explanation of long-term fatigue for elastic– plastic materials, such as metals, was proposed by Suh14 in the delamination theory of wear. In this theory, it is nicely shown that, due to repeated stress cycles, plastic deformation starts in the subsurface region at the location of maximum shear stress. Namely, it is well known that dislocations start to grow at lower loads due to shear, rather than normal stresses. Once plastic deformation in the form of dislocation is initiated, predominantly originating at some stress concentration site, such as voids and impurities, this will slowly grow with new stress cycles. The strain accumulates and cracks propagate parallel to the surface until they reach a critical size, which allows them to turn toward the surface and form a sheetlike thin wear particle just after the pass of the slider, producing a tensile stress behind the contact. The delami nation of the material through the formation of these wear debris particles causes formation of pits at the surfaces, which is widely known as a pitting process in gears, bearings, and so on. The model for the delamination theory of wear can be explained using Figure 7.4(i), which shows a crack in the subsurface region. The wear rate will depend on crack growth on the left (L) and right (R) ends, as well as on coefficient of friction (μ), which affects the shear stress and depth of the maximum shear stress, the effec tive depth (d), crack length (c), and material properties. The solution leads to two equations, defining wear rate (Eq. 7.8) and wear coefficient (Eq. 7.9):
V ∆L2 d ( ∆CL + ∆CR ) = (wear rate), S λlc 3H∆L2 d ( ∆CL + ∆CR ) K= (wear coeficient ), Wλlc
(7.8) (7.9)
where N╯=╯S/λ is the number of stress cycles, Nc╯=╯ΔL/lc is the number of cracks in the loading zone, λ is the distance between asperities, lc is the distance between cracks, ΔL is the contact length, S is the sliding distance needed to remove one layer of material, ∆CL is the average crack growth from the left side after N cycles, ∆CR is the average crack growth from the right side after N cycles, H is hardness, and W is the load. The schematic of the progress of delamination wear is shown in Figure 7.4(ii). From the materials point of view, it is clear that the wear rate will depend on the rate of crack propagation through the subsurface matrix, as well as on the number of initiating sites, that is, crack formation. So, the slower of the two processes will determine the fatigue properties.
7.2 Classification of Wear Mechanisms
–q Lc
L
â•… 79
d R c (i)
(a)
(b)
(c)
(d)
(ii)
Figure 7.4â•… Schematic illustration of (i) subsurface crack below a slider and (ii) the formation of wear sheets due to delamination in the context of the description of delamination wear.14
It should also be noted that this process occurs slowly and in the long run when the surfaces are not damaged or when they do not wear with a high wear rate. Namely, if the surface experiences wear, that is, material removal, then the subsur face shear stresses will change location and suppress crack growth. In this case, the cracks will not develop to a state to form delaminated wear debris. Therefore, this process is more typical for lubricated sliding (and rolling) contacts, which keeps the surfaces undamaged. Accordingly, the effect of nonlubricated sliding wear on fatigue failure of bulk material is often masked by other processes such as adhesion, abra sion, and tribochemical reactions, which produce wear debris at a faster pace than the cracks due to the delamination wear process. From the preceding discussion, it is clear that plastic strain accumulation both at leading and trailing edges of asperi ties, in combination, can contribute to total wear volume. The wear volume (V ) for predominantly plastic contact of the asperities on the mated surfaces of a sliding pair is given by2,15 m
V ϕ F = c N , ϕf H x
(7.10)
where c is a tribosystem-dependent constant (e.g., surface topography), x is the sliding distance, FN is the normal load, H is the hardness, ϕ f is the average plastic
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strain to failure by fracture in one loading cycle, and ϕ is the average strain in the static condition. The value of the coefficient m is about 2–3.
7.2.4â•… Oxidation and Tribochemical Wear Chemical reactions that are initiated and occur in the contacting zone of two moving bodies are a subject of tribochemistry. These reactions are often not expected or feasible under the same temperature and static-load conditions; therefore, the reac tions themselves and the kinetics of the tribochemical reactions are different from thermochemical reactions. Tribochemical wear results from the removal of the reac tion products formed in situ from the contacting surfaces. The reaction products are formed from the reactions of solid surfaces with the environment or between mating materials (see Fig. 7.5). Many times, the tribochemical reactions are associated with a positive effect on wear, also because all the additive reactions in oils rely on tri bochemistry, and these are designed for surface protection and friction reduction. These tribochemical layers act as load-bearing surfaces and possess low shearing strength, resulting in low wear and low friction. The whole boundary and in great part also mixed lubrication regime depend entirely on the formation of tribochemical films, so-called boundary films. The additives that are used in lubricants or other fluids (coolants, emulsions, etc.) are tailored for the proper chemistry to react with surfaces in specific applications and to form the tribochemical layers that could protect the surfaces. These reactions occur only under specific contact conditions, when the temperatures and loads are suitable to form an appropriate layer, with just the right performance. Thus, “the best” performance is always a compromise and a combination of materials, chem
(a)
(b) adhesive contact
protective layer
reactive environment
(c)
adhesive contact
(d) protective layer
protective layer
wear debris adhesive contact
adhesive contact
wear debris
Figure 7.5â•… Schematic illustration of the mechanisms involved in tribochemical wear.1
7.2 Classification of Wear Mechanisms
â•… 81
istry, and operating conditions. From this, it is clear that good performance needs to be well designed and understood. In all cases where this is not done, we may expect rather poor behavior, high wear, and high friction. It is thus even more important to understand that when the contacts are not lubricated—where we can select from a large variety of additives to improve the behavior—but consist of only two materials and the environment, the proper selection of materials and operating conditions is even more critical and may easily lead to very poor performance and catastrophic failures. Thus, to form “protective” boundary films in these cases, the behavior of materials under a variety of conditions is essential, and subsequent chapters present part of these cases for a selection of advanced ceramics and composites. The formation of tribochemical layers depends on many parameters: chemistry and physics of the mating materials, environment, shear, formation of clean reactive surfaces, and the contact temperatures generated due to friction heating. The aspect of contact temperature and its consequent influence on various tribochemical reac tions are further discussed in various chapters of this book. It can be understood that the tribochemical wear is strongly dependent on the kinetics of formation of surface layers and the properties that determine their removal, for example, ductility, strength, hardness, shear, and adhesion to the surface. For example, sliding in oxidiz ing atmosphere leads to formation of oxides, which are, therefore, probably the most common type of tribochemical layers formed in many applications and are very typical for ceramics and composites as well. The combination of testing parameters, such as load, speed, temperatures, humidity, and partial pressures, favors the forma tion of tribochemical oxide layers at the sliding interface. It is well known that the structure and properties of the oxides influence the wear of the materials. It has been reported that Fe3O4 results in lower wear rate than Fe2O3 oxide films. Hong and coworkers16 developed a model for quantifying the wear loss due to tribo-oxidation:
V=
A Al exp − (Q / RgTf ) ρo fo
ts ,
(7.11)
where V is the wear volume, A is the area of contact, Al is the Arrhenius constant, Q is the activation energy for oxidation, Rg is the molar gas constant, Tf is the flash temperature, ρo is the average density of the oxide in contact, fo is the mass fraction of the oxide that is oxygen, and ts is the sliding time.
7.2.5â•… Fretting Wear Fretting is one of the wear mechanisms that is many times misunderstood or is obscured behind some other more conventional mechanisms, such as fatigue, various forms of sliding wear, and corrosion. However, the reasons and possible solutions for fretting damages are very different from those of the conventional. Because of this, and the fact that in this book we refer often to fretting wear results, we feel that it is required to explain the fretting in a bit more detail at this point. This also coin cides with less information being available in the literature on various forms of fretting damage.17
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In 1911, Eden et al. first reported the occurrence of fretting on the formation of brown oxide debris in fatigue experiments of steel against steel.18 However, it was Tomlinson, in 1927, who coined the term fretting, while constructing two sepa rate machines for producing a small-amplitude rotational movement between two annuli and between an annulus and a flat, respectively.19 Tomlinson coined the term “fretting corrosion” because of the chemical reaction of steel with oxygen in ambient air, resulting in the formation of red iron oxide (α-Fe2O3) debris. However, it was found that the damage can also occur on non-oxidizing materials, such as iron and gold,20,21 and that contradicts Tomlinson’s definition. The current understanding of fretting refers to “any situation in which the contacts between materials are subjected to a low amplitude oscillatory relative motion that is lower than a few hundred µm.”22 The displacement amplitudes (typically up to 300╛µm) and relative velocities between contacting bodies in fretting are smaller than those of reciprocating sliding.2,23,24 This means that the contact is maintained over most of the tribosurface during fretting. As a result, much of the wear debris produced by fretting remains trapped at the interface, which can cause even seizure.25 The damage due to fretting can be categorized into three main forms. (1) Fretting corrosion accounts for the degradation due to chemical reactions between surface constituents and the environment that result in the formation of corroded debris.22 (2) Fretting fatigue refers to fatigue of materials due to cyclic changes of stress field22 under fretting conditions, typically in just few micrometers of sliding amplitude. (3) Fretting wear represents the surface damage, originating from the fretting process under conditions of larger amplitudes, where slip occurs over the whole or most of the contact. However, as a result of the improved understanding of the fretting concept, the three categories—fretting wear, fretting corrosion, and fretting fatigue—are referred to as fretting damage, particularly in connection with fretting regimes.22,23,26 The occurrence of fretting can be found in any technical system where contact vibrations are present or which enable small tangential relative motion. Holland27 divided an even longer list of potential fretting damage situations into two scenarios: (1) where the contacting surfaces are not designed to move relative to each other, for example, shrink fits, bolted flanges, keys, and riveted joints; and (2) where rela tive movement occurs for part of the time, for example, bearings, flexible couplings, and reciprocating cams. Some instances of fretting in engineering applications include hubs and disks press-fitted to rotating shafts, in riveted and bolted joints, between the strands of wire ropes, between the rolling elements and their tracks in stationary ball-and-roller faces, the flanges between the beveled gear and the drive shaft in gas turbine trans mitters in helicopters, damage at the femoral stems in total hip replacements, and dental restorations. The list of consequences of fretting is noteworthy as the list includes loss of helicopters, failure of power station generator motors, loss of eleva tors and cable cars due to wire rope failures, turbine disk failure in a gas turbine aero-engine, failure of wire reinforcements in radial tires, failure of a supporting joint of a railway line, and failure of an artificial hip joint. 7.2.5.1 Fretting Modes. Based on the mode of trajectories of oscillatory motion at contacting surfaces, the displacement field of an arbitrary vibration, fret
7.2 Classification of Wear Mechanisms
P
â•… 83
P
Variables: P, ∆P, f
Variables: P, d, f
2amin D (a)
2amax
(b) P
Variables: P, θ, f
θ (c)
Figure 7.6â•… Schematic illustrating principles of fretting modes.
ting is divided into three main modes28 (see Fig. 7.6 for a schematic of the three fretting modes): 1. Fretting mode I (linear mode):╇ The fretting is induced by a small-amplitude linear relative displacement with a constant frequency of oscillation. The tangential displacement at a ball-on-flat fretting contact is schematically pre sented in Figure 7.6. Most of the laboratory experiments reported in the litera ture as well as discussed in various chapters of this book are related to mode I. When the sliding is confined to an outer ring-shaped zone slip annulus sur rounding the central part (stick zone), the displacement is called partial slip. When the sliding occurs throughout the contact area, then the displacement is called gross slip. 2. Fretting mode II (radial mode):╇ In this fretting mode, the radius of the con tacting boundary oscillates between amin and amax, due to normal force oscil lations. The induced displacement is zero at the center and increases to maximum at the contacting periphery. Unlike fretting mode I, the exposure time of the slip region varies with 1/f, while the locked region remains always unexposed. Thus, there exists no gross slip in mode II. As an example, ballbearing applications and electrical contacts often encounter mode II fretting.
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3. Fretting mode III (circumferential mode):╇ The circumferential displacement amplitude increases from zero at the center to a maximum value at the contact periphery. Thus, the normal force and twisting angle can determine the maximum amplitude of displacement. Since the contact area is not exposed to the environment, occurrence of fretting corrosion is less probable. Kennedy and coworkers, however, reported for mode III the occurrence of fretting fatigue in the slip region for ceramics,29 which is not typical like the commonly used mode I fretting behavior. 7.2.5.2 Mechanics of Elastic Contacts under Fretting Conditions. To explain fretting mechanics, a simple elastic model based on Hertzian theory is used here. A stress–strain field can be very well defined at the contact of an ideally smooth ball and a flat surface. Accordingly, the contact surface of these two bodies is a circle, having a radius a, as seen from Figure 7.7. The value of a depends on the normal force FN, the radius of the ball R, the elastic modulus E, and the Poisson ratio ν:
a=
3
3 (1 − ν2 ) FN R 2E
.
(7.12)
The distribution of normal pressure over the contact area as a function of a certain radial distance r from the center is given by
p (r ) =
3FN r2 1 − . 2 πa 2 a2
(7.13)
Figure 7.7â•… Elastic model for surface contact under normal load and tangential force.23
7.2 Classification of Wear Mechanisms
â•… 85
Under elastic conditions, the contact pressure reaches a maximum at the center of the contact circle and falls to zero at the edges. If a small cyclic tangential force FT is superimposed on the normal load, some microslip can occur at the outer edges of the contact circle and a distribution of the shear traction τ(r) can be expressed in the following form: FT τ (r ) = . (7.14) 2 πa a 2 − r 2 From Equation 7.14, it is clear that the point of singularity of the tangential force is at the outer rim of the contact (r╯=╯a). This means that, at this point, the value of the tangential force reaches infinity. Since the shear traction τ(r) cannot overcome the stress defined on the whole contact area as τ (r ) ≤ µ p (r ) ,
(7.15)
it follows that, at a radius larger than r╯=╯a′╯≤╯a, slip occurs: a′ = a 3 1 −
FT . µFN
(7.16)
When the tangential force FT╯<╯μFN is applied (where μ represents a static coefficient of friction), a contact surface has in its inner part a circle of radius a′ that is in a stick zone, while the outer part of the surface in the form of an annulus is in a slip zone. Based on the preceding description, the surface traction distribution across the annulus in the slip zone (a′╯≤╯r╯≤╯a) can be defined as τ (r ) =
3µFN r2 1− 2 , 2 2 πa a
(7.17)
while, within the stick zone (r╯≤╯a′), τ(r) can be described by r2 a′ r2 3µFN 1− 2 − 1 − 2 . (7.18) 2 a a a ′ 2πa The elastic deformation of the ball and the flat will result in a tangential dis placement δ of the center of the ball relative to a fixed reference point in the flat, far from the contact zone. The displacement is expressed by
τ (r ) =
δ=
3kµ 2
3
2 FN2 3 FT , − − 1 1 µFN E2 R
(7.19)
where k is a material constant, given as a function of Poisson’s ratio ν:
k=
(1 + ν) (2 − ν) 3 2
2 . 3 (1 − ν2 )
(7.20)
From Equations 7.11–7.20 and Figure 7.7, it can be seen that, in the transient phase of applying the tangential load, microslip starts at the outer rim of the contact
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circle and penetrates inward as the slip annulus forms. The corresponding displace ment δ is also a function of the normal load FN, as indicated in Equation 7.19, and consists of elastic and slipping components. It is shown that the inner radius of the annulus approaches zero when the applied tangential force approaches the friction force FT╯=╯μFN. This is the condition for incipient gross slip over the entire contact area. 7.2.5.3 Mechanics of Elastic–Plastic Contacts under Fretting Conditions. A more general description of the fretting process compared with the elastic model included the contribution of plastic deformation. The consideration of elastic–plastic deformation behavior of the asperity junctions implies that the contact surface is divided into three zones, where the contact conditions are modified as shown in Figure 7.8. Under traction, the asperities are deformed elastically in a central stick zone, which is surrounded by a yield annulus in which the asperities have yielded plastically but have not fractured. The yield annulus, in turn, is sur rounded by a slip annulus, where the asperities are subjected to shear fracture in the same sense as in the elastic model. Therefore, in the elastic–plastic model, it is assumed that the displacement is made up of three components, that is, elastic displacement between the protrusion and the half-space (flat), plastic deformation in the bulk of the contact zone, and slip in the slip annulus. The shear traction τ(r) distribution curve in Figure 7.8 demon strates the yield annulus by the rounded transition between stick and slip zones, compared with the sharp transition for the elastic case, as shown in Figure 7.7. 7.2.5.4 Fretting Regimes. It was the elastic–plastic theory of fretting contact that confirmed the aforementioned features of plastic deformation in such contacts and allowed a more realistic explanation of the stick and slip zones.26 A connection
Figure 7.8â•… Elastic–plastic conditions in a fretting contact.23
7.2 Classification of Wear Mechanisms
â•… 87
of the contact zones with the amplitude of oscillation was determined by Vingsbo26,29,30 by defining various fretting regimes. This was an important step forward in the understanding of fretting, since it provided a background for presenting the contact conditions and the response of materials in a new way, that is, through fretting maps,17,26,29,30,31,32 which will be discussed later. Stick Regime (a′╯≈╯a).â•… The stick regime is characteristic of lower amplitudes of oscillation. The displacement is accommodated by elastic deformation of the asperi ties. The slip and yield annuli are negligible. No energy is dissipated in the process, as seen from the FT–δ curve (line) (Fig. 7.9a). The contact conditions are determined by the elastic deformation of the bulk material and asperities in the stick zone across the entire contact circle (a′╯≈╯a) (Fig. 7.10a). In the stick regime, limited surface damage is caused by oxidation and wear, and an absence of fatigue crack formation has been observed for some studied systems up to 106 cycles. Consequently, wear occurring in the stick regime is some times called low-damage fretting.26 Partial Slip Regime (0╯<╯a′╯<╯a).â•… With increasing amplitude of oscillation, plastic deformation of the contact surface becomes more important. There is still a central stick region, where the asperities elastically deform, and an outer annulus region develops, where microslip occurs. Thus, plastic deformation occurs in terms of plastic yield of asperities in the yield annulus and in terms of fracture of asperities in the slip zone. The FT–δ curve, shown in Figure 7.9b, changes shape from a line to a hysteresis loop. The area of the FT–δ curve corresponds to dissipated energy due to friction and plastic deformation. The contact area is characterized by two regions, as presented in Figure 7.10b. In the partial slip regime, wear and oxidation effects are small, and accelerated crack growth may result in strongly reduced fatigue life. The highest alternating stress occurs at the surface of the boundary between the stick and slip zones; hence, it is at this point where fatigue cracks would be expected to initiate. Hence, it is in the annular microslip region where damage with characteristic fatigue cracks occurs in this fretting regime.
FT
FT
(a)
FT
(b)
(c)
Figure 7.9â•… FT–δ plots for different fretting regimes: (a) stick, (b) partial slip, (c) gross slip.23
88â•…
CHAPTER 7â•… Wear Mechanisms
(a)
(b)
(c)
WEAR RATE (m3 (Nm)–1)
Figure 7.10â•… A schematic of the contact conditions in different fretting regimes: (a) stick, (b) partial slip (stick–slip), (c) gross slip.23
10–14
STICK PARTIAL SLIP
GROSS SLIP RECIPROCATING SLIDING
10–15
10–16
1
3
30 100 300 10 DISPLACEMENT (µm)
1000
Figure 7.11â•… Variation in fretting wear rate with displacement amplitude.23,26
Gross Slip Regime (a′╯=╯0).â•… When the amplitude reaches a critical value and the condition FT╯=╯μFN is satisfied, the coefficient of friction (static) drops from its maximum value to a lower value, which corresponds to the kinetic coefficient of friction. A hysteresis loop represented by the FT–δ curve in Figure 7.9c changes its shape in accordance with a coefficient of friction shift in every half-cycle. This occurs at the moment when the displacement overcomes the critical value and the gross slip conditions are satisfied. The entire contact area is in the macroslip region (Fig. 7.10c). In the gross slip regime, there is severe surface damage by wear, which is assisted by oxidation, and fretting fatigue is suppressed by the continuous elimina tion of the contact fatigue cracks by the wear process. This type of wear is usually termed fretting wear. As mentioned previously, the term fretting corrosion can be used when oxidation takes place. But due to more general considerations and the frequent simultaneous occurrence of different surface damage features, fretting wear is now generally preferred. Reciprocating Sliding Regime.â•… When the amplitude of oscillation is high enough, the wear mechanisms and wear rate become equal to those in reciprocating sliding.21,22,26. The point at which this occurs is referred to as the transition to the reciprocating sliding regime. Figure 7.11 shows wear rates for various fretting
7.2 Classification of Wear Mechanisms
â•… 89
regimes for a wide range of steels, represented by the dimensional wear coefficient. Mixed Fretting Regime.â•… Later, a new fretting regime was suggested31 to comple ment the already presented basic fretting regimes. It is named the mixed fretting regime, since it occurs in partial slip and gross slip regimes, and thus, both fretting wear and fretting fatigue damage are observed. It was reported that this type of fret ting regime could cause the highest fretting damage and is characterized by the elliptical shape of the hysteresis loop and by the fact that partial slip and gross slip conditions can be established alternately many times. 7.2.5.5 Determination of Fretting Regimes. The points of transition of fret ting regimes can be determined by observable changes in wear scar morphology or by calculated criteria, which use friction–displacement data from the tests or purely theoretically obtained values. The transition from stick to partial slip regime cor responds to an opening up of the fretting hysteresis loop. In the stick regime, the energy is accommodated by elastic deformation and no energy is dissipated, making the area of the loop equal to zero. Thus, a value for dissipated energy that is greater than zero corresponds to the transition from the stick regime to the partial slip regime. However, the determination of the transition from partial to gross slip is slightly more complicated, and two methods are now discussed. The point of incipient gross slip, in terms of displacement amplitude, can be found from either the friction force–displacement relations or from frictional energy dissipation. When the increasing displacement δ has reached the critical gross slip value δCR, there is a drop in tangential force FT to a critical value FT,CR that corre sponds to the transition from static to kinetic friction. Hence, the point that coincides with the position (δCR, FT,CR) is classified as the critical transition coordinate. However, it was found11 that the critical transition coordinate is easier to discern with a dissipated energy criterion. If the dissipated energy Ed is plotted as a function of displacement δ, taken every time from the same number of cycles, it generally shows a monotonically increasing pattern. However, it was revealed that, at a certain displacement, there is always a sudden increase in the slope dEd/dδ of the energy curve for roughly the same δ as the δCR of the force curve, described previously. Thus, the sudden increase in the slope of the energy curve indicates the point of incipient gross slip (Fig. 7.12). 7.2.5.6 Fretting Maps. One of the major reasons that fretting still remains an industrial problem in spite of numerous investigations is that the experimental results are difficult to compare because of the complexity of the process itself and the many influencing parameters.23,26,33,34,35,36 Because of so many contact and working condi tions, it is also difficult to predict the working life of a machine component that is subjected to fretting. Vingsbo26 showed that it is possible from the experimental data, based on dynamic measurements of tangential force and displacement, to distinguish and recognize different fretting regimes, which are characterized by different surface damage as discussed previously. This is a promising finding for the prediction of possible surface damage and the working life of various machine components.
90â•…
CHAPTER 7â•… Wear Mechanisms
Dissipated energy (µJ)
140 120 100 80 60
δcr
40 20 0 0
20
40
60
80
100
Displacement amplitude (µm)
Figure 7.12â•… Dissipated energy versus displacement diagram showing the transition from partial to gross slip regime.23
Normal force (N)
30 Stick
Gross slip
20
Reciprocating sliding
10
Mixed stick and slip
0 1
3
10
30
300
1000
Displacement amplitude (µm)
Figure 7.13â•… Displacement amplitude versus normal force fretting map.23,26
A fretting map is an illustration that portrays the pertinent regimes in two variables, where the regime boundaries represent the critical values for the transition from one regime to another.26 According to many influencing parameters, various fretting maps can be plotted, such as displacement versus normal force, displacement versus frequency, and displacement versus wear. Figure 7.13 shows one of the most informative and widely used maps, describing fretting regimes as a function of normal load and displacement amplitude. The use of two types of fretting maps is suggested, running condition fretting maps (RCFMs) and material response fretting maps (MRFMs), as shown in Figure 7.14. The RCFMs plot the normal load versus displacement amplitude for a given frequency. The three zones identified in the RCFM are stick, partial slip, and gross slip. The MRFM, on the other hand, plots stress or equivalent stress versus ampli tude. The three zones identified in the MRFM are no degradation, cracks, and particle-detachment wear zone. Fretting maps are therefore used for interpretation of experimental data and determining the boundaries for tribological and working parameters in order to get an appropriate fretting regime or even to avoid it.
7.2 Classification of Wear Mechanisms
â•… 91
(a)
(b)
Figure 7.14â•… (a) Running condition fretting map (RCFM); (b) material response fretting map (MRFM). L, normal load; a, displacement amplitude; S, equivalent stress.23,24
Figure 7.15â•… Velocity accommodation sites and modes.23,24
7.2.5.7 Velocity Accommodation in Fretting. The concept of velocity accommodation means that the relative displacement and velocity difference between the contacting bodies are accommodated at different sites and according to different modes.24,33 The four basic accommodation modes (M1–M4) such as elastic deforma tion, fracturing, shearing, and rolling and the five accommodation sites (S1–S5) in the contact are schematically indicated in Figure 7.15. Combination of modes and sites results in 20 individual velocity accommoda tion mechanisms (VAMs). The changes in the operative VAMs as well as the material properties are reflected in the evolution of the macroscopically measured friction behavior. Different mechanisms may be operative simultaneously at different locations in the contact zone. This generally happens, for instance, in contacts subjected to partial slip conditions. The macroscopically measured tangential force thereby
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reflects the integral action of all VAMs operative in the fretting contact as well as the material properties involved. 7.2.5.8 Friction Logs. When fretting maps were confirmed to be applicable as tools in recognizing global fretting conditions in terms of running conditions or material response on a macroscale, a new approach was developed, that is, friction logs,24,33 which helped to gain an understanding of the contact conditions on a microscale. Friction logs, sometimes called fretting logs, are three-dimensional graphs that give the variation in the coefficient of friction versus displacement stroke length and time or number of cycles, as shown in Figure 7.16. Friction logs reveal the actual development of the coefficient of friction during a fretting experiment. This allows the coefficient of friction to be correlated to other influencing factors, different contact conditions as well as underlying mechanisms.
7.2.6â•…Solid Particle Erosion Solid particle erosion is a wear mechanism that is highly relevant for advanced materials such as ceramics and composites.37–51 Since many practical applications suffering from erosion are associated with the use of these materials, we put some emphasis on erosion in this chapter. When solid particles entrained in a fluid stream strike a surface, the damage is referred to as erosive wear (see Fig. 7.17). The fact that erosion is mainly studied in reference to specific problems in industries, as well as the unavailability of a complete experimental database, limits the understanding of the basic principles of erosion wear.5 Namely, the erosion problems can be found in a number of engineer ing instances, such as equipment used in catalytic cracking of oil, grit blasting nozzles, coal turbines, hydraulic turbines, coal hydrogenation equipment, and pipe lines used for coal transportation. On the other hand, erosion can be used in useful applications in such processes as sand blasting, abrasive deburring, and the erosive drilling of hard materials.37 The wear of metal plates by the entrained soil in air and injuries due to flying pebbles during rowing were some of the reported instances of erosion. The erosion of cutting tools or workpiece surfaces by chipped-off debris during machining is one of the serious problems in modern manufacturing processes. Even though there were instances of quoting erosive wear since 1873,38–40 the first systematic investigation was reported by Wahl and Hartstein in 1946.41 The erosion experiments of Kashcheev on Cu–Al alloys at different angles showed that the erosion depends on the angle of impact.42 Note that erosion differs from threebody abrasion, which also involves loose particles, in the origin of the forces between the particles at the wearing surface. In erosion, several forces act on the eroding particle at the contact.43 The dominant forces acting on eroding particles include gravitational force due to the weight of the particle, drag force due to flowing fluid, contact forces exerted by the surface, and interparticle contact forces in a stream of particles. In contrast, abrasion involves the movement of hard particles along the surface under contact pressure and hence the amount of material removal depends on the normal load and the distance traveled, as well as the size and shape of particles, as discussed previously.
7.2 Classification of Wear Mechanisms
â•… 93
+750 N F
0
–750 N –50 µm
10 102 103 104 105
N
(a) 50 N F
0
–50 N –15 µm
0 +15 µm
10 102 103 104 105 N (b)
Figure 7.16â•… Representative friction log when (a) debris is formed and when (b) cracks are present.23,24
94â•…
CHAPTER 7â•… Wear Mechanisms
Figure 7.17â•… Schematic of solid particle erosion.5
Therefore, in erosion, the number and mass of particles and their impact veloc ity at the striking contact are considered to be the primary parameters in determining the extent of wear. Moreover, Finnie,44 after his pioneering work on erosion of ductile and brittle materials, reported that the factors influencing solid particle erosion include: (1) fluid flow conditions such as angle of impingement, particle velocity, particle rotation, particle concentration in the fluid, and nature of the fluid and its temperature; (2) particle properties such as size, shape, hardness, and strength (resistance to fragmentation); and (3) surface properties such as stress–strain, strain rate and temperature, hardness, fracture toughness, stress level and residual stresses, and microstructure. It is convenient to divide the erosion process into two major parts: determina tion of fluid flow conditions and quantification of material removal. The first part involves the number, mass, and direction of the striking particles at the surface; the second part involves understanding the wear mechanisms. This book, focusing on materials behavior, is thus limited to discussing only the second part. It can be understood that the dominant mechanism of one material cannot be expected to be the same for different materials, and often several mechanisms were reported to be acting simultaneously. Therefore, for a better understanding of the importance and response of different materials, we divide the range of materials into two types: ductile and brittle. Ductile materials wear mainly by plastic deformation; brittle materials wear by flow or fracture, depending on the impact conditions. 7.2.6.1 Erosion of Ductile Materials. The erosion or erosive behavior of ductile metals was reported in detail by several researchers.45–47 During erosion of ductile materials, a large number of abrasive particles strike the surface. As it may not be possible that every particle will cut the striking surface, we will start with the consideration that the motion of a single rigid abrasive particle will result in displace ment or cutting away of part of the surface. Erosion wear at normal incidence can be possible with certain assumptions.3 The particle does not deform and the deformation of surface is perfectly plastic with a constant indentation pressure (hardness). At time t after initial contact, the particle of mass dm with an initial velocity v indents the surface to a depth x such that the
7.2 Classification of Wear Mechanisms
V
â•… 95
Mass dm Area A(x) x d Time = 0
t
t0
Figure 7.18â•… Schematic of erosion by a hard single particle striking soft surface.5
cross-sectional area of the indent impression is A(x), which is dependent on the shape of the particle (see Fig. 7.18). If the particle comes to rest at a depth in the subsurface region, the work done by the retarding force is equal to the initial kinetic energy of the particles: (7.21) dV = dmv 2 / 2 H , where dV is the volume of the material displaced from the indentation and H is the hardness of the material surface being eroded. If the total mass of the impinging particles is M and ke is the proportion of the resulting displaced material, then total wear volume is
V = ke Mv 2 / 2 H .
(7.22)
The steady-state erosion ratio is usually written as the ratio of mass of material removed to mass of erosive particles striking the surface. Therefore,
Er = ke ρv 2 / 2 H ,
(7.23)
where ρ is the density of the material being eroded. Note that the abrasion wear equation is partly similar to that of erosion (Eq. 7.23), as both are inversely proportional to hardness, but different as the normal load in abrasion is replaced by mv2 in erosion. Thus, it is not significantly different from abrasion, except that the contact stresses arise from the kinetic energy of the particles flowing in air or a liquid stream as it encounters the surface. The kinetic energy of particles can be measured with the particle velocity and impact angle combined with the size of the particle. However, the equation is crude in its form, as the effect of impact angle, size, and shape of the particles are included by a single ke value. The typical ke value for metals ranges from 10−5 to 10−1. The erosion reaches to maximum at 20–30° and then decreases to reach almost one-third of the peak erosion value at normal incidence.48 Several researchers reported cutting and deformation (plowing) as two basic mechanisms when eroding ductile materials. During cutting erosion by angular particles, crater formation is caused by single or multiple impacts of the micromachining or lip formation. The rounded particles deform the surface by plowing or surface fragmentation by several indentation-type impacts. During single indentation, the material is displaced as
96â•…
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small crater lips; subsequent indentations flatten and strain further, creating thin platelets of highly stressed metal that are finally removed from the surface by suc ceeding particles. At low grazing angles, cutting erosion dominates; at normal incidence, the deformation mechanism dominates. These phenomena are clearly similar to those in abrasion. The following more realistic model was proposed by Hutchings,42 according to which the point of action of the forces on the particles is allowed to move during the impact, leading to a complex expression for Er:
Er = g1ρv n / Hf (α ) .
(7.24)
The value of g1 depends on the geometry of the particles and on the fraction of the particles actually cutting in an idealized manner. The value of exponent n is between 2 and 2.5; f(α) is a function of the angle of impingement, α. The erosion by spherical particles at normal angles can be modified with ductility term as (7.25) Er = ( g2ρσ1 / 2 v 3 ) / ε c2 H 3 / 2, where σ is the density of spherical particles and εc is the critical plastic strain at which detachment of wear debris occurs. The experimental evidence for erosion of pure metals reveals that the erosion is related to particle velocity as
Er ∝ v n ,
(7.26)
where n ranges between 2.3 and 3. With an increase in velocity, the particle dips deeper and the resultant force on the particle changes position and the rotation decreases. This leads to an additional volume removal term, which is a function of the cube power of velocity. The higher values of n are associated with the steeper angles of impact. Surface hardness is important since the surface materials become heavily work-hardened by particle impacts. Further, it was reported that the degree to which plastic flow is localized around each particle impact site is also important, since this influences the susceptibility of displaced material to removal.43 This can be observed in several metals and alloys. 7.2.6.2 Erosion of Brittle Materials. The erosive wear of brittle materials occurs by the propagation and intersection of cracks produced by impacting par ticles onto the surface of brittle materials. It was observed in a number of inves tigations that severe cracking occurs when a particle impacts the surface of a brittle material at normal incidence. As shown in Figure 7.19, erosion is rapid at angle of 90°, indicating brittle fracture of a classical brittle solid, that is, alumina. Once fracture occurs, the extent of material removal depends on the propagation and intersection of surface and subsurface cracks. Therefore, it is important to under stand erosion behavior with initial cracking conditions. The literature indicates two types of theories of material removal for brittle fracture, depending on the shape of the indenter.51 Brittle fracture due to concentrated (counter formal) contact can be understood in terms of Hertzian elastic stress distributions. Static as well as dynamic (such as
7.2 Classification of Wear Mechanisms
â•… 97
20 V a Erosion-grams per gram of abrasive × 10–4 (Al) × 10–3 (Al2O3)
16
Al
12 Al2O3 8
4
0
0
30 60 Angle of impingement α, in degrees
90
Figure 7.19â•… Erosion loss of aluminum and aluminum oxide when impinged by silicon carbide particles.30
in sliding) indentations with a blunt (e.g., spherical) indenter at sufficiently high loads produce surface ring cracks, which intersect with the subsurface conical cracks, resulting in material removal. The relevance of this model was reported when a glass surface was eroded by spherical steel shots and ring cracks were formed in the initial stages. With increasing particle impacts the cone-shaped fracture surfaces associated with ring cracks begin to intersect and material is removed. Based on this model, the erosion rate of brittle materials can be explained by the expression (7.27) Er ∝ v m r n , where the exponent values should be related to the statistical distribution of fracture stress raisers on the tribological surface. Thus, according to such models, the velocity (v) and particle size (r) exponents (m and n) would be typically around 0.1 and 3.0, respectively.5 For sharp particles, erosion by elastic–plastic indentation fracture theory of brittle materials is applicable, according to which the intersection of lateral cracks with each other and with the surface eventually result in material removal.51,52 The sequence of crack formation and material removal in brittle materials due to a sharp indenter can be explained by the schematic illustration in Figure 7.20. This phenom enon also occurs in the case of abrasive wear. The high contact stresses are relieved by plastic flow around the tip of the indenter. When the contact stresses reach a critical value, tensile stresses across the vertical midplane initiate a median vent crack, which extends further with load.
98â•…
CHAPTER 7â•… Wear Mechanisms
W Sh Sharp asperity it
Motion of indenter
Plastic groove Surface
Potential wear zone
c
d Lateral crack
Plastic zone
Median crack
Figure 7.20â•… Schematic illustrating the mechanism of wear by a sharp asperity sliding on the flat surface of a brittle material causing lateral fracture.5
During unloading, the median crack closes and the relaxation of the deformed mate rial around the contact region produces residual stresses, resulting in lateral cracks. The intersection of lateral cracks with the surface leads to material removal as chips or lips. The formation of rims or lips inside the crater formed due to erosion can be explained on the basis of this theory. In practical situations, the particles hit the surface at lower angles and slide over it for a finite distance, resulting in wider scratches that are larger in length than width. Thus, the brittle–ductile transition is likely to initiate lateral cracking. The impact of a sphere at an angle α on semi-infinite body would result in the following relation of particle velocity (v) and impact force with particle size (D):
D = dc Fcr1 / 2 ( v sin α )
−3 / 5
,
(7.28)
where dc is a material-dependent constant and Fcr is the threshold force required or lateral cracking obtained in scribing. Another approximation can be derived by equating kinetic energy to the work done by deformation as
D = dc′Fcr1 / 2 ( v sin α )
−2 / 3
,
(7.29)
where d′c is another material-dependent constant. Scribing with an edge-leading indenter tip leads to a lower threshold than with a face-leading tip.
7.3 CLOSING REMARKS As a closing note, it needs to be emphasized here that, under a given combination of operating parameters, a number of competing wear mechanisms can result in material removal from the contacting surfaces in a tribosystem, and it is therefore important to identify all the dominant wear mechanisms. It is equally important to investigate the role of various material parameters, for example, hardness, elastic
REFERENCES
â•… 99
modulus, and fracture toughness, on the severity of individual wear mechanisms. Such an approach of investigation into various wear mechanisms is essential to the development and design of new wear-resistant materials.
REFERENCES ╇ 1â•… K.-H. Zum Gahr. Microstructure and Wear of Materials. Elsevier, Oxford, 1987. ╇ 2â•… A. Tewari, R. K. Bordia, and B. Basu. Model for fretting wear of brittle ceramics. Acta Mater. 57 (2009), 2080–2087. ╇ 3â•… J. K. Lancaster. The influence of substrate hardness on the formation and endurance of molybdenum disulphide films. Wear 10 (1967), 103–107. ╇ 4â•… J. F. Archard. Contact and rubbing of flat surfaces. J. Appl. Physics 24 (1953), 981–988. ╇ 5â•… B. Bhushan. Principles and Applications of Tribology. John Wiley & Sons, New York, 1999. ╇ 6â•… T. Kayaba and K. Kato. Adhesive transfer of the slip-tongue and the wedge. ASLE Trans. 24(2) (1981), 164–174. ╇ 7â•… T. Sasada. Tribology in the 80s, Vol. I. NASA Conference Publication 2300, NASA Lewis Research Centre, Cleveland, OH, 1984, 197–218. ╇ 8â•… D. A. Rigney, H. Chen, M. G. S. Naylor, and A. R. Rosenfield. Wear processes in sliding systems. Wear 100 (1984), 195–219. ╇ 9â•… E. Rabinowicz. Friction and Wear of Materials, 2nd ed. Wiley, New York, 1995. 10â•… B. R. Lawn and D. B. Marshall. Hardness, toughness, and brittleness: An indentation analysis. J. Am. Ceram. Soc. 62 (1979), 347–350. 11â•… A. G. Evans and D. B. Marshall, Wear mechanisms in ceramics, in Fundamentals of Friction and Wear of Materials, D. A. Rigney (Ed.). American Society for Metals, Metals Park, OH, 1981, 439–452. 12â•… S. G. Roberts. Depths of cracks produced by abrasion of brittle materials. Scr. Mater. 40(1) (1999), 101–108. 13â•… S. Suresh. Fatigue of Materials. Cambridge University Press, Cambridge, 1988. 14â•… N. P. Suh. The delamination theory of wear. Wear 24 (1973), 111–124. 15â•… J. F. Tavernelli and L. F. Coffin. A compilation and interpretation of cyclic/strain fatigue tests. Trans. ASM 51 (1959), 438–450. 16â•… H. Hong, R. F. Hochman, and T. F. J. Quinn. A new approach to the oxidational theory of mild wear. STLE Trans. 31(1) (1988), 71. 17â•… M. Odfalk and O. Vingsbo. An elastic-plastic model for fretting contact. Wear 157 (1992), 435–444. 18â•… E. M. Eden, W. N. Rose, and F. L. Cunningham. Endurance of metals. Proc. Inst. Mech. Eng. 4 (1911), 839–974. 19â•… G. A. Tomlinson. The rusting of steel surfaces in contact. Proc. R. Soc. Lond. A 115 (1927), 472–4834. 20â•… D. Godfrey and J. M. Bailey. Early stages of fretting of copper, iron and steel. Lubrication Eng. 10 (1954), 155. 21â•… D. Godfrey. Investigation of fretting corrosion by microscopic observation. NACA Report 1009, NASA Center for AeroSpace Information, Hanover, MD, 1951. 22â•… R. B. Waterhouse. Fretting Corrosion. Pergamon Press, Oxford, England, 1972, 3. 23â•… M. Kalin. Fretting wear mechanisms in contact of steel and silicon nitride ceramics. PhD thesis, University of Ljubljana, 1999. 24â•… L. Vincent, Y. Berthier, and M. Godet. Testing methods in fretting fatigue: A critical appraisal, in Standardization of Fretting Fatigue Test Methods and Equipment, ASTM STP 1159, M. H. Attia and R. B. Waterhouse, Eds. American Society for Testing and Materials, 1992, 33–48. 25â•… A. D. Sarkar. Friction and Wear. Academic Press, London, 1980. 26â•… O. Vingsbo and S. Soderberg. On fretting maps. Wear 126 (1988), 131–147. 27â•… C. D. Holland. Fretting: A survey of present-day knowledge. J. Jr. Inst. Eng. 76 (1965), 68–98. 28â•… H. Mohrbacher. The tribological performance of advanced hard coatings under fretting condition. PhD thesis, Katholeike Universiteit Leuven, 1995.
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29â•… O. Vingsbo. Fretting and contact fatigue studied with the aid of fretting maps, in Standardization of Fretting Fatigue Test Methods and Equipment, ASTM STP 1159, M. H. Attia and R. B. Waterhouse, Eds. American Society for Testing and Materials, 1992, 49–59. 30â•… O. Vingsbo and J. Schon. Gross slip criteria in fretting. Wear 162–164 (1993), 347–356. 31â•… Z. R. Zhou and L. Vincent. Mixed fretting regime. Wear 181–183 (1995), 531–536. 32â•… S. Fouvry, P. Kapsa, and L. Vincent. Analysis of sliding behaviour for fretting loadings: Determi nation of transition criteria. Wear 185 (1995), 35–46. 33â•… Y. Berthier, L. Vincent, and M. Godet, Velocity accomodation in fretting. Wear 125 (1988), 25–38. 34â•… J. M. Dobromirski. Variable of fretting process: Are there 50 of them?, in Standardization of Fretting Fatigue Test Methods and Equipment, ASTM STP 1159, M. H. Attia and R. B. Waterhouse, Eds. American Society for Testing and Materials, 1992, 60–66. 35â•… Z. R. Zhou and L. Vincent. Effect of external loading on wear maps of aluminium alloys. Wear 162–164 (1993), 619–623. 36â•… S. Fouvry, P. Kapsa, and L. Vincent. Quantification of fretting damage. Wear 200 (1996), 186–205. 37â•… I. Finnie. Erosion of surfaces by solid particles. Wear 3 (1960), 87–103. 38â•… O. Reynolds. On the action of a blast of sand in cutting hard materials. Philos. Mag. 46 (1873), 337–343. 39â•… L. Rayleigh. The sand blast. Nature 93 (1912), 188. 40â•… T. Young. A Course of Lectures on Natural Philosophy and the Mechanical Arts. J. Johnson, (1807). 41â•… H. Wahl and F. Hartstein. Strahlverschleiss Franckhsche Verhandlung, Stuttgart, 1946. Translated into English, January 1979 for Lawrence Livermore National Lab., UCRL Translation 11447. 42â•… V. N. Kashcheev. Destruction of a metal surface as a function of the angle of impact of abrasive particles (in Russian). Zhur. Tekh. Fiz. 25 (1955), 2365. 43â•… I. M. Hutchings. Wear by particulates. Chem. Eng. Sci. 42(4) (1987), 869–878. 44â•… I. Finnie. Some reflections on the past and future of erosion. Wear 186–187 (1995), 1–10. 45â•… W. F. Adler (Ed.). Erosion: Prevention and Useful Applications, ASTM STP 664, 3rd ed. Effects Technology, Inc., Santa Barbara, CA, 1979. 46â•… P. Shewman, G. Sundararajan, R. A. Huggins, et al. The Erosion of Metals Annual Review of Material Science. Annual Reviews, Palo Alto, CA, 1983, 301–318. 47â•… I. Finnie. The mechanism of erosion of ductile metals, in Proceedings of the 3rd US National Congress of Applied Mechanics, American Society of Mechanical Engineers, New York, 1958: 527–532. 48â•… J. E. Ritter (Ed.). Erosion of Ceramic Materials. Trans. Tech. Publications, Zurich, Switzerland, 1992. 49â•… I. M. Hutchings. Mechanisms of the erosion of metals by solid particles, in Erosion: Prevention and Useful Applications, ASTM STP 664, 3rd ed., Adler, W. F. (Ed.). American Society for Testing in Materials, Philadelphia, 1979, 59–76. 50â•… I. Finnie, J. Wolak, and Y. Kabil. Erosion of metals by solid particles. ASTM J. Mater. 2 (1967), 682–700. 51â•… B. R. Lawn and M. V. Swain. Microfracture beneath point indentations in brittle solids. J. Mater. Sci. 10(1) (1975), 113–122. 52â•… A. G. Evans, M. E. Gulden, and M. Rosenblatt. Impact damage in brittle materials in the elasticplastic response regime. Proc. R. Soc. Lond. A 361 (1978), 343.
CHAPTER
8
LUBRICATION Under nonlubricated conditions, the mating materials typically experience high friction and wear, and the prime role of lubricants is to reduce friction and wear. Lubricants can achieve these goals through various mechanisms that will be briefly presented here in various lubrication regimes, which depend on the operating conditions such as load, contact pressure, velocity, and temperature, as well as materials and lubricants used. Lubricants are used to prevent direct contact between moving surfaces. This can be done through oil film separation or by providing molecularlevel boundary films at the surfaces, which act as low-shear and/or wear-resistant protective layers. Conventional lubricants are compounds consisting of a base lubricant oil and additives. However, different compounds in the lubricant have different roles, and these are not limited to tribological behavior only. The purpose of the base oil is also to remove heat from the contact and distribute it around the system, to carry the additives, to remove the wear debris from the contacts, to cool the system, and to reduce noise and vibrations. However, base oils also have the ability to reduce wear and friction and/or to separate the contacting surfaces via hydrodynamic (HD) oil film. Additives, on the other hand, take the key role in providing the proper tribological performance, when surface asperities of the moving bodies come into contact; they also keep the oil clean without significant degradation during its expected life. Depending on the type of application, the required amount of additives can vary from about 5% to more than 20% in the oil.
8.1 LUBRICATION REGIMES A regime of lubrication in which a thick layer of lubricant is maintained between two interacting surfaces with little or no relative motion is called hydrostatic lubrication. Hydrostatic bearings support load on a thick film of fluid. As a bearing with convergent shape in the direction of motion starts to move in the longitudinal direction from rest, a thin layer of fluid is pulled through because of viscous entrainment. The constrained fluid layer is subsequently compressed between the bearing surfaces, creating a sufficient HD pressure to support the load without any external pumping agency. This kind of lubrication with formation of a thick film at the interface is known as HD lubrication. Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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The hydrodynamic (HD) was first explained by Osborne Reynolds1 in 1886, who presented the model that has proved that continuous oil film separation can be established between two moving bodies and carry the load in such a contact, if some conditions are satisfied. The Reynolds equation can be obtained through solution of the Navier–Stokes equations or by solving the equilibrium of forces acting on the fluid. The equation is very complex in its full form, but there exist many simplified Reynolds equations. For demonstrative purposes, Equation 8.1 is just one of the simplified one-dimensional equations that can be obtained from countless literature sources: ∂ 3 ∂p ∂ 3 ∂p dh h = 6Uη , h + ∂x ∂x ∂y ∂y dx
(8.1)
where p is the lubricant film pressure, η is the dynamic viscosity at pressure p, U is the relative fluid velocity, and h is the fluid film thickness. In his theory, Reynolds showed that a certain relative velocity must exist between the moving bodies, which would eventually cause a wedge between them and satisfy the pressure distribution in the oil film that will carry the load. When this occurs, the surfaces can be fully separated by the oil film, which means that the “no-wear” conditions are achieved, and the friction can be extremely low, on the order of 0.001 or nonmeasurable. This regime is called the HD regime and is typical for a system having conformal contacts, that is, contacts for which the shape and size of the surfaces are very similar. For example, this can be valid for flat-on-flat (Fig. 8.1a) or shaft-in-a-bore (Fig. 8.1b) contacts. Another key reason for the HD lubrication in such contacts is the low pressure that is generated between the surfaces due to very high contact area. Thus, relatively high loads can be applied to these contacts, but the contact pressure will remain low, on order of megapascals, and a relatively thick lubricating film can be formed. The HD regime is very important for journal bearings used in pumps, turbines, large engines, and bearings, where the operation is steady and long term, with little variation, such as power stations. Moreover, the HD lubrication phenomenon controls the buildup of water films under the tires of automobiles and airplanes on wet roadways or landing strips.2,3
(a)
(b)
Figure 8.1â•… Conformal contacts where hydrodynamic regime can be achieved. (a) Flat on flat and (b) shaft in a bore contact.
8.1 Lubrication Regimes
(a)
(b)
â•… 103
(c)
Figure 8.2â•… Schematic of some typical nonconformal contacts in machine elements: (a) gears, (b) cams, and (c) ball bearings.
However, if the surfaces are nonconformal, such as point or line contacts, the contact stresses are high, in the range of hundreds of megapascals or even gigapascals. This is a typical condition for rolling bearings, gears, cams, and so on, as shown in Figure 8.2. If the Reynolds theory is applied to nonconformal concentrated contacts, the results would show that there is almost “no oil film” between the surfaces or is far too thin to separate the surfaces. However, even components under such conditions can run at almost “no-wear” conditions for a long time, such as bearings or gears. This is a clear indication that conventional HD regime cannot be used for nonconformal contacts. In 1949, Ertel and Grubin4 explained and resolved the mystery about this problem. They found that other parameters must be simultaneously considered. First, elastic deformation of the heavily loaded nonconformal contacts is not negligible and should be considered in the model. Second, oil viscosity is by far more dependent on the pressure than, for example, the temperature, as we usually think. Namely, in an average engineering range of oil temperatures, the viscosity changes typically by about 10 times. However, in the range of pressures that oil can experience in concentrated contacts, the oil viscosity could change over several orders of magnitude, and thus oil becomes quasi-solid at very high pressures, as evident also from the Barus exponential model5:
η = η0 eαp,
(8.2)
where p is the lubricant film pressure, α is the pressure–viscosity coefficient, η,η0 are the dynamic viscosity at pressure p and at atmospheric pressure, respectively. With this understanding, the calculated values for film thickness became feasible, in the range around 1â•›µm; the lubrication regime that is characterized by these high contact pressures with relevant elastic deformation and oil viscosity change is known as the elastohydrodynamic (EHD) regime or elastohydrodynamic lubrication (EHL). Therefore, in EHD lubrication, a viscous oil film is formed that separates the contacting surfaces, similarly as in the HD regime, but with the distinct difference that the film thickness is much lower. This further means that maintaining the perfect
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conditions where “1 micron” separates the heavily loaded surfaces in machine elements is much more difficult and, of course, imperfections are more frequent. However, even in EHL, the friction is very low (much less than 0.1) and also the wear could be very low or negligible, if the running conditions are stable and not too severe. Although calculation of the film thickness under EHL conditions is complex and cannot be solved analytically, there exist equations that are very useful and rather accurate. The most frequently used equation for general elliptical contacts is the Hamrock–Dowson equation6
hmin = 3.63 RxU 0.68 G 0.49W −0.073 (1 − e −0.68 k ) ,
(8.3)
where hmin is the minimum oil film thickness, U is a nondimensional parameter for speed, W is a nondimensional parameter for load, G is a nondimensional parameter for materials, k is a nondimensional parameter for contact geometry, Rx effective radius in the x direction. Some more elaborated versions of this equation and many other details can also be found in the literature.7,8 From Equation 8.3, it is clear that for a given set of mating surfaces, that is, contact geometry, the minimum oil film thickness that ensures lubrication is more sensitively dependent on the nondimensional parameter related to speed than that of material properties. Also, hmin has a weak inverse dependence on load. Despite this rather elegant solution for the EHL problem and having the film thickness formula for concentrated contacts, there still exists one problem. Namely, the surfaces are not ideally (atomistically) smooth. The average surface roughness for most engineering surfaces is in the range from nanometers to micrometers. Thus, for a film thickness of only a micrometer or less, roughness can be even larger than this value; even if it is smaller, it is certainly still very important. Namely, due to machine vibrations, temperature and consequently viscosity change, surface damage, wear debris, and so on, the oil film thickness varies and cannot be kept ideal, as calculated. Thus, the surface asperities can frequently touch each other, and the probability will depend not only on film thickness, but also on roughness. The larger the difference between film thickness and roughness, the more seldom the asperity collisions and the safer the mechanical operation will be. A way to determine how safe the operation is and how large the consequences of these collisions will be was proposed by Tallian,9 who introduced the λ parameter, which is in essence a ratio between the oil film thickness and the roughness of the two moving surfaces:
λ=
hmin R + Rq22 2 q1
,
(8.4)
where Rq is the root mean square (RMS) roughness value and the subscripts 1 and 2 denote two mating solids.
8.1 Lubrication Regimes
â•… 105
As defined in Equation 8.4, the λ parameter indicates reasonably well the ability of the oil film to separate the surfaces and thus the expected wear and friction in certain contacts, and consequently the effect on the lifetime of components. Namely, it was found that the service life of machine components strongly depends on the λ ratio. In accordance with this, new lubricating regimes were defined. There exist different relations between λ and surface damage, but the regime transitions are broadly accepted as follows: λâ•›≥â•›3 3â•›≥â•›λâ•›≥â•›1 λâ•›≤â•›1
HD/EHD lubrication—no reduction in life, no wear; mixed lubrication (ML)—no or limited reduction in life, small wear; boundary lubrication (BL) regime—reduction in life, (severe) wear.
The mixed lubrication (ML) regime is thus the regime where the asperities of the surfaces often come into contact, because the oil film is not high enough to separate them. Nevertheless, the HD effect is still very pronounced and most of the load is carried by the oil film, especially at λ╯=╯3.0. The severity of the contact conditions will change as a function of λ. The lower the λ value, the more severe the conditions, friction will be higher and so will the wear; thus, the life of the component will decrease. The efficiency of the tribological systems in the ML regime will depend also on the shape of asperities, not only on their height and number. For example, if the asperities are rather flat, they can partially act as small EHD spots and the conditions will be better than expected from the λ value (Fig. 8.3). However, if they are too flat, then the HD effect will be lost due to oil entrapment failure. On the contrary, if the asperities are very sharp, then the damage on the surface will be high and the conditions consequently will be poor. EHD lubrication is most readily induced in
Figure 8.3â•… Mixed lubrication regime with possible micro-elastohydrodynamic (EHD) contact spots.
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heavily loaded contacts, where load acts over a relatively small area of contact, such as the point contacts of ball bearings and of gear teeth. As is evident, ML is a rather complex regime with many parameters acting, which are, however, very difficult to control or model. Tribological performance in this regime depends strongly on both the fluid film and the asperity contacts, with all their specifics and interplay. Moreover, because several asperities come into contact, they will partially deform and some wear will occur; thus, the contact condition will change with run-in and the regime may even change. The lubrication regime for λ values below 1 is called the boundary lubrication (BL) regime. In the HD effects in carrying the load are negligible, since the oil film is so thin that the number of asperities in contact is high enough to carry the load. Accordingly, in terms of contact mechanics, this regime is very similar to nonlubricated conditions. However, the tribological properties of boundary lubricated contacts are significantly better. What then is the function of the lubricant in this regime? Asperity contacts under BL conditions depend on surface protective layers that form at the surfaces, typically with the help of additives. Without these layers, the conditions would indeed be similar to those in nonlubricated contacts, thus too severe for the operating conditions and the contacts would completely fail. Therefore, lubricants must provide a source to form the surface layers, which will be wear resistant and provide low friction, that is, low-shear and high-strength boundary films. The effectiveness of these films under specific conditions defines the success of a particular oil and, consequently, the tribological system. The boundary layers could be physically adsorbed, chemically adsorbed, or chemically reacted with the surfaces. Of course, the strength of protection is very different for these mechanisms. Moreover, also the chemistry of compounds used and the operational conditions to form any of these types of layers is very different. Since the boundary surface layers are formed in situ in the tribological system, when needed, this clearly indicates a vast variety of options of additives, surfaces, and conditions possible to form a specific protective boundary film. In addition, the additives have sometimes synergistic but sometimes antagonistic effects. In this respect, BL is much more of an art than a science. Many additives are used—and have been for decades—without fully understanding their detailed protective mechanisms. Luckily, recent development of several extremely surface-sensitive analytical techniques, and nanotribology, changes this situation, and BL is becoming increasingly understandable. There exists a significant literature pool on specific additives, systems, mechanisms, and so on, for BL, but hardly any papers or books can be found that would summarize most of them. The reason is not only in the number of possible systems studied, but also in the differences of opinion among scientists and different findings that are difficult to indisputably prove. Nevertheless, the three mechanisms addressed here are typical routes of interactions between the surfaces and additives and/or oils. Physically adsorbed boundary layers are the most vulnerable, because they are based on physical interactions between the surfaces and lubricants. They can only be effective in mild conditions with low temperatures, low pressures, and low shear forces. They are fully reversible, since no permanent changes occur. Most base oils can also provide these layers, sometimes even with the help of impurities in these
â•… 107
8.2 Stribeck Curve
oils. On the other hand, additives that can help to form these layers are called friction modifiers. Chemically adsorbed boundary layers are stronger than physisorbed layers because of the stronger interaction and exchange of molecules of lubricants and surfaces, which are however, still rather limited. With this interaction mechanism, permanent changes occur in the structure of oil molecules and surfaces, and the process is thus irreversible. Because of stronger interaction, the effectiveness of these layers is better, and they can last longer under more severe conditions, including ceramic surfaces.10 Surface active molecules such as fatty acids and aliphatic alcohols are required in the oil and additives. Both friction modifiers and antiwear additives are used to enable this mechanism. The most effective additives are based on chemical reactions between the surface and additives. With this mechanism, a new protective compound is formed at the surface as a consequence of reaction between the surface active substance (additive) and the surface. Chemically reactive boundary lubricant agents containing sulfur, chlorine, and phosphorus atoms in the molecule react with solid (metal) surfaces to form lubricating films (sulfide, chloride, and phosphide) of low shear strength. These layers are usually very effective in wear protection. Sometimes they need some time to form on the surface in satisfactory amounts (induction period), especially in the case of antiwear additives, but sometimes the low-shear layers are formed in the contact instantaneously and may even be removed during the same contact. The latter are called sacrificial layers, typical for extreme-pressure additives, which are active only under the most severe conditions at extreme pressures and extreme temperatures. Due to reaction with the surface and removal of the layer, this also causes a corrosive effect, so these additives should be used with care. However, it should also be noted that for contacts under severe conditions, as in many current applications involving exposure of surfaces at high loads and temperatures, these contacts and the entire system would not be able to run without highly efficient and high-tech additives that form these protective boundary layers. Nevertheless, they are mostly suited for metal surfaces and it is generally very difficult to obtain such chemical reaction films with extreme pressure agents for chemically stable ceramics.11
8.2 STRIBECK CURVE The curve that can experimentally show the regimes and transitions between them is well known as the Stribeck curve, which is schematically presented in Figure 8.4. The Stribeck curve plots the coefficient of friction as a function of (Stribeck) parameter, which is proportional to viscosity and velocity and inversely proportional to load. It is clear that the lower the load, or the higher the velocity, the position on the curve will move toward the HD or EHD regime. In contrast, at low velocity or high load, the conditions will approach BL. It is also clear that higher viscosity will improve the functioning of oil film. Therefore, the Stribeck parameter also represents a qualitative equivalent of oil film thickness. In the HD regime, the film thickness is large; in the boundary regime, the film thickness is low. Accordingly, with the help of the Stribeck curve, it is possible to define transitions between
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µ ML EHL BL
ηv F Figure 8.4â•… The Stribeck curve.11 BL, boundary lubrication; ML, mixed lubrication; EHL, elastohydrodynamic lubrication.
regimes—which are defined as a function of λ. In fact, with the Stribeck curve, it is actually possible to better understand the effects discussed in previous sections such as film thickness, asperity contacts, and oil film friction. If the film thickness is high enough to prevent contacts, the wear will be low (or none), but very high film thickness also causes frictional losses due to viscous drag, as evident at the very right-hand side of the Stribeck curve. Therefore, it is not the best to have too large a film thickness, but as low as possible—without asperity contacts. In this case, the Stribeck curve reaches its minimum value. However, if the oil film were then further decreased, the asperities would start to come in contact, and this is characterized as ML. This transition (minimum friction) is defined to occur at λ╯=╯3.0. It is clear that friction will then increase due to solid contacts. If the film thickness decreases even further, more and more solid contacts would occur, friction would further increase, as would wear, and the regime would eventually change into BL (λ╯=╯1.0). These effects can be regulated with load, velocity, and viscosity. At very low Stribeck parameter values, friction reaches the value of unlubricated solid contacts. The transitions can thus qualitatively be represented also with a slightly different diagram than shown in Figure 8.4, that is, with friction as a function of the lambda value. Figure 8.5 shows even more detailed relations between the regimes, lambda, and effects on lifetime based on occurrence of contacts.12 Since the lambda parameter was designed to determine the effect of contact conditions on lifetime, a more precise classification also includes λâ•›≥â•›4 for fully EHD with no effects on wear, and the ML regime is also divided up from 1 to 1.5 and from 1.5 to 3, indicating the different influences on wear and lifetime.9,11 The Stribeck curve is slightly different for the tribological systems where the contacts are conformal (HD) and nonconformal (EHD). Namely, if the oil film is thick, the viscous drag will be higher; thus, friction increases in the region of λ╯>╯3 will be greater than at highly concentrated contacts, where the film is very thin. On the other hand, the transition from the EHD regime with low oil film thickness and high loads will be much faster than in the case of the HD regime. Thus, the friction curve slope in the mixed regime will be steeper (λ between 1 and 3); also, the friction in the boundary regime will be higher due to high contact pressures. Nevertheless, the overall lubrication strategy and shape of the Stribeck curve would remain the
â•… 109
REFERENCES
Figure 8.5â•… A schematic of a typical classification of different lubrication regimes depending on λ parameter. BL, boundary lubrication; ML, mixed lubrication; EHL, elastohydrodynamic lubrication.12
same. However, more recent studies have shown that Stribeck curve and lubrication regimes change with solid–liquid interface properties, such as wetting and/or surface energy, and that these effects should be included in proper lubrication design.12
REFERENCES ╇ 1â•… O. Reynolds. On the theory of lubrication and its application to Mr. Beauchamp Tower’s experiments, including an experimental determination of the viscosity of olive oil. Philos. Trans. R. Soc. 177 (1886), 191–203. ╇ 2â•… W. A. Gross, L. A. Matsch, V. Castelli, A. Eshel, J. H. Vohr, and M. Wildman. Fluid Film Lubrication. Wiley, New York, 1980. ╇ 3â•… B. J. Hamrock. Fundamentals of Fluid Film Lubrication. McGraw Hill, New York, 1994. ╇ 4â•… A. M. Ertel and A. N. Grubin. Russian Texts on EHL Investigation of the Contact of Machine Components. Central Scientific Research Institute for Technology and Mechanical Engineering (Moscow), 1949. Book no. 39 (DSIR translation). ╇ 5â•… C. Barus. Isothermals, isopiestics and isometrics relative to viscosity. Am. J. Sci. 45 (1893), 87–96. ╇ 6â•… B. J. Hamrock and D. Dowson. Ball Bearing Lubrication—The Elastohydrodynamics of Elliptical Contacts. Wiley-Interscience, New York, 1981. ╇ 7â•… R. J. Chittenden, D. Dowson, J. F. Dunn, and C. M. Taylor. A theoretical analysis of the isothermal elastohydrodynamic lubrication of concentrated contacts. Part 1 and 2. Proc. R. Soc. Lond. 245–269 (1985), 271–294. ╇ 8â•… G. Nijenbanning, C. H. Venner, and H. Moes. Film thickness in elastohydrodynamically lubricated elliptic contacts. Wear 176(2) (1994), 217–229. ╇ 9â•… T. E. Tallian. On competing failure modes in rolling contact. ASLE Trans. 10 (1967), 418–439. 10â•… S. Hironaka. Boundary lubrication, in Tribology of Mechanical Systems: A Guide to Present and Future Technologies, J. Vizintin, M. Kalin, K. Dohda, and S. Jahanmir (Eds.). ASME Press, New York, 2004, 41–51. 11â•… G. W. Stachowiak and A. W. Batchelor. Engineering Tribology, 3rd ed. Elsevier ButterworthHeinemann, Oxford, UK, 2005. 12â•… M. Kalin, I. Velkavrh, and J. Vizintin. The Stribeck curve and lubrication design for non-fully wetted surfaces. Wear 267 (2009), 1232–1240.
SECTION
II
FRICTION AND WEAR OF STRUCTURAL CERAMICS
CHAPTER
9
OVERVIEW: STRUCTURAL CERAMICS In one of the preceding chapters, ceramics are introduced as a class of materials with specific properties and their mechanical properties were subsequently discussed. Before discussing wear properties of some of the structural ceramics, it is imperative to discuss in detail how toughness properties, which are one of the major drawbacks of ceramics, can be enhanced in many of the brittle ceramics. In this context, this chapter discusses transformation toughening in zirconia as well as the development of some non-oxide ceramics, such as Si3N4/SiAlONs. In discussing the physics and mechanics of the transformation-toughening phenomenon, different microstructural variables governing the stability and the transformability of tetragonal zirconia are emphasized. Finally, this chapter also mentions the properties of TiB2 ceramics.
9.1 INTRODUCTION Toughening mechanisms in ceramics1 can be broadly classified into two major types: one involving a process zone around the crack tip, and the other being associated with bridging of crack faces by reinforcements (fibers, whiskers, particulates, etc.). The process zone mechanism is considered to increase toughness in ZrO2-based materials, and important examples of this are transformation toughening2 and microcracking. The mechanism of crack bridging is realized in toughening the ceramic matrix composites (CMCs) and to an even greater extent in ceramics having microstructure with characteristic coarse elongated grains (e.g., Si3N4 or sialon ceramics). Among engineering ceramics, ZrO2 has an excellent combination of high fracture strength (∼1â•›GPa) and good fracture toughness (∼10â•›MPa m1/2). This can be attributed to the transformation-toughening mechanism, and therefore, this mechanism has received considerable attention during the last few decades. It may be noted here that the toughness obtainable with the transformation-toughening mechanism is more than in the particle-reinforced (e.g., ZrO2-toughened Al2O3 [ZTA]) or whiskerreinforced (e.g., Al2O3/SiCW composites) composites.3 From this perspective, the discussion in this chapter is primarily limited to the toughening mechanisms
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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pertinent to the process zone mechanism as observed in monolithic tetragonal zirconia (t-ZrO2) ceramics. Transformation toughening in stabilized zirconia was discovered in the 1970s4 and, since then, ceramic materials based on zirconia have received increasing attention. In the last few decades, partially stabilized zirconia (PSZ), tetragonal zirconia polycrystals (TZPs), and fully stabilized zirconia (FSZ), as well as zirconiatoughened/dispersed ceramics (ZTC/ZDC) have started influencing the ceramic market.5 Also, excellent mechanical properties combined with biocompatibility, wear resistance, and high chemical and corrosion resistance make ZrO2-based materials potential candidates for biomedical applications.6 In this context, the high toughness of ZrO2, compared with Al2O3 and other monolithic bioceramics is particularly relevant. Recent research results also show that fracture toughness significantly influences the friction and wear properties of yttria-stabilized TZP (Y-TZP) materials.7 Hence, toughness optimization is one of the key criteria in the development of Y-TZP materials for structural applications. In view of the potential applications, progress has been made both in understanding the underlying mechanisms of transformation toughening3,4,8–13 and in exploiting it to develop toughened materials. This has also been reflected by several international conferences devoted wholly to zirconia, along with a large volume of research literature published in scientific journals. This chapter, therefore, discusses the role of different microstructural variables in influencing transformation toughening. This is needed also to show how the toughness of TZP ceramics can affect wear properties.
9.2 ZIRCONIA CRYSTAL STRUCTURES AND TRANSFORMATION CHARACTERISTICS OF TETRAGONAL ZIRCONIA The crystal structures of the three major zirconia phases are shown in Figure 9.1a, and detailed information regarding crystallography of different ZrO2 polymorphs can be found in the literature.14 Cubic zirconia has the ideal fluorite structure, while other polymorphs (tetragonal and monoclinic) have distorted fluorite structure.15 Undoped zirconia undergoes the following phase transition during thermal cycling: 1170° C monoclinic(m − ZrO2 ) ← → tetragonal(t − ZrO2 ) 950° C °C °C 2370 → cubic(c − ZrO2 ) 2680 → liquid
The transformations between the polymorphic forms are important so far as the processing and mechanical properties (strength, toughness, etc.) of zirconia ceramics are concerned. It has been well documented in the literature3,4,6 that the tetragonal-to-monoclinic (t→m) transformation in zirconia is a reversible athermal martensitic transformation, associated with a large temperature hysteresis of around 200°C, a volume change of 4–5% and a large shear strain (14–15%). This results in crumbling of the sintered part made from undoped zirconia. Several dopants, such
9.2 ZIRCONIA CRYSTAL STRUCTURES AND TRANSFORMATION CHARACTERISTICS
â•… 115
TEMPERATURE
1170°C
2370°C
monoclinic
tetragonal
cubic
ALLOY OXIDE CONTENT ELASTIC CONSTRAINT GRAIN SIZE (a)
metastable t-ZrO2 transformed m-ZrO2 crack tip stress field
(b)
Figure 9.1â•… (a) The phase transformation in zirconia induced by thermal heating or cooling, by addition of dopant oxides, or by elastic constraints and grain size. (b) Schematic showing the stress-induced phase transformation of metastable tetragonal zirconia particles in the crack tip stress field. The arrows in (b) indicate the generation of compressive residual stress due to the transformation-induced volume expansion and the microstructural constraint.37
as yttria, ceria, magnesia, and calcia, are added to stabilize the high-temperature tetragonal and/or cubic phase in the sintered microstructure.4 Although c-ZrO2, t-ZrO2, and m-ZrO2 are the most commonly observed phases, other zirconia phases such as nontransformable tetragonal (t′-ZrO2)16 and rhombohedral (r-ZrO2)17,18 have been found to exist under certain conditions. The extreme stability of the t′-ZrO2 phase is related to a large extent to both the higher amount of stabilizer and microstructural features, such as finer domain size (∼0.1â•›µm). The t′-ZrO2 does not undergo stress-induced phase transformation. The martensitic transformation of tetragonal zirconia (t-ZrO2) to monoclinic (m-ZrO2) can be induced either during cooling or by external loading under isothermal conditions.3,4 Both transformation routes have their own relevance. While thermally induced transformation controls the amount of tetragonal phase that can be
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retained after thermal cycling, the stress-induced martensitic transformation of ZrO2 enhances the toughness of zirconia ceramics. The latter phenomenon is known as transformation toughening.
9.3 TRANSFORMATION TOUGHENING Transformation toughening is the phenomenon in which the tetragonal zirconia phase, retained in a thermodynamically metastable condition, transforms to stable monoclinic phase in the tensile stress field around a propagating crack.4 The associated volume expansion (4–5%) involved in the t→m transition introduces a net compressive stress in the process zone around the crack tip.19 This decreases the local crack tip stress intensity factor and the driving force for crack propagation. This results in a subsequent increase in the toughness of the material (see Fig. 9.1b). The mechanism of transformation toughening can be characterized by a change in the stress intensity factor
∆K = K tip − K ∞.
(9.1)
Crack tip shielding can only occur when the crack tip stress intensity factor (Ktip) is less than that due to the applied stress (K∞), that is, ΔK╯<╯0. The resultant toughening is generally expressed by the parameter ΔKc, which is nothing but −ΔK. For the effective contribution of transformation toughening, the retention of the maximum amount of t-ZrO2 at room temperature or at the application temperature with optimum transformability is required. In the transformation-toughening literature,3,4,12,17,19 transformability is defined as the transformation potential or the ability of the t-ZrO2 to transform to m-ZrO2 in the crack tip stress field.
9.3.1â•…Micromechanical Modeling In the last few decades, considerable effort has been invested in the development of a theoretical framework for predicting the toughness increase due to stress-induced transformation.4,15,17,20–23 There are two different mechanistic approaches to predicting the toughness increments (ΔK): the stress intensity approach and the conservation integral approach. A detailed mathematical study of the stress intensity formulation is given elsewhere.5,6,24 It is now well understood that enhanced transformation toughening is only achieved during crack growth, that is, as the transformation zone develops around the crack tip (see Fig. 9.2). This phenomenon is typically illustrated in a resistance curve (popularly known as an R-curve). The formation of an R-curve is depicted in three consecutive stages: the frontal zone, the partial zone, and the extended zone. If the long-range strain field around the transforming particles is purely dilatational, there would be no contribution from the frontal zone to the toughness (ΔK╯=╯0). Crack tip shielding will, however, be achieved during subsequent crack growth, that is, in the partial zone, in which only a finite fraction of the transformable tetragonal phase transforms. Finally, shielding of the crack tip will reach the asymptotic maximum level in the fully developed extended zone inside
â•… 117
9.4 Stabilization of Tetragonal Zirconia
0.22 E Vf εt h1/2(1 – v) extended zone
∆Kc partial zonc h ∆a
crack
transformation zone
frontal zone ∆a/h
Figure 9.2â•… Schematic representation of the evolution of R-curve in transformation-toughened ceramics.2
which the maximum amount of the tetragonal particles are assumed to have transformed into monoclinic phase. Both the analytical models formulated by McMeeking and Evans11 and Budiansky et al.15 predict a similar supercritical plane strain toughness increment (ΔKc),
∆K c = 0.22 fEε t h /(1 − ν),
(9.2)
where f is the volume fraction of the tetragonal phase transformed within the transformation zone, E is the composite modulus, and εt is the dilatational strain involved in the transformation. Hence, from the microstructural point of view, the extent of crack tip shielding due to stress-induced transformation is linked to the transformation zone size and the volume fraction of the transformable t-ZrO2 particles. For simplicity, most of the micromechanical models assumed that stressinduced transformation is totally controlled by hydrostatic stresses. However, the transformation being martensitic in nature, shear stresses play a vital role, particularly when the transformation involves deformation-induced twinning phenomena.25 Evans and Cannon2 modified the existing formulation to take into consideration the possible influence of the shear strain:
∆K c = 0.38 fEε h /(1 − ν).
(9.3)
This shows that, when both the shear and dilatational components of transformation strain are considered in the micromechanical model, a higher level of transformation toughening is predicted.
9.4 STABILIZATION OF TETRAGONAL ZIRCONIA In the context of transformation toughening, the transformability of stabilized tetragonal phase plays an important role12,17,19 and hence the stability and transformability
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of t-ZrO2 need to be discussed in greater detail. Subbarao suggested that the size, charge, and concentration of dopant cations have an influence on the stabilization.26 In experiments with different stabilizer additions, Kim observed27 that the tetragonality, expressed by the “c/a” ratio, is significantly modified by dopant additions and influences the tetragonal phase stabilization. Morinaga et al.28 argued that the displacement of the oxygen anions from their ideal positions in the fluorite structure results in a change in the electronic structure, which influences t-ZrO2 stability. Hillert29 and others30 proposed that the oxygen vacancies also play a significant role in the stabilization of the high-temperature tetragonal phase. Adopting the Kröger–Vink notation,1 one can write the defect reaction for zirconia doped with yttria as
2 Y2 O3 ZrO → 2YZr′ + 3Oox + Vo ,
(9.4)
where YZr′ indicates a Y atom substituting and occupying the Zr-lattice site, O represents an oxygen atom occupying a normal lattice site, and Vo expresses the vacancy formation in the oxygen lattice site. The increase in oxygen vacancies raises the disorder of the ZrO2–Y2O3 system and increases the stability of the tetragonal phase.31 It is now generally accepted that tetragonal zirconia can only be stabilized at room temperature, if a critical grain size is retained in the sintered microstructures. Garvie further suggested that there exists a critical size range (dcl–dcu), irrespective of the martensitic start temperature, within which the retained tetragonal zirconia particles can undergo transformation under the applied stress field.34 This argument implies that there is a lower bound to this critical size (dcl) below which a particle cannot transform even if the interaction energy density is more than the dilatational strain energy density related to the transformation. The upper limit of the range (dcu) occurs when the particles, being larger than dcu, would spontaneously transform to m-ZrO2 during cooling from the sintering temperature. This would lead to degradation in mechanical properties, particularly toughness. Finally, the important factors concerning the stability of tetragonal zirconia particles embedded in a ceramic matrix (Mg-PSZ, ZTA, etc.) include matrix constraint, residual stress due to thermal expansion and E-modulus mismatch, chemical composition, and the transformational nucleation barrier.33 Garvie has observed that the critical size range for the Ca-PSZ and the Al2O3–ZrO2 composite is 62–95â•›nm and 0.38–0.45â•›µm, respectively.34 x o
9.5 DIFFERENT FACTORS INFLUENCING TRANSFORMATION TOUGHENING Extensive efforts have been made in the recent past to understand the different factors such as microstructure, alloying, influencing the tetragonal zirconia transformation. The different parameters influencing the transformability of the tetragonal zirconia and thus the toughness are schematically shown in Figure 9.3. In the following subsections, the influence of different microstructural parameters on properties of t-ZrO2 and toughness is discussed. In the context of transformation toughening,
â•… 119
9.5 Different Factors Influencing Transformation Toughening
Transformation toughening (Stress induced t-ZrO2 transformation in crack tip process zone) -Grain size - Stabilizer amount/distribution
Toughness of ZrO2based ceramics
Residual stress - CTE and E-modulus mismatch between ZrO2 matrix and reinforcement phase
Microcrack Toughening ((Microcrack growth in crack tip process zone)
Figure 9.3â•… Summary of various toughening mechanisms as well as various factors contributing to toughness of the monolithic zirconia ceramics.
the transformability is defined as the ability of t-ZrO2 to transform to m-ZrO2 in the crack tip stress field.
9.5.1â•… Grain Size It is an established fact that the grain size34 has a strong influence on the transformability and toughness of tetragonal-zirconia ceramics. As discussed earlier, t-ZrO2 phase can be retained below a critical grain size or particle size. Ruiz et al. studied the effect of heat treatment on the fracture toughness of 3Y-TZP and explained the experimental results obtained in terms of the grain size.35 By annealing a 2Y-TZP ceramic at 1500°C for different holding times, Swain obtained tetragonal grains of sizes ranging between 0.4 and 1.9â•›µm.36 An investigation of the reported mechanical properties shows that transformation toughness increases linearly with grain size. The highest toughness, around 12â•›MPa m1/2, is obtained at tetragonal grain size of 1.9â•›µm. While studying the influence of yttria content on the critical grain size (dc), Lange observed that dc increases from 0.2 to 1â•›µm as the yttria content increases from 2 to 3â•›molâ•›%, respectively.38 In the two-phase (c╯+╯t) field, specifically in the compositions ranging between 3.0 and 7.5â•›molâ•›% yttria, the toughness decreases from 6.3 to 3.0â•›MPa m1/2 as the volume fraction of the retained transformable tetragonal phase decreases to zero. Lange also found that the tetragonal grain size in Y-TZP materials (sintered under identical conditions: 1400°C, 1 hour in air) having an yttria content in the range of 0.8–6.6â•›molâ•›% was dependent on the composition.38 The average zirconia grain size also decreased abruptly as yttria content increased from 0.8 to 1.4â•›molâ•›%, remained relatively constant between 1.4 and 4.5â•›molâ•›% yttria and increased at yttria content higher than 4.5â•›molâ•›%. In Table 9.1, literature data revealing the influence of grain size on toughness are summarized. In general, it is observed that the fracture
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TABLE 9.1â•… Summary of the Toughness of Stabilized ZrO2 Ceramics1
Yttria, molâ•›% 2
PS-1500°C, 1 hour PS-1500°C, 10 hours PS-1500°C, 80 hours
2.5
1.8 3 2 2.6 3 3 2
3 2.8
2 3 3
Grain size, (µm)
KIc (MPa m1/2)
Toughness measurement
0.51
5.0
ISB
0.89
6.8
1.87
12.4
0.29╯±â•¯0.1
4.8╯±â•¯0.36
Indentation
60% monoclinic
0.35╯±â•¯0.15
4.2╯±â•¯0.32
Indentation
— — — — — — —
— — 1.5 0.5 — — —
13.5╯±â•¯0.5 11.6 12 7 9.5 5.6 6
Indentation — — — SENB SENB Indentation
— ∼35% tetragonal ∼55% tetragonal ∼99% tetragonal
— —
10.6 5.0
SENB Indentation
—
5.5
Indentation
—
11.0
Indentation
∼100% tetragonal
—
11.0
Indentation
Tetragonal
0.9
4.8╯±â•¯0.2
Indentation
—
—
4.6╯±â•¯0.2
Indentation
— — —
0.4 0.3 0.83
Indentation
— — —
0.62 0.86 1.14
5.9╯±â•¯0.1 2.5╯±â•¯0.1 6.34 5.05 5.8 6.30 8.2
Processing
PS: 1400°C, 30 minutes (unmilled powder) PS: 1400°C, 30 minutes (milled, 1 hour) HP — — — — — PS- 1400°C, air, 15â•›minutes — Presintering at 1200°C, 12 hours Presintering at 1250°C, 12 hours 1200°C, 12 hours and HP 1450°C, 2 hours 1250°C, 12 hours and HP 1450°C, 2 hours 1650°C, 3 hours, air, PS HP 1400°C 40 minutes, Ar, 40â•›Mpa HP 1450°C, 1 hour, vac PS-1500°C, air
Phases Monoclinic 10% Monoclinic 15% Monoclinic 20% 78% monoclinic
Indentation SEPB Indentation Indentation Indentation
â•… 121
9.5 Different Factors Influencing Transformation Toughening
TABLE 9.1â•… Continued
Yttria, molâ•›%
Processing
Phases
2
PS-1500°C, 2 hours, air
—
1.5 2 3 2.8 2
1150°C, 2–5 hours, air
Fully tetragonal
HP 1450°C, 1 hour, vac
Fully tetragonal
Grain size, (µm)
KIc (MPa m1/2)
Toughness measurement
—
6.1╯±â•¯0.2 Kplateau 11.9–13.9 12.6–14.8 3.5╯±â•¯0.1 8.7╯±â•¯0.3 10.2╯±â•¯0.5
—
— — 0.3 0.19 0.49
Indentation Indentation
Different processing routes: pressureless sintering (PS) and hot pressing (HP) are indicated. Various toughness measuring techniques: single-edge notched beam (SENB), single-edge precracked beam (SEPB), and indentation strength in bending (ISB).
toughness increases with grain size. One important observation is that a very high toughness of 17â•›MPa m1/2 (measured by single-edge notched beam [SENB]) can be obtained with 2Y-TZP of grain size 1.4â•›µm.39 For 3Y-TZP ceramics, the highest toughness, 8.2â•›MPa m1/2 (indentation method), has been measured at a grain size of 1.14â•›µm.40
9.5.2â•… Yttria Content The effectiveness of transformation toughening is greatly influenced by the nature and amount of the dopant cations.24,25,41 To optimize the fracture toughness of arcmelted zirconia alloys with yttria contents between 1 and 3â•›molâ•›%, Sakuma et al. measured a maximum toughness of 15â•›MPa m1/2 in a sample with 1.8â•›molâ•›% yttria.42 Sakuma et al. also compared the evolution of toughness as a function of the yttria content obtained in Y-TZP ceramics processed by conventional sintering and arc melting.42 The toughness peak is obtained at around 2â•›molâ•›% yttria. Masaki et al. fabricated zirconia ceramics stabilized with yttria ranging between 1.5 and 5.0â•›molâ•›% using hot isostatic pressing (HIP).43 The fracture toughness of the ceramics thus obtained was observed to increase nonlinearly with decreasing yttria content from 2.5 to 2.0â•›molâ•›%, with a maximum of 20â•›MPa m1/2 for the 2Y-TZP. This toughness value is the highest ever reported for Y-TZP ceramics by any research group. Furthermore, the toughness value measured by Vickers indentation and indentation strength in bending (ISB) methods were comparable. Matsui et al. observed that the difference in mechanical behavior of Y-TZP ceramics with the equivalent average grain size should be related to the existence of critical yttria content (Xcr).44 The Xcr for a TZP ceramic with 0.3â•›µm average grain size lies in the range of 4–5â•›wtâ•›% yttria content. For the tetragonal phase with yttria content lesser than Xcr, the driving force for tetragonal to monoclinic transformation will be high during cooling from sintering. As a result, microcracking will take place due to the thermally induced spontaneous tetragonal transformation. Gao et al. obtained high strength of around 1000â•›MPa with a toughness of around 14â•›MPa m1/2
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in a 2.1â•›molâ•›% Y-TZP, with a grain size of 2â•›µm and 70% t-ZrO2 retention.45 Similarly, a critical grain size needs to be achieved in order to retain 90% tetragonal phase in the sintered microstructure.46 For example, the critical grain size at which 90% tetragonal phase can be retained in 2Y-TZP is around 0.25â•›µm. This critical grain size is found to increase with the stabilizer content. In general, toughness is observed to increase with decreasing yttria content (see Table 9.1). The maximum variation in toughness is noted both at 3â•›molâ•›% and at 2â•›molâ•›% yttria stabilization. This is mainly related to the variation in grain size and the use of different toughness measurement techniques. A high value of indentation toughness (more than 10â•›MPa m1/2) is measured at lower yttria content (less than 2â•›molâ•›%). A 2003 report reveals that indentation toughness up to 14â•›MPa m1/2 can be obtained with 1.5Y-TZP nanoceramics, sintered at 1150°C.47
9.5.3â•… Yttria Distribution Besides grain size and yttria content, research results also indicate that yttria distribution is another important factor influencing the t-ZrO2 transformation48 and fracture toughness of Y-TZP monoliths.32,49–53 While critically assessing the published literature, it is observed that Y-TZP materials, processed from yttria-coated ZrO2 powders and Y-ZrO2 co-precipitated powders reveal different fracture toughness trends.53,54 Finer tetragonal grains of coated Y-TZP resulted in high fracture toughness, whereas the reverse effect was observed in the co-precipitated materials, which show increased toughness with increasing grain size above a certain lower limit. The achievement of high toughness in co-precipitated ceramics is only possible by growing the t-ZrO2 grains by sintering at high temperature or by sintering followed by annealing for a longer time at high temperature.35,36 In the work of Basu and coworkers, a new route to develop the microstructure and to tailor the toughness of Y-TZP ceramics by means of mixing zirconia powders with varying yttria content (3 and 0â•›molâ•›%) has been reported.49–52 In this method, the toughness of TZPs can be considerably improved (up to 10â•›MPa m1/2) compared with co-precipitated 3Y-TZP ceramics by tuning the composition of starting powders, when all are sintered under identical conditions (1450°C, 1 hour). Detailed microstructural analysis revealed that the difference between the average grain size of the co-precipitated 3Y-TZP (Tosoh grade, 0.3â•›µm) and the powder-mixture-based 2Y-TZP (hereafter referred to as TM2 grade ceramic, 0.5â•›µm) ceramics remain limited to less than 200â•›nm.49 However, a significant difference in indentation toughness is obtained between co-precipitated 3Y-TZP (2.5â•›MPa m1/2) and the TM2 ceramic (10â•›MPa m1/2). The available data in the literature1 indicate that the fracture toughness for a co-precipitated powder based Y-TZP ceramic with an average grain size of around 0.5â•›µm is about 6â•›MPaâ•›m1/2. This value, however, is much less than the excellent fracture toughness of 10â•›MPa m1/2 obtained with the mixed grade 2Y-TZP. For coprecipitated Y-TZPs, the increase in transformation toughening due to an increased mean grain size is also expressed in an increased Ms temperature. The Ms temperatures of the co-precipitated 2Y-TZP and mixed grade 2Y-TZP (TM2) ceramics were found to be at 390 and 312°C, respectively.32 Based on the accepted theory for co-
â•… 123
9.5 Different Factors Influencing Transformation Toughening
precipitated powder-based ceramics, the Ms temperature for the 2Y-TZP ceramic (TM2) having an overall yttria content of 2â•›molâ•›% should have been higher because of its larger grain size. Another interesting observation is that the Tio3 ceramic shows excellent fracture toughness of 9â•›MPa m1/2 at an average t-ZrO2 grain size of 0.19â•›µm (Fig. 9.4). At this grain size, co-precipitated t-ZrO2 grains are very stable and hardly
(a)
(b) 3
Ceramic grade Tio3
6
Relative frequency of observations
Relative frequency of observations
7
5 4 3 2 1
Ceramic grade TM2
2
1
0
0 0.0
0.2
0.4
0.6
0.8
1.0
0.0
1.2
0.2
0.4
0.6
(c)
1.0
1.2
(d)
30
30 Ceramic grade Tio3
Ceramic grade TM2
25
25
20
20 Frequency
Frequency
0.8
Average grain size (µm)
Average grain size (µm)
15
15
10
10
5
5
0
0 0
1
2
3
4
5
6
Y2 O3 content (mol %)
(e)
7
8
0
1
2
3
4
5
6
7
8
Y2 O3 content (mol %)
(f )
Figure 9.4â•… The sintered microstructures (a,b), grain size distributions (c,d), and EPMA yttria distribution (e,f) of two high-toughness Y-TZPs processed from 2.8â•›molâ•›% Y-coated starting powder (Tio3) and mixture of 3 and 0â•›molâ•›% Y-containing powders (TM2 ceramic).49
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susceptible to transformation. This indicates that Y-TZPs, based on mixing of 3YZrO2 powders with undoped ZrO2, show different toughening behavior compared with Y-TZPs processed from co-precipitated or Y-coated powders. The yttria distribution in the TiO3 ceramic (Fig. 9.4e) is broad and symmetric around 3â•›molâ•›% yttria, with quite a large number of electron probe microanalyzer (EPMA)–analyzed areas having an yttria content below 3â•›molâ•›%. EPMA data are reflection of the fact that yttria diffusion took place during the sintering process. Grains with low yttria content are considered to be very susceptible to transformation, since transformability increases with decreasing yttria content. In contrast, TM2 ceramics exhibit a coarser microstructure with average grain sizes around 0.5â•›µm (Fig. 9.4). Although the yttria distribution in the TM2 ceramic reveals a higher frequency at 2â•›molâ•›% yttria, a number of EPMA-analyzed areas having less than 2â•›molâ•›% yttria were also recorded. The predominant presence of t-ZrO2 measured in TM2 ceramic together with the acquired yttria distribution (Fig. 9.4f) clearly indicates that the yttria is redistributed during sintering, resulting in a microstructure containing tetragonal grains with nonuniform yttria distribution. Furthermore, the high toughness of the TM2 ceramic has been attributed to the increased transformability caused by the larger grains with lower yttria content (2â•›molâ•›% or less). It can be noted here that the co-precipitated 3Y-TZP ceramic shows a narrow yttria distribution with the highest frequency of yttria content of around 3â•›molâ•›%.49 Hence, the difference in yttria distribution could only explain the observed variation in toughness. Based on the experimental observations, the effect of grain size and yttria distribution on t-ZrO2 transformability is explained (see Fig. 9.5). The presence of
Figure 9.5â•… Schematic illustration showing the combined influence of grain size and yttria distribution on the t-ZrO2 transformability in Y-TZP ceramics.
â•… 125
9.6 Stress-Induced Microcracking
wider distribution of yttria and grain size leads to high transformability and tougher Y-TZP ceramic. Due to larger driving force, grains with lower yttria content and with larger tetragonal grain sizes transform to monoclinic phase. Subsequently, the transformation in the higher Y-containing grains occurs in an autocatalytic manner, a characteristic feature of martenstic transformation. More discussion of this plausible explanation of how the yttria and grain size distribution can contribute to enhanced toughness can be found in the literature.32,49,51
9.6 STRESS-INDUCED MICROCRACKING Conceptually, the microcrack toughening mechanism involves energy dissipation as microcracks nucleate and grow in size in the crack tip stress field.56,57 Microcrack toughening can initiate in the presence of the residual stresses, if the local tensile stress associated with the t–m zirconia transformation is sufficiently high. The shear strain involved in the martensitic t-ZrO2 transformation is released or accommodated in the form of twins and/or microcracks. These microcracks are preferentially nucleated at the grain boundaries of transformed m-ZrO2 grains, as shown in Figure 9.6. However, it has been argued that radial microcracking does not accompany the stress-induced (t–m) ZrO2 transformation. It was also commented that a given tetragonal grain either can transform in the crack tip stress field or, if already transformed, can cause microcracking. In the presence of microcracks around a primary crack, toughness can be enhanced by the growth of microcracks (dissipation of crack tip stress) and their interaction with the crack tip stress field. This lowers the local stress concentration at the crack tip, giving rise to primary crack tip shielding and resulting in toughening of the matrix. However, the presence of microcracks leads to a decrease in stiffness (effective elastic modulus). Following a continuum approach, Hutchinson studied the effect of microcracking on lowering the crack tip stress intensity, or shielding.56 It was also noted that the shielding contribution from microcracks is more for steadily growing cracks than for the stationary crack. It was predicted that the strong resistance curve behavior
Transformed m-ZrO2
Microcracks
Figure 9.6â•… Microcrack toughening induced in the zirconia ceramics as a result of the formation and the subsequent growth of microcracks, as a consequence of stress-induced t-ZrO2 phase transformation in the process zone.1
126â•…
CHAPTER 9â•… Overview: Structural Ceramics
associated with microcracking arises mainly from the release of residual stress. Faber presented a theoretical analysis to develop the microcrack toughening of zirconiabased ceramics for two conditions.58 The first is the case of toughening caused by stress-induced microcracking of residually strained monoclinic zirconia particles formed during cooling from sintering temperature. The second condition assesses the occurrence of microcracking as a result of stress-induced zirconia phase transformation. Faber58 formulated the toughness increment due to the second condition as
∆K M = 0.25 fEθε s h ,
(9.5)
where f is the microcrack density, θ is the dilatational strain involved with the stressinduced microcracking, εs is the residual strain at the saturated level of microcracking, E is the composite elastic modulus, and h is the half-width of the microcracked zone. Quantification of the toughness increase has shown that toughness can only be increased as a result of the stress-induced tetragonal zirconia transformation, and the magnitude of such increase strongly depends on the relative increase in permanent strain (due to dilatational crack opening) compared with the decrease in modulus. At this point, it is worthwhile to mention that the occurrence of microcracking is seen at tribological interfaces, as is discussed in a subsequent chapter.
9.7 DEVELOPMENT OF SIALON CERAMICS Among non-oxide ceramics, Si3N4/SiAlONs possess superior combinations of low density and high fracture toughness, hardness, and strength, as well as thermal shock resistance.59 These ceramics are being widely researched for their potential engineering applications.60 These materials are commonly used in diverse applications that require wear resistance and chemical stability at elevated temperatures, such as cutting tools, seal rings, valve seats, and cylinder liners, and in a variety of highefficiency engines and other mechanical systems.61 In view of such an array of applications, Si3N4-based materials represent a promising ceramic system. These materials also offer a number of opportunities for compositional or processing tailoring, thereby developing new Si3N4-derived materials (sialons) with varying microstructure and properties.59–62 In the SiAlON family, α-SiAlON and β-SiAlON are two well-studied phases with particular importance.59,62 The α- and β-SiAlONs are solid solutions of α-Si3N4 and β-Si3N4, in which some of the Si and N atoms are replaced by Al and O. The β-SiAlON ceramics have good sinterability and relatively high fracture toughness, associated with the typical rodlike grains. The α-SiAlON ceramics exhibit higher hardness, but their sintered microstructures are usually composed of equiaxed grains, which result in lower values of fracture toughness compared with the β-phase. The α- and β-SiAlON phases are totally compatible, and they are readily prepared by a single-stage sintering of the appropriate mixture of nitrides (Si3N4, AlN) and oxides (Al2O3). Therefore, mixed α–β SiAlON materials have received increasing attention due to their easier fabrication and higher chemical resistance compared with Si3N4 materials. Moreover, good mechanical properties can be obtained due to the combination of high hardness of α-SiAlON and good strength and toughness of β-SiAlON.
9.8 Microstructure of S-sialon Ceramics
â•… 127
The majority of engineering applications for β′/α′ ceramics remain in the lower temperature regime (generally less than 1000°C) and a greater market entry has been inhibited mainly by economic factors. In view of this, continuing research exists for novel SiAlON ceramics, which could be sintered at lower temperatures and might be more tolerant to less expensive starting powders. Of the new phases that have been identified within M-Si-Al-O-N systems (M is a cation in Group II of the periodic table), many have relatively high oxygen-to-nitrogen ratios, with lowered covalence and intrinsic properties, and those with high nitrogen content are difficult to sinter or to obtain with controlled phase content.54 A recently discovered SiAlON phase is the “S-phase,” having the composition M2AlxSi12−xN16−xO2+x with x╯≈╯2, where M is a cation in Group II of the periodic table.63 The larger cation (Ba or Sr) replaces some of the Al or Si and thus changes the possible compositions and the material properties. Research on this class of material is limited and the only work known to us is the characterization of Sr-doped S-phase with the help of x-ray diffraction (XRD) and scanning electron microscopy–energy-dispersive x-ray spectroscopy (SEM-EDS).54 Basu and coworkers have investigated the aspects of microstructure development, effect of microstructure on mechanical properties, and wear resistance properties of S-SiAlON ceramics.55,64,65
9.8 MICROSTRUCTURE OF S-SIALON CERAMICS A representative SEM image of a hot pressed and polished (unetched) surface of Ba-doped S-sialon ceramic (hot pressed at 1700°C, 2 hours) is presented in Figure 9.7. The acicular β phase and the residual silicate glass phase in an S-SiAlON
Figure 9.7â•… Backscattered electron (SEM) image showing minor phases (β′ and residual glass) in dark and light contrast relative to the major S-phase in an S-SiAlON ceramic, hot pressed at 1700°C for 2 hours.65
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CHAPTER 9â•… Overview: Structural Ceramics
ceramic can be distinguished in darker and lighter contrast, respectively, compared with the S-phase, which has an intermediate Ba content, dominating the atomic number contrast. Extensive phase analysis (semiquantitative) study with the help of SEM-EDS indicates that the hot pressed microstructure contains (4–5%) residual glass and (6–8%) α/β-Si3N4, besides the characteristic S-phase. Bright field transmission electron microscopy (TEM) imaging shows a microstructure dominated by a continuous array of elongated platelet crystals of S-phase with a minor intercrystalline liquid residue in triple junction channels (Fig. 9.8). The
glass glass
S β'
S APB
β'
200 nm
(a) ops Si
20
Ba
15 10
O Al
Ba
5 Ba 0
1
2
3
4
Ba 5
Ba Ba
6 Energy (keV)
(b)
Figure 9.8â•… (a) Bright field TEM image of S-phase crystals with residual intercrystalline glass, included β′, and antiphase boundaries (APBs) in an S-SiAlON ceramic, hot pressed at 1700°C for 2 hours. (b) TEM-EDS analysis of the glass phase.65
9.9 MECHANICAL PROPERTIES AND CRACK BRIDGING OF SiAlON CERAMIC
â•… 129
Figure 9.9â•… Phase equilibria of interest showing the possible composition range of the residual liquid phase, formed at the triple pocket during hot pressing of the investigated S-SiAlON ceramic.65
presence of fine β′ crystals within the S-phase suggests that they are formed initially in the sintering reaction and subsequently act as heterogeneous nucleants for the main phase. TEM-EDS analysis indicates that an average composition for S-phase within a specimen sintered typically at 1700°C is Ba2Si12−xAlxO2+xN16−x (x╯=╯2.0╯±â•¯0.2). The presence of the residual glass phase is clearly visible in Figure 9.8a. TEM-EDS analysis of the glass, as shown in Figure 9.8b, shows that the glass phase is Ba-rich aluminosilicate. Also, the possible composition range of the triple pocket glass is given in Figure 9.9.
9.9 MECHANICAL PROPERTIES AND CRACK BRIDGING OF SiAlON CERAMIC It is also important to assess how the mechanical properties of S-SiAlON ceramic compare with the existing SiAlONs. Like many other ceramics, the properties of various SiAlON ceramics are strongly sensitive to microstructural characteristics, in particular, grain shape or aspect ratio, intergranular phase, amount of seed crystals, ratio of α to β phases, and so on (e.g., see References 66–68). Table 9.2 presents the summary of the mechanical property data of some selected SiAlON ceramics; such data show a large variation in apparent hardness and toughness properties. Such a wide spectrum of properties can be uniquely described by two end members: the equiaxed α-SiAlON material with a combination of high apparent hardness (HV10╯=╯22â•›GPa) and low toughness (Kc╯=╯3â•›MPa m1/2); and β-SiAlON ceramic with high toughness (>6â•›MPa m3/2) and low hardness (∼15â•›GPa). The advantages associated with the two phases have been exploited in the development of “in situ selfreinforced multiphase α−β structure.” However, the optimization of fracture toughness occurs at the expense of high hardness of α-grain matrix. This fact led to
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TABLE 9.2â•… A Summary of the Literature Results on Mechanical Properties of Some Competing Si3N4-Based Materials along with Properties of S-SiAlON Ceramic, Hot Pressed at 1750°C
Material under study
Starting material and/or method of fabrication
S-SiAlON
α-Si3N4, AlN, Al2O3, BaCO3: HP at 1750°C, 2 hours in N2 atmosphere α-Si3N4, AlN, Al2O3, Y2O3: HP at 1900°C for 1 hour in N2 β-Si3N4, AlN, Al2O3, Si3N4: HP at 1950°C and heat treated for 12 hours at 1650°C α-Si3N4, Al2O3, AlN, SiO2, Lu2O3: HP at 1950°C for 2 hours in N2 α-Si3N4, AlN, Al2O3, Dy2O3: PLS at 1800°C and then GPS at 1900°C for 1 hour By post-sintering the nitrided compact at 1900°C for 3 hours α-Si3N4, with 20% AlN, Al2O3, AlN: Y2O3 in 9:1 molar ratio: GPS at 1900°C, 2 hours in N2 α-Si3N4, AlN, (Y,Yb)O3: HP at 1600–1750°C in vacuum HIP-ed Si3N4╯+╯MgO Si6−zAlzOzN8−z, z╯=╯0.3 α-Si3N4, AlN, Al2O3 with CeO2 (SHS-ed) CeO2: Y2O3 in 1:0:0.375 molar ratio: HP at 1750°C, for 1 hour in N2 α-SiAlON 76.92â•›wtâ•›% Si3N4, 13.46â•›wtâ•›% AlN, 5.77â•›wtâ•›% Y2O3, 3.85 Al2O3, reinforced with 10% β-SiAlON fiber Starting powders α-Si3N4, Al2O3, AlN, and Y2O3 Hot pressed (1650°C, 1 hour) Hot pressed (1650°C, 1 hour)
Y-α SiAlON (Y0.5S i9.3Al9.3O1.2N14.8) Nb-stabilized α-SiAlON Lu╯+╯5% Lu2SiO5 Dy-α SiAlON
Y-α SiAlON
Si3N4╯+╯20â•›volâ•›% (9:1 AlN/Y2O3) Y-Yb-α-SiAlON Si3N4 β-SiAlON (Ce-Y) α-SiAlON
α-SiAlON −10% β-SiAlON
α-SiAlON (equiaxed) Y2O3-Si3N4 Yb2O3-Si3N4
Hardness
Fracture toughness
Hv10 (GPa)
Kc (MPa m1/2)
14.0
11.6
21.0
12.0
21.7
6.3
18.9
4.4
18.8
6.3
18.5
5.1
18.5
5.1
18.9
4.6
17.0 15.5 15.2
6.0 4.7 4.9
14.8
5.9
22.0
3.0
15.2 13.9
7.1–8.1 7.4–8.5
Different processing routes mentioned are hot pressed (HP), pressureless sintering (PLS), hot isostatic pressing (HIP), gas pressure sintering (GPS), and self-propagating high-temperature synthesis (SHS).55
9.9 MECHANICAL PROPERTIES AND CRACK BRIDGING OF SiAlON CERAMIC
â•… 131
Figure 9.10â•… SEM image illustrating crack–microstructure interaction in S1750 ceramic. While the single pointed arrow indicates typical crack–S-phase interaction, the dotted arrow indicates the pullout of β-Si3N4 needles.55
the development of “single-phase in situ toughened α-SiAlON,” in which the coarse elongated α-SiAlON grains exhibit self-reinforcement and thus lead to improved toughness without impairing the high hardness of the α-SiAlON matrix. Although the toughness of S-phase SiAlON (11.6â•›MPa m1/2) is much higher than that of several Si3N4-derived ceramics, the apparent hardness of S-SiAlON (hot pressed at 1750°C), 15.3â•›GPa is much less compared with that of many of the Si3N4-based ceramics, possessing elongated α-SiAlON grain morphology. However, the hardness properties of Ba-doped S-SiAlON is comparable or modestly better than some competing silicon nitrides. To understand the toughening mechanisms, a representative SEM image showing the interaction of the propagating crack with the microstructure is presented in Figure 9.10. A typical, but not representative, case of multiple indentation cracking has been chosen to present evidence of important aspects of crack–microstructure interaction. Various observations can be summarized as (1) sinusoidal type trajectory or tortuous crack path (see upper crack) indicating increased crack growth resistance, (2) evidence of grain bridging and pullout, and (3) frequent occurrence of debonding of elongated S-phase grains surrounding the crack. From the preceding observations, it is obvious that crack bridging and grain pullout both result in moderately high toughness of Ba-S-SiAlON ceramic. Such observations can be analyzed in the light of a reported model that gives the correlation of fracture resistance with microstructure, containing elongated ceramic grains.
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Such a model is useful in predicting the frictional work, which is the main source of fracture energy. Also, the near tip toughening phenomenon, for example, elastic bridging caused by grain bonding, is ignored. Another assumption is that the decohesion of an elongated or platelet-shaped grain, before pullout, preferentially occurs along the grain boundary by intergranular fracture. On the basis of these assumptions, Becher and coworkers69 derived the following expression, correlating fracture toughness (K) with microstructural features:
K =C
∑ V W ( AR ) i
i
i
2
(9.6)
i
In this expression, V the is volume fraction of the elongated grains, W is the grain width, AR is the aspect ratio of elongated grains, and the suffix i denotes a group of the ith type of elongated grain. Also, the same value of C is taken for different elongated phases. Following the original model,70 the fitting parameter/constant C in Equation 9.6 can be expressed as
C=
Eτ . 6(1 − ν2 )
(9.7)
In this expression, E is elastic modulus, ν is Poisson’s ratio, and τ denotes interfacial friction. It is also clear from Equation 9.6 that the aspect ratio of the elongated phase, compared with width or volume fraction, should have more influence on toughness. On the basis of the preceding theoretical understanding, it should also be clear why elongated grain morphology can lead to better toughness in SiAlON-based ceramics. To summarize, the S-SiAlON ceramic represents a new class of SiAlON material, which offers a combination of E-modulus (more than 200â•›GPa), apparent hardness (up to 16â•›GPa), and toughness (up to 12â•›MPa m1/2). Also, this new material is different from many of the previously developed SiAlONs, as they show both increased apparent hardness and toughness with indentation load or contact pressure.
9.10 PROPERTIES OF TITANIUM DIBORIDE CERAMICS Structural ceramics are suitable materials for many tribological applications due to their unique combination of properties such as low density, low thermal expansion, high hardness and elastic modulus, good abrasion, and corrosion resistance.17 Titanium diboride (TiB2) is a non-oxide ceramic, with excellent combination of attractive properties: low density (4.52â•›g/cc), high hardness (∼22–25â•›GPa), high melting point (2200°C), and high elastic modulus (∼500â•›GPa).72 The combination of
9.10 Properties of Titanium Diboride Ceramics
â•… 133
high hardness and elastic modulus makes borides especially attractive for tribological and other engineering applications, such as cutting tools, wear-resistant parts, electrodischarge machining (EDM) electrodes, conductive coatings, cathode material for Hall–Heroult cells, and ballistic armor material.72 However, the large-scale use of monolithic TiB2 in engineering applications is limited due to its poor sinterability, brittleness, oxidation at high temperatures, and poor machinability. Titanium diboride is a stable intermetallic compound, as is evident from the Ti–B binary phase diagram (see Fig. 9.11a). In the Ti–B binary system, three
0
5
wt. % B 20
10
30
40 50 60
100
3225 L
3000
2200
2080
2000 1670 1500
2092
1540
βTi
1000
TiB
884
B
Ti3B4
Temperature (°C)
2500
TiB2
αTi 500 0
10
20
30
40
Ti
50 60 at. % B (a)
70
80
90 100 B
Ti
120° c
B
90°
90° a
b=a (b)
Figure 9.11â•… (a) Ti-B binary equilibrium phase diagram,74 indicating the possibility of formation of three intermetallic compounds, that is, TiB, Ti3B4, and TiB2. It can be noted that TiB2 has the highest melting point (3225°C) and little off-stoichiometric compositional variation. (b) The hexagonal unit cell of single crystal TiB2, a╯=╯b╯≠╯c, (a╯=╯b╯=╯3.029â•›Ao; c╯=╯3.229â•›Ao), α╯=╯β╯=╯90°, γ╯=╯120°, 1 formula unit per cell, Ti at (0,0,0), B at (1/3,2/3,1/2), and (2/3,1/3,1/2).75
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intermetallic phases, namely, orthorhombic TiB, orthorhombic Ti3B4, and hexagonal TiB2, are known to exist. While TiB and Ti3B4 undergo peritectic decomposition at 2180 and 2200°C, respectively, TiB2 melts congruently at 3225°C. Both TiB and TiB2 phases have a narrow range of homogeneity, whereas Ti3B4 has a fixed stoichiometric composition. TiB2 is present over a stoichiometry range of 28.5–30â•›wtâ•›% B. Owing to its stability and high melting point, TiB2 is considered as an important candidate material for high-temperature structural applications. From the point of view of crystal structure, the Ti atoms form a hexagonal close-packed (HCP) structure in TiB2. The hexagonal unit cell of single crystal TiB2, having space group P6/ mmm (a╯=╯b╯=╯3.029Å, c╯=╯3.229Å; α╯=╯β╯=╯90°, γ╯=╯120°) is presented in Figure 9.11b. Ti atoms are located at (0,0,0) and B atoms at (1/3,2/3,1/2) and (2/3,1/3,1/2) lattice sites. The high hardness and elastic modulus of TiB2 and also its chemical resistance are attributed to its inherent crystal structure and atomic bonding. In Table 9.3, the important physical and mechanical properties of various engineering ceramics are presented and compared with those of TiB2. From the listed data, it can be seen that TiB2 is superior to other advanced structural ceramics, such as Al2O3, B4C, and SiC, in terms of its hardness, toughness, and electrical and thermal conductivity. Particularly, the hardness of TiB2 (35â•›GPa) is greater than TiC (32â•›GPa), WC (24â•›GPa), and Si3N4 (25â•›GPa), except B4C (47â•›GPa). Although having high hardness, the engineering applications of monolithic TiB2 are restricted due to its poor sinterability, exaggerated grain growth, and poor oxidation resistance (at T╯>╯1000°C). Among various hard materials, TiB2-containing materials have a great potential for tribological applications due to high hardness and the lubrication efficiency of the boric acid film produced at the tribological interface. To make a comparison of the tribological properties of monolithic TiB2 with other structural ceramics, the reported data from the literature, as recorded with several other structural ceramics, such as Al2O3, SiC, B4C, Si3N4, WC, and WC-based cermet, are summarized in Table 9.4. It may be mentioned here that all the tribological data were obtained with unidirectional sliding test (mostly pin-on-disk configuration) at sliding speeds of 0.01– 4â•›m/s and with a load from 5–214â•›N under dry conditions. So far as the sliding wear rate is concerned, the material removal rate of all the tribocouples occurs under mild wear regime and varies in the order of ∼10−5–10−8â•›mm3/Nâ•›m, despite variation in load and sliding speed. While a very low wear rate of 10−8â•›mm3/Nâ•›m has been measured with a WC/Al2O3 tribocouple, the wear rate of the tribocouples with TiB2 is moderate and around 10−5 to 10−6â•›mm3/Nâ•›m. So far as the frictional property is concerned, TiB2 under sliding contact experiences a higher coefficient of friction (COF) of 0.5 or more, with the exception of a TiB2/SiC tribocouple. In addition it can be noted in Table 9.4 that both self-mated SiC and WC/Al2O3 sliding couples exhibit much lower COF of around 0.3. The reported variation in measured value of COF or wear rate should be attributed to the difference in tribological interaction at sliding interface. Apart from brittle fracture or grain pullout, as could be expected for brittle materials, the tribochemical reactions strongly dominate the tribological interaction in many sliding couples involving non-oxide ceramics. Particularly, the chemistry of tribochemical layer and protectiveness of such layer under dynamic sliding conditions are important factors to be considered to scientifically analyze the differences in
135
23.03 9.2 — 350 22–26 — 326.6 1100
10–30
5–7
560 25–35 700–1000
324.1 1100
hex 3000 6.1 6.83
ZrB2
hex 3225 4.52 αa-6.6 αc-8.6 60–120
TiB2
72 1100
450 37–47 300 71.6 1400
480 20–35 300–800
2.5–6
35.2 800
720 20–24 480–830
—
17
>105
106 3–3.5
29–121
hex 2600 15.7 5.2–7.3
WC
15–155
hex 2200 3.2 5.68
SiC
27.63
rhom 2450 2.52 4.5
B4C
fcc, face-centered cubic; hex, hexagonal; rhom, rhombohedral/trigonal; tet, tetragonal.76
Thermal conductivity (W/m/K) Electrical resistivity (10−6â•›Ωcm) Fracture toughness, KIC (MPa m1/2) Elastic modulus (GPa) Hardness (GPa) Three-point flexural strength (MPa) Enthalpy (kJ/m) Oxidation resistance (°C)
Crystal structure Melting point (°C) Density (g/cm3) Linear thermal expansion, α (10−6â•›K−1)
Property
183.8 1200
400 24–32 240–270
4
52
17–32
fcc 3067 4.93 7.42
TiC
Si3N4
750.5 1200
210 14–25 1000–1200
4–6
1018
20–24
hex 1900 3.44 2.4
TABLE 9.3â•… Summary of Important Physical and Mechanical Properties of TiB2 and the Other Important Ceramics
1580.1 >1700
400 18–21 323
2.5–4
1020
30.1
hex 2043 3.99 8.0
Al2O3
108.9 1700
384 13 —
2–2.5
21
50–221
tet 2050 6.3 8.4
MoSi2
136â•…
CHAPTER 9â•… Overview: Structural Ceramics
TABLE 9.4â•… Summary of the Unlubricated Tribological Data, along with Testing Conditions, Measured with Various Structural Ceramics and Comparison with the TiB2 Monolith76
Material TiB2 TiB2 TiB2 TiB2 SiC SiC SiC SiC Si3N4 WC WC–(6â•›wtâ•›%)Co B 4C Al2O3
Counterbody
Load (N)
SiC Al2O3 SiC Al2O3 Al2O3 SiC Al2O3 SiC Si3N4 Al2O3 Al2O3 B4C Al2O3
10 10 214 214 214 10 10 10 10 10 5,20,40 15 40
Sliding speed (m/s) 0.01 0.01 0.12 0.12 0.12 0.01 0.01 3.0 3.0 0.1 0.03,0.3 0.01 4
COF
Wear rate (mm3/Nm)
0.38 0.52 0.63 0.77 0.72 0.3 0.38 0.59 0.61 0.26 0.28–0.44 0.7 0.9
0.7╯×╯10−6 0.7╯×╯10−6 11.2╯×╯10−5 19.3╯×╯10−5 12.5╯×╯10−5 0.8╯×╯10−6 5.0╯×╯10−6 5.0╯×╯10−6 1.0╯×╯10−5 2.8╯×╯10−8 0.01–0.1╯×╯10−6 5.0╯×╯10−5 1.9╯×╯10−5
tribological properties. For example, oxidative wear together with adhesive and abrasive wear was the major mechanism for material removal in the case of a TiCcermet/steel sliding couple.73 In a tribological study70 on WC-based composites against steel, it has been observed that there is a transition from low COF of 0.1 to high COF of 0.5 with increase in the number of fretting cycles as well as load. Such a transition takes place at an early stage under higher load. Also, the wear involving WO3-rich tribolayer formation was observed on worn WC composites at higher load (>5â•›N).2 In Figure 9.12, the published tribological data, that is, steady-state COF and specific wear rate are provided. Considering the fact that fretting wear is a distinct type of reciprocatory wear process, the tribological data obtained with the conventional pin-on-disk sliding wear test or other type of wear tests involving longer relative displacement stroke between counterfaces under severe operating conditions are not included in Figure 9.12. From Figure 9.12a, it is clear that TiB2-based materials experienced much higher material loss due to wear in contacts with steel, even at lower Hertzian contact pressure of 750–800â•›MPa. However, at higher Hertzian contact pressure of 850–1200â•›MPa, much lower wear rate (that in contact with steel) as well as little change in wear rate is measured for TiB2-based materials. The wear rate data, presented in Figure 9.12a, therefore, indicate that TiB2-based materials are more wear resistant at higher contact pressure. Higher wear rate in contact with steel, compared with other counterfaces (alumina, hard metal, SiAlON) is primarily due to difference in wear mechanisms, as is discussed later. The wear rate and COF data, presented in Figure 9.12b, show that COF varies over a wide range of 0.3–0.7, depending on the counterfaces. An interesting observa-
9.10 Properties of Titanium Diboride Ceramics
â•… 137
100 –6
Wear rate, mm /N m (× 10 )
90 TIB2-Alumina TIB2-Hardmetal TIB2-Sialon TIB2-Steel TIB2Cermet-Alumina TIB2Cermet-Hardmetal TIB2Cermet-Sialon TIB2Cermet-Steel
80 70 60
3
50 40 30 20 10 0
–10 750 800 850 900 950 1000 1050 1100 1150 1200 1250 1300
Initial hertzian contact pressure, MPa (a)
3
–6
Wear rate, mm /N m (× 10 )
100 90 80
TiB2-Alumina TiB2-Hardmetal TiB2-Sialon TiB2-Steel TiB2 Cermet-Alumina TiB2 Cermet-Hardmetal TiB2 Cermet-Sialon TiB2 Cermet-Steel TiB2-Steel (Oil lubrication) TiB2-Steel (Water lubrication) TiB2 Cermet-Steel (Water lubrication)
70 60 50 40 30 20 10 0 –10 0.0
0.1
0.2
0.3
0.4
0.5
0.6
0.7
0.8
0.9
1.0
Coefficient of friction (COF) (b)
Figure 9.12â•… Variation in wear rate with Hertzian contact pressure under unlubricated fretting conditions (a) and with COF (b) for TiB2 and TiB2-based cermet against different counterbodies. TiB2-cermet has a sintered composition of TiB2–(16â•›volâ•›%)Ni3(Al,Ti); while that of monolithic TiB2 has TiB2–(2.5â•›wtâ•›%)SiC. Fretting parameters include a load of 8â•›N for 100,000 cycles with a frequency of 10â•›Hz and a displacement of 200â•›µm.76
tion is that a combination of lower COF (up to 0.5) and lower wear rate (up to 10╯×╯10−6â•›mm3/Nâ•›m) can be obtained during fretting of TiB2 against alumina and SiAlON counterfaces in an unlubricated condition. Also, both COF and wear rate are lower as expected, in the case of water and oil lubrication, when compared with the unlubricated fretting condition (see Fig. 9.12b). Additionally, to study the unlubricated tribological performance, it is also important to analyze the role of water or oil lubrication on friction and wear of tribological materials, such as TiB2. Basu et al.76 studied the mechanisms of material removal of TiB2-containing composites at fretting contacts against ball bearing steel in a lubricating medium. With a steel counterface, the frictional properties of TiB2
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(COF ∼0.08–0.12) measured in paraffin oil were much less compared with those in water (COF ∼0.3–0.4). In contrast to the water lubrication condition, the worn surface on TiB2 flat was found to be totally covered by an adherent carbon-rich (graphite) tribochemical lubricating layer after the fretting tests in paraffin oil lubrication. The tribodegradation of the paraffin oil, caused by frictional heat generation, is attributed to the carbon layer deposition. This observation matches well to the measured low friction and wear rate. The effect of relative humidity on friction and wear behavior of pressureless sintered TiB2 ceramic was studied against SiC and Al2O3 counterfaces under unlubricated conditions at room temperature.76 TiB2 showed low wear rates of less than 1╯×╯10−6â•›mm3/Nm, that is, high wear resistance, when rubbed against SiC and Al2O3 under dry unlubricated conditions. Interestingly, the COF showed a distinct dependence on relative humidity, but not the wear. Thus, experimentally it is proved that friction and wear are not necessarily correlated. The reason behind such an observation lies in the composition of the third body, mainly oxides, formed at the sliding interface. During the running-in period, the cavities formed due to grain pullout and microcracking act as reservoirs of oxide. As a result, tribo-oxidation was considered as the main wear mechanism.
REFERENCES ╇ 1â•… B. Basu. Toughening of Y-stabilized tetragonal zirconia ceramics. Int. Mater. Rev. 50(4) (2005), 239–256. ╇ 2â•… A. G. Evans. R. M. Cannon. Toughening of brittle solids by martensitic transformations. Acta Materialia. 34(5) (1986), 761–800. ╇ 3â•… M. Rühle and A. G. Evans. High toughness ceramics and ceramic composites. Prog. Mater. Sci. 33 (1989), 85–167. ╇ 4â•… R. C. Garvie, R. H. Hannink, and R. T. Pascoe. Ceramic steel? Nature 258 (1975), 703–704. ╇ 5â•… R. H. J. Hannink, P. M. Kelly, and B. C. Muddle. Transformation toughening in zirconia containing ceramics. J. Am. Ceram. Soc. 83(3) (2000), 461–487. ╇ 6â•… L. L. Hench. Bioceramics. J. Am. Ceram. Soc. 81 (1998), 1705–1728. ╇ 7â•… T. E. Fischer, M. P. Anderson, and S. Jahanmir. Influence of fracture toughness on the wear reduction of yttria doped zirconium oxide. J. Am. Ceram. Soc. 72 (1989), 252–257. ╇ 8â•… R. H. J. Hannink and M. V. Swain. Progress in transformation toughening of ceramics. Mater. Sci. 24 (1994), 359–408. ╇ 9â•… I.-W. Chen. Model of transformation toughening in brittle materials. J. Am. Ceram. Soc. 74(10) (1991), 2564–2572. 10â•… J. B. Watchman. Mechanical Properties of Ceramics. John Wiley & Sons, New York, 1996, 391–408. 11â•… R. M. Mcmeeking and A. G. Evans. Mechanics of transformation-toughening in brittle materials. J. Am. Ceram. Soc. 65(5) (1982), 242–246. 12â•… D. M. Stump and B. Budiansky. Crack-growth resistance in transformation-toughened ceramics. Int. J. Solids Struct. 25 (1989), 635–646. 13â•… I. R. Gibson and J. T. S. Irvine. Qualitative x-ray diffraction analysis of metastable tetragonal (t′) zirconia. J. Am. Ceram. Soc. 84(3) (2001), 615–618. 14â•… E. H. Kisi and C. J. Howard. Crystal structures of zirconia phases and thin inter relation. Key Eng. Mater. 153-154 (1998), 1–36. 15â•… B. Budiansky, J. W. Hutchinson, and J. C. Lambropoulos. Continuum theory of dilatants transformation toughening in ceramics. Int. J. Solids Struct. 19 (1983), 337–355.
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16â•… A. H. Heuer, R. Chaim, and V. Lanteri. The displacive cubic → tetragonal transformation in ZrO2 alloys. Acta Mater. 35(3) (1987), 661–666. 17â•… N. Mitra, K. Vijayan, B. N. Pramila Bai, and S. K. Biswas. Phase transformation introduced by mechanical and chemical surface preparations of tetragonal zirconia polycrystals. J. Am. Ceram. Soc. 76 (1993), 533–535. 18â•… D. P. Burke and W. M. Rainforth. Intermediate rhombohedral (r-ZrO2) phase formation at the surface of sintered Y-TZP’s. J. Mater. Sci. Lett. 16 (1997), 883–885. 19â•… D. J. Green, R. H. J. Hannink, and M. V. Swain. Transformation Toughening of Ceramics. CRC Press, Boca Bacon, FL, 1989. 20â•… J. C. Lambropoulos. Shape and orientation effects in transformation toughening. Int. J. Solids Struct. 22 (1986), 1083–1106. 21â•… A. G. Evans and A. H. Heuer. Review—Transformation toughening in ceramics: Martensitic transformations in crack-tip stress fields. J. Am. Ceram. Soc. 63 (1980), 241. 22â•… I. W. Chen and P. E. Royes Morel. J. Am. Ceram. Soc. 69(3) (1986), 181–189. 23â•… P. E. Royes Morel and I. W. Chen. Transformation plasticity of Ce2O3-stabilized tetragonal zirconia polycrystals I. Stress assistance and autocatalysis. J. Am. Ceram. Soc. 71 (1988), 343. 24â•… H. Schubert and G. Petzow. Microstructural investigations on the stability of yttria-stabilized tetragonal zirconia, in Science and Technology of Zirconia III, Advances in Ceramics, Vol. 24A. S. Somiya, N. Yamamoton and H. Yanagida (Eds.), American Ceramic Society Columbus, OH, 1988, 21–28. 25â•… K. Tsukuma, Y. Kubota, and T. Tsukidate. Thermal and mechanical properties of Y2O3-stabilized tetragonal zirconia polycrystals, in Science and Technology of Zirconia II, Advances in Ceramics, Vol. 12, N. Claussen, M. Rühle, and A. H. Heuer (Eds.). American Ceramic Society, Columbus, OH, 1984, 382–390. 26â•… E. C. Subbarao, H. S. Maiti, and K. K. Srivastava. Martensitic transformation in zirconia. Phys. Status Solidi A 21(9) (1974), 9–40. 27â•… D.-J. Kim. Effect of Ta2O5, Nb2O5, and HfO2 alloying on the transformability of stabilized tetragonal ZrO2. J. Am. Ceram. Soc. 73(1) (1990), 115–120. 28â•… M. Morinaga, H. Adachi, and M. Tsukada. Electronic structure and phase stability of ZrO2. J. Phys. Chem. Solids 44(4) (1983), 301–306. 29â•… M. Hillert. Thermodynamic model of the cubic tetragonal transition in nonstoichiometric zirconia. J. Am. Ceram. Soc. 74(8) (1991), 2005–2006. 30â•… S. P. S. Badwal and N. Nardella. Formation of monoclinic zirconia at the anodic face tetragonal zirconia polycrystalline solid electrolysis. Appl. Phys. A49 (1989), 13–24. 31â•… P. Kountouros and G. Petzow. In Advances in Ceramics 3, Science and Technology of Zirconia, A. H. Heuer and L. W. Hobbs (Eds.). American Ceramic Society, Columbus, OH, 1981, 202–216. 32â•… B. Basu, J. Vleugels, and O. Van Der Biest. Transformation behaviour of tetragonal zirconia: Role of dopant content and distribution. Mater. Sci. Eng. A 366(2) (2004), 338–347. 33â•… A. H. Heuer, N. Claussen, W. M. Kriven, and M. Rühle. Stability of tetragonal ZrO2 particles in ceramic matrics. J. Am. Ceram. Soc. 65(12) (1982), 642–650. 34â•… R. C. Garvie and M. F. Goss. Intrinsic size dependence of the phase transformation temperature in zirconia microcrystals. J. Mater. Sci. 21 (1986), 1253–1257. 35â•… L. Ruiz and M. J. Ready. Effect of heat treatment on grain size. Phase assemblage and mechanical properties of 3 mol% Y-TZP. J. Am. Ceram. Soc. 79(9) (1996), 2331–2340. 36â•… M. V. Swain. Grain-size dependence of toughness and transformability of 2 mol% Y-TZP ceramics. J. Mater. Sci. Lett. 5 (1996), 1159–1162. 37â•… T. H. Elmer. Engineered Materials Handbook, Vol. 4, Ceramics and Glasses. ASM International, The Materials Information Society, 1991. 38â•… F. F. Lange. Transformation-toughened ZrO2: Correlations between grain size control and composition in the system ZrO2-Y2O3. J. Am. Ceram. Soc. 69(3) (1986), 240–242. 39â•… J. Wang, W. M. Rainforth, T. Wadsworth, and R. Stevens. J. Eur. Ceram. Soc. 10 (1992), 252–257. 40â•… H. Kondo, T. Sekino, T. Kusunose, T. Nakayama, Y. Yamamoto, M. Wada, T. Adachi, and K. Niihara. Solid-solution effects of a small amount of nickel oxide addition on phase stability and mechanical properties of yttria-stabilized tetragonal zirconia polycrystals. J. Am. Ceram. Soc. 86(3) (2003), 523–525.
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41â•… S. Ramesh, C. Gill, and S. Lawson. The effect of copper oxide on sintering, microstructure, mechanical properties and hydrothermal ageing of coated 2.5Y-TZP ceramics. J. Mater. Sci. 34 (1999), 5457–5467. 42â•… T. Sakuma, H. Eda, and H. Sato. In Science and Technology of Zirconia III, Advances in Ceramics, Vol. 24A, S. Somiya, N. Yamamoto, and H. Yanagida (Eds.). American Ceramic Society, Columbus, OH, 1988, 357–363. 43â•… T. Masaki and K. Sinjo. Mechanical properties of highly toughened ZrO2–Y2O3. Ceram. Int. 13 (1987), 109–112. 44â•… M. Matsui, T. Soma, and I. Oda. Stress-induced transformation and plastic deformation for Y2O3containing tetragonal zirconia polycrystals. J. Am. Ceram. Soc. 69(3) (1986), 198–202. 45â•… L. Gao, T. S. Yen, and J. K. Guo. In Science and Technology of Zirconia III, Advances in Ceramics, Vol. 24A, S. Somiya, N. Yamamoto, and H. Yanagida (Eds.). American Ceramic Society, Columbus, OH, 1988, 405–414. 46â•… F. F. Lange and D. J. Green. Effect of inclusions size on the retention of tetragonal ZrO2: Theory and experiments, in Science and Technology of Zirconia V, S. P. S. Badwal, M. J. Bannister, and R. H. J. Hannink (Eds.). American Ceramic Society, Columbus, OH, 1981, 217–225. 47â•… O. Vasylkiv, Y. Sakka, and V. V. Skorokhod. Low temperature processing and mechanical properties of zirconia and zirconia-alumina nanoceramics. J. Am. Ceram. Soc. 86(2) (2003), 299–304. 48â•… T. C. Lei, G. Y. Lin, Q. L. Ge, Y. Zhou, and X. J. He. Morphologies of monoclinic phase in ZrO2 (2 mol% Y2O3) revealed by TEM in situ continuous observations. J. Mater. Sci. 32 (1997), 1105–1111. 49â•… B. Basu, J. Vleugels, and O. Van Der Biest. Toughness tailoring of yttria-doped zirconia ceramics. Mater. Sci. Eng. A 380 (2004), 215–221. 50â•… B. Basu, J. Vleugels, and O. Van Der Biest. Transformation behaviour of yttria stabilized tetragonal zirconia polycrystal-TiB2 composites. Mater. Res. 16(7) (2001), 2158–2169. 51â•… B. Basu. Zirconia-titanium boride composites for tribological applications. PhD thesis, Katholieke Universiteit Leuven, Belgium, 2001. 52â•… B. Basu, J. Vleugels, and O. Van Der Biest. Microstructure-toughness-wear relationship of tetragonal zirconia ceramics. J. Eur. Ceram. Soc. 24(7) (2004), 2031–2040. 53â•… H. Mandal. New developments in α-SiAlON ceramics. J. Eur. Ceram. Soc. 19 (1999), 2349–2350. 54â•… J. Grins, Z. Shen, M. Nygren, and T. Ekstrom. Preparation and crystal structure of LaAl(Si6−zAlz) N10−zOz. J. Mater. Chem. 5 (1995), 2001–2006. 55â•… B. Basu, Manisha, and N. K. Mukhopadhyay. Understanding the mechanical properties of hot pressed Ba-doped S-phase SiAlON ceramics. J. Eur. Ceram. Soc. 29 (2009), 801–811. 56â•… J. W. Hutchinson. Crack tip shielding by micro-cracking in brittle solids. Acta Metall. 35(7) (1987), 1605–1619. 57â•… D. L. Porter and A. H. Heuer. In Science and Technology of Zirconia II, Advances in Ceramics, Vol. 12, N. Claussen, M. Rühle, and A. H. Heuer (Eds.). American Ceramic Society, Columbus, OH, 1984, 653–662. 58â•… K. T. Faber. In Science and Technology of Zirconia II, Advances in Ceramics, Vol. 12, N. Claussen, M. Rühle, and A. H. Heuer (Eds.). American Ceramic Society, Columbus, OH, 1984, 293–305. 59â•… M. Zenotchkine, R. Shuba, and I.-W. Chen. Effect of seeding on the microstructure and mechanical properties of alpha-SiAlON: III comparison of modifying cations. J. Am. Ceram. Soc. 86 (2003), 1168–1175. 60â•… F. L. Riley. Silicon nitride and related materials. J. Am. Ceram. Soc. 83 (2000), 245. 61â•… S. Kurama, M. Herrmann, and H. Mandal. The effect of processing conditions, amount of additive and composition on the microstructures and mechanical properties of α-SiAlON ceramics. J. Eur. Ceram. Soc. 22 (2003), 109–119. 62â•… M. Zenotchkine, R. Shuba, J.-S. Kim, and I.-W. Chen. Effect of seeding on the microstructure and mechanical properties of alpha-SiAlON: I, Y-SiAlON. J. Am. Ceram. Soc. 85(5) (2002), 1254–1259. 63â•… Z. Shen, J. Grins, S. Esmaeilzadeh, and H. Ehrenberg. Preparation and crystal structure of a new Sr containing SiAlON phase Sr2AlxSi12−xN16−xO2+x (x∼2). J. Mater. Chem. 9 (1999), 1019–1022.
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64â•… Manisha and B. Basu. Tribological properties of a hot pressed Ba-doped S-phase SiAlON ceramic. J. Am. Ceram. Soc. 90(6) (2007), 1858–1865. 65â•… B. Basu, M. Lewis, M. E. Smith, M. Bunyard, and T. Kemp. Microstructure development and properties of novel Ba-doped S-phase SiAlON ceramics. J. Eur. Ceram. Soc. 26 (2006), 3919–3924. 66â•… Y. S. Zheng, K. M. Knowles, J. M. Vieira, A. B. Lopes, and F. J. Oliveira. Microstructure, toughness and flexural strength of self-reinforced silicon nitride ceramics doped with yttrium oxide and ytterbium oxide. J. Microsc. 201 (2001), 238–249. 67â•… Z. Y. Deng, Y. Inagaki, J. She, Y. Tanaka, Y. F. Liu, M. Sakamoto, and T. Ohji. Long crack R-curve of aligned porous silicon nitride. J. Am. Ceram. Soc. 88(2) (2005), 462–465. 68â•… M. Zenotchkine, R. Shuba, J. S. Kim, and I. W. Chen. R-curve behavior of in situ toughened α-SiAlON ceramics. J. Am. Ceram. Soc. 84(4) (2001), 884–886. 69â•… P. F. Becher, H. T. Lin, S. L. Hwang, M. J. Hoffmann, and I.-W. Chen. The influence of microstructure on the mechanical behavior of silicon nitride ceramics. Mater. Res. Soc. Symp. Proc. 287 (1993), 147–158. 70â•… R. B. Bhagat, J. C. Comway, Jr., M. F. Aateaue, and R. A. Brezler. Tribological performance evaluation of tungsten carbide-based cermets and development of a fracture mechanics wear model. Wear 201 (1996), 233–243. 71â•… B. Basu, J. Vleugels, and O. Vanderbiest. Unlubricated tribological performance of advanced ceramics and composites at fretting contacts with alumina. J. Mater. Res. 18(6) (2003), 1314–1324. 72â•… R. Riedel. Handbook of Ceramic Hard Materials, Vol. 2. Wiley-VCH Verlag GmbH, D-69469, Weinheim, Federal Republic of Germany, 2000, 968–990. 73â•… Z. Xingzhong, L. Jiajun, Z. Baoliang, O. Jinlin, and X. Qunji. Tribological properties of TiC-based ceramic/high speed steel pairs at high temperature. Ceram. Int. 24 (1998), 13–18. 74â•… J. L. Murray, P. K. Liao, and K. E. Spear. The B-Ti (boron-titanium) system. Bull. Alloy Phase Diagrams 7(6) (1986), 550–554. 75â•… R. G. Munro. Material properties of titanium diboride. J. Res. Natl. Inst. Stand. Technol. 105(5) (2000), 709–720. 76â•… B. Basu, G. B. Raju, and A. K. Suri. Processing and properties of monolithic TiB2-based materials. Intl. Mater. Rev. 51(6) (2006), 352–374. 77â•… B. Basu, J. Vleugels, and O. Van Der Biest. Fretting wear behaviour of TiB2-based materials against bearing steel under water and oil lubrication. Wear 250 (2001), 631–641.
CH A P T E R
10
CASE STUDY: TRANSFORMATIONTOUGHENED ZIRCONIA In an earlier chapter, it was stated that the friction and wear of ceramics depend on the material combination, their mechanical properties (combination of hardness and toughness) and microstructure (grain size, porosity, phase content, phase distribution), experimental parameters (normal load, sliding speed), contact geometry, the thermomechanical stress state in the contact area, the interaction with surrounding atmosphere (relative humidity [RH], water or other lubricants), and so on. From this perspective, this chapter demonstrates the important role of fracture toughness and atmospheric humidity in the wear resistance of yttria-stabilized tetragonal zirconia (Y-TZP) ceramics, an important class of ceramics having a good combination of transformation toughness and strength. Under the selected tribological conditions, phase transformation (tetragonal to monoclinic zirconia)-induced microcracking and spalling are found to play a major role in the wear of high-toughness TZP ceramics. An analysis of the surface topography in combination with Raman spectroscopy and x-ray photoelectron spectroscopy (XPS) will be shown to confirm the dominance of tribochemical wear in moist environments.
10.1 BACKGROUND One of the important aspects of ceramics tribology is to understand the effects of material property (fracture toughness) and environmental chemistry on the tribological properties. It has already been reported that this influence is different for oxide and non-oxide ceramics, as reported elsewhere.1 For example, it was found that microfracture and tribochemical reaction are the principal mechanisms in the wear of self-mated silicon nitride ceramics.2 It was observed that the reduction in wear rate is due to the formation of a lubricious silica layer formed by the tribochemical reaction of self-mated Si3N4 with the environment, especially water. However, the environmental effects on the friction and wear of Si3N4 ceramics largely depend on the type of sintering aids, the fabrication process, and the friction conditions.3 On Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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â•… 143
the other hand, it is observed that the wear rate is higher during the sliding of oxide ceramics in water, compared to that in air.4 Transformation-toughened Y-TZP ceramics with an excellent combination of mechanical properties (high fracture toughness and strength with moderate hardness) have been recognized as a promising candidate material for tribological applications.5 However, experimental results of the tribological behavior of self-mated TZP ceramics under dry sliding conditions have been disappointing.6 This is because of the extremely low thermal conductivity of tetragonal zirconia ceramics (≈2â•›W/m/°C) which causes high friction and wear in air and vacuum, thereby restricting their use in practical applications. However, zirconia ceramics show lower friction and improved wear resistance under different lubricating conditions. The coefficient of friction (COF) of self-mated zirconia contacts reduces to 0.1 in lubricating medium with hexadecane, in hexadecane with stearic acid, and in n-tridecane with stearic acid.1 Lower friction usually accompanies a decrease in wear rate. Similar decrease in friction and wear was also observed in other lubricating media with a mineral oil as the base oil7 and in more sophisticated synthetic lubricants, such as perfluoropolyalkylether.8 Since many technological applications (e.g., pump valves, bearings, biomedical implants) involve prolonged contact with water and/or water-based solutions, wear investigations on zirconia ceramics under water-lubricated conditions have also been carried out by several groups.9,10 The results are quite contradictory in the sense that wear of zirconia in water is found to increase8 and decrease9 compared with wear in air. It is difficult to compare the published results, owing to the various frictional conditions, material combinations, and test configurations involved in the experiments. Nevertheless, hydrothermal transformation of pure zirconia was undoubtedly shown to result in high wear, unacceptable for practical applications, when used in aqueous solutions.11–14 Since grain size critically determines the toughness of transformation-toughened ceramics, tribologists have been stimulated to understand the influence of grain size on friction and wear mechanisms of this class of ceramics.1–4 Birkby et al. investigated the impact of transformation toughening on the lubricated sliding wear behavior of 2 and 3â•›mol% Y-TZP ceramics against steel counterfaces.15 These investigators have proposed a qualitative relationship between the transformability and the wear of these ceramics. It has been reported that grain pullout and subsequent removal of material are the major mechanisms of wear for highly transformable ceramics. The dependence on grain size of the unlubricated sliding wear of 5.7â•›mol% Y-TZP ceramics against SiC was studied in dry nitrogen.16 This study showed that the wear follows a Hall–Petch type of relationship, that is, the wear rate is inversely proportional to the square root of the grain size for TZP ceramics having grains finer than 0.7â•›µm. Wear resistance of Y-TZP materials having a coarser grain size (>0.9â•›µm) was found to be inversely proportional to grain diameter. The same group of researchers also studied the influence of porosity of Y-TZP ceramics on the wear behavior of a TZP–SiC tribosystem under nitrogen gas.17 The wear rate of the TZP materials was found to increase by a factor of 5 as porosity increased from 1.5 to 7.0â•›vol%. In a different study,18 Jansen et al. investigated the effect of grain size and ceria doping on the tribological behavior of 5Y-TZP ceramics against SiC in distilled
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8N
WC-Co Y-TZP
200 µm, 10 Hz, 100,000 Cycles Figure 10.1â•… Schematic of the reciprocatory sliding experiment with low-amplitude displacement at interface.20
water and dry nitrogen. Surface fatigue-induced cracking and plastic deformation were identified as the predominant wear mechanisms in dry nitrogen. Grain pullout was the dominant wear mechanism in water. In a previous study, Fischer et al.19 established an inverse fourth power relationship between the fracture toughness and the wear rate of self-mated yttria-doped zirconia ceramics in dry air (35–45% RH). The investigated zirconia materials contained different levels of yttria stabilization (3–6â•›mol%), which is reported to change the phase assemblage from fully tetragonal to fully cubic in the materials studied. From the preceding perspective, this chapter summarizes the experimental results obtained while investigating the influence of microstructure and toughness20 as well as of humidity21 on wear properties. The experimental details can be found elsewhere.20,21 The study on microstructure/toughness–wear correlation was conducted on a Y-TZP/WC-Co couple, while the influence of humidity was studied on a self-mated Y-TZP couple. A schematic of the testing conditions is provided in Figure 10.1. By varying the yttria content, a range of Y-TZP ceramics was sintered to full density using the hot pressing route. Details of the materials processing can be found elsewhere.22 Table 10.1 summarizes the mechanical properties of the representative Y-TZP ceramics.
10.2 WEAR RESISTANCE The wear volume measured on the Y-TZP flats in the Y-TZP/WC-Co friction couples is plotted versus the t-ZrO2 grain size and fracture toughness in Figures 10.2 and 10.3, respectively. The error bars in the wear data represent the standard deviation in the data for at least three test results. The wear volume of the Y-TZP ceramics exhibits a strong correlation with grain size (see Fig. 10.2). In general, volumetric wear loss increases with an increase in the tetragonal grain size. It is also reported in the literature12 that the sliding wear rate of self-mated alumina ceramics increases linearly with increase in grain size. It is well known that the toughness of Y-TZP ceramics is controlled by grain size; hence it can be expected that the fretting wear resistance of such materials
10.2 Wear Resistance
â•… 145
TABLE 10.1â•… Mechanical Properties of the Y-TZP Ceramics with Varying Yttria Content, Hot Pressed to Full Density (>99.5% ρth) at 1450°C for 1â•›Hour in Vacuum
Ceramic grade
Overall yttria content
HV10 (GPa)
E (GPa)
KIC (MPaâ•›m1/2)
3.0 3.0 3.0 2.0 2.5 2.5 2.0 2.0
12.1╯±â•¯0.2 12.5╯±â•¯0.2 11.9╯±â•¯0.2 10.9╯±â•¯0.4 12.6╯±â•¯0.1 12.2╯±â•¯0.1 11.9╯±â•¯0.1 11.5╯±â•¯0.2
204 203 194 215 203 225 215 214
8.7╯±â•¯0.3 3.5╯±â•¯0.1 2.5╯±â•¯0.1 5.9╯±â•¯0.1 5.7╯±â•¯0.1 5.7╯±â•¯0.1 10.3╯±â•¯0.5 10.1╯±â•¯0.5
Tio3 D3 T3 T2 TM2.5 DM2.5 TM2 DM2
Tio3 grade is processed from commercial 2.7â•›mol% yttria-coated zirconia powders, while grades T3, D3 are from commercial 3â•›mol% yttria co-precipitated zirconia powders, respectively, and TM, DM grades are essentially processed from the mixture of T3/D3 powders with T0 powders (containing 0â•›mol% yttria) with the aim to obtain an overall yttria content in sintered ceramics of 2.5â•›mol% (TM 2.5, DM2.5) and 2â•›mol% (TM2, DM2), respectively.21
Wear volume, µm3
7e+5 6e+5
T2
Tio3 Yttria-coated power-based Y-TZP Y-TZP processed via “mixing route”
5e+5
DM2
Co-precipitated power-based Y-TZP
TM2 4e+5 3e+5
TM2.5 T3
D3
2e+5 0.15
0.20
0.25
0.30
0.35
0.40
0.45
0.50
0.55
Grain size, µm Figure 10.2â•… Plot of the average grain size (Y-TZP flats) against the measured wear volume (flats) for the investigated Y-TZP/WC-Co friction couples. The testing parameters are the same as in Figure 10.1. To study the influence of grain size on wear, two dashed lines connect and group the materials processed via two different routes: co-precipitation (T3, D3, and T2) and powder mixing (TM2.5, TM2, and DM2).20
should be dependent on grain size as well. Close observation of the data in Figure 10.2 further shows that at much finer tetragonal grain size, wear loss of yttria-coated Tio3 ceramic is quite high. The high volumetric wear loss of Tio3 ceramic can be in part related to the presence of cubic zirconia grains, which are inferior in wear resistance, as observed by Fischer et al.1 Considerably higher wear volume (after 100,000 fretting cycles) is measured after testing under dry conditions compared with that under ambient humidity (see
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50–52 % RH, after 100,000 cycles
1.2e+6
yttria-coated powder-based Y-TZP co-precipitated powder-based Y-TZP powder mixture-based Y-TZP
wear volume, µm3
1.0e+6 8.0e+5
T2
6.0e+5
Tio3 DM2
DM2.5
4.0e+5
T3
D3
2.0e+5 2
4
TM2
TM2.5 6
8
toughness, MPa m1/2
10
12
(a)
5–8 % RH, after 100,000 cycles
wear volume, µm3
1.2e+6 1.0e+6
Tio3
T2 TM2
8.0e+5 6.0e+5 D3 4.0e+5
yttria-coated powder-based Y-TZP co-precipitated powder-based Y-TZP powder mixture-based Y-TZP
T3 2.0e+5 2
4
6
8
toughness, MPa m1/2
10
12
(b)
Figure 10.3â•… Correlation between toughness and wear volume of Y-TZP flats in contact against WC-Co balls for 100,000 cycles under ambient (a) and dry (b) humidity conditions. The dashed and bold lines connect the materials processed via co-precipitation and powder mixing routes, respectively. The testing parameters are the same as in Figure 10.1.20
Fig. 10.3). This observation is not dependent on the toughness of the flat materials. Under unlubricated conditions too, wear loss of the co-precipitated ceramics increases with increasing toughness. The maximum in volumetric wear is noted with the Tio3 flat, the lowest with the T3 flat among the materials investigated. Again, the TM2 flat exhibits much less wear loss than the Tio3 ceramic.
10.3 MORPHOLOGICAL CHARACTERIZATION OF THE WORN SURFACES The worn surfaces were subjected to ultrasonic cleaning prior to microstructural investigation. Only mild polishing marks were observed on the worn low-toughness
10.3 Morphological Characterization of the Worn Surfaces
â•… 147
(a)
(b)
Figure 10.4â•… SEM micrographs, taken from the ultrasonically cleaned worn surfaces, showing the center of the wear scar on high-toughness Tio3 plate (a,b) after being tested against WC-Co in ambient humidity (50–52% RH) for 100,00 cycles. Fretting direction is indicated by a double-pointed arrow. The arrow in (a) indicates the occurrence of microcracking perpendicular to the sliding direction. The details of the cracking are clearer in (b). It can be noted that the cracking is not observed at the edge of the wear pit on Tio3 ceramic.20
materials (T3, D3) after testing for 10,000 cycles in both humidity conditions. Also, the occurrence of cracking is not pronounced on high-toughness materials (Tio3) even under dry conditions after 10,000 cycles. The morphology and details of the fretting pits on two different grades of Y-TZP flats (T3 and Tio3) after fretting for 100,000 cycles in ambient and low humidity conditions are shown in Figures 10.4 and 10.5, respectively. Abrasive grooves parallel to the fretting direction were formed in all the worn surfaces. The wear debris mainly accumulated around the edge of the wear pit. Compositional analysis, using energy-dispersive x-ray spectrometry (EDS), of the transfer layer on
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(a)
(b)
Figure 10.5â•… SEM micrographs showing the center of the wear pit on a high-toughness Tio3 plate (a,b) after it was tested against a WC-Co ball in dry conditions (5–8% RH) for 100,000 cycles.5 The fretting direction is indicated by a double-pointed arrow. Several microcrack patterns are observed to be oriented perpendicular to the sliding direction (b). The arrow in (a) indicates material removal by cracking-induced localized spalling. The spalled region is shown in detail in (b). No observable cracking is noted at the edges of the wear pit on Tio3 ceramic.
the flats reveals the presence of W and Co. Formation of WO3 and CoO in the transfer layer on the flat worn surface is indicated by XPS analysis of the tribolayer. Thus, it is evident that tribo-oxidation of the counterbody occurs during the process of fretting wear. The morphology of the cleaned worn surface on the high-toughness Tio3 plate shows the presence of microcracks in the central area of the wear pit. The microcracks (see Fig. 10.4b) appear perpendicular to the sliding direction. Similar microcracks were also observed on the fretted surfaces of T2, DM2, and TM2 ceramics (not shown). The factors affecting cracking on the worn surface are discussed in a
10.4 Zirconia Phase Transformation and Wear Behavior
â•… 149
subsequent section. Cracking was, however, not observed on the materials having low toughness (see Fig. 10.4a). A comparison of Figures 10.4 and 10.5 shows that the severity of cracking is more pronounced in worn Tio3 under conditions of low humidity (5–8% RH). Microcracking is only observed in the center of the worn surface. No pronounced microcracking pattern was seen at the edges of the fretting pit. Occasionally, the microcracking results in localized spalling and delamination of the material in the center of the worn surface on the Tio3 flat. This observation could also be correlated with the high friction recorded for the Tio3/WC-Co couple. The worn surface of the T3 flat having low toughness shows more severe scratches compared with that in ambient humidity conditions. However, no cracking is visible on the worn surfaces of T3 and D3 flats even in dry fretting conditions. Since greater volumetric wear is measured under dry atmospheric conditions, much attention is paid to investigating the worn surfaces after testing. The occurrence of cracking on the worn Tio3 ceramic is more obvious in Figure 10.6. The microcracks having a typical length of 2–3â•›µm are mostly seen perpendicular to the sliding direction. In addition, typical features of the worn surface on Tio3 and TM2 ceramics are also illustrated in Figure 10.6. Under severe conditions of microcracking, material removal from the worn Tio3 occurs by spalling across abrasive grooves. Although spalling induced by cracking is also observed on worn TM2, the severity is less compared with that on Tio3, as shown in Figure 10.6. This also correlates well with the measured wear data (see Fig. 10.3).
10.4 ZIRCONIA PHASE TRANSFORMATION AND WEAR BEHAVIOR The role of the tetragonal zirconia (t-ZrO2) to monoclinic zirconia (m-ZrO2) phase transformation during fretting wear of Y-TZP against hardmetal has been investigated with Raman spectroscopy. The Raman spectra, as shown in Figure 10.7, are obtained from different regions, at an interval of 20â•›µm, inside the wear pit (ambient humidity) on the worn Tio3, after being fretted against WC-Co for 100,000 cycles in ambient humidity (50–52% RH). The characteristic Raman peaks related to the tetragonal phase (located at 145 and 267â•›cm−1) and the monoclinic phase (located at 185 and 197â•›cm−1)23 have also been indicated in Figure 10.7. On comparing the Raman spectra taken from the native surface with those taken from the worn surface, it is observed that there is a considerable amount of monoclinic phase produced during the fretting process. Although, the intensity of the m-ZrO2 peak is quite low, the intensity of m-ZrO2 Raman bands acquired from the worn surface is clearly much higher compared with that from the unworn surface, where the monoclinic Raman bands can hardly be discerned from background. Also, x-ray diffraction (XRD) studies could not detect any monoclinic zirconia on the polished Tio3 specimen. The amount of transformed m-ZrO2 is clearly much higher in the center than at the edge of the wear pit, as shown by the intensities of the Raman band concerned with m-ZrO2. These observations provide clear evidence that the t-ZrO2–to–m-ZrO2 phase transformation occurs during the fretting wear process.
(a)
(b)
(c)
Figure 10.6â•… SEM micrographs showing the occurrence of material removal at the center of the worn surface by severe cracking-induced localized spalling at the center of the wear pit on Tio3 (a,b) and TM2 (c) flats after testing with WC-Co balls under dry conditions for 100,000 cycles.5 The fretting direction is indicated by a double-pointed arrow. The arrow in (a) indicates the details of microcracking.
150
10.4 Zirconia Phase Transformation and Wear Behavior
m
â•… 151
t
t
Relative intensity
center
20 µm
100 µm bulk
100
150
200
250
Raman shift,
300
350
400
cm –1
Figure 10.7â•… Comparison of the Raman spectra taken from the polished unworn surface (bulk) with those taken from different points in the worn surface (at an interval of 20â•›µm from the center of the wear pit) formed on the Tio3 ceramic after testing against WC-Co for 100,000 cycles in ambient humidity.20
It is a well-known fact24 that the martensitic transformation of tetragonal to monoclinic phase in Y-TZP ceramics is induced either thermally or by mechanical stresses (combination of both shear and tensile stress). As regards heat generation at the contacting surfaces, two interfacial temperatures should be considered: bulk and flash temperature.25 In general, the contact surface temperature is a function of the size and shape of the real contact area, the friction coefficient, normal load, sliding velocity, and thermal properties of the contacting bodies. The bulk interfacial temperature under unlubricated sliding conditions could increase to a steady-state temperature, which is mainly controlled by the contact geometry. The flash temperature is the short-duration microscopic temperature pulse mainly caused by the asperity–asperity contact and frictional dissipation of heat.25 It has been reported in the literature that the flash temperature for self-mated Y-TZP under dry conditions could rise to 500–800°C, depending on the combination of sliding speed and load.26 This flash temperature has been calculated for the pin-on-disk configuration and for a range of sliding speeds (0.002–0.570â•›m/s). Since the flash temperature strongly depends on the thermal conductivity of the mating surfaces, it is reasonable to assume that the flash temperature should be below 500°C for TZP/WC-Co tribocontact (thermal conductivity of WC-Co∼100â•›W/m/K) under the present fretting conditions (maximum sliding speed 0.004â•›m/s). Furthermore, dilatometry measurements22 showed that the TM2 and DM2 grade ceramics have martensitic start temperature (Ms) around 300°C. The Ms temperature of the Tio3 and all other TZP grades that have been investigated, however, lies below room temperature. From the preceding discussion, it is obvious that the t-ZrO2 phase transformation in high-toughness Y-TZP, such as Tio3 ceramic, under fretting conditions is only triggered by mechanical stress. On the other hand, both the repeated thermal cycling and mechanical stress can induce the t-to-m transformation at the fretting contacts of TM2 and DM2 grade ceramics.
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10.5 WEAR MECHANISMS Wear mechanisms as shown by the morphologies of the worn surfaces are investigated. If the microstructures of fretting wear pits under two different humidity conditions are compared (see Figs. 10.4–10.6), it is observed that the wear mechanisms are similar under both ambient and dry sliding conditions. The difference observed is in the degree of severity in abrasive scratches or microcracking and the associated material damage. In both cases, the area at center of the wear pit on the hightoughness Tio3 plate appears to be rather rough, with the presence of microcracks (oriented perpendicular to the sliding direction). It may be noted here that there is a higher concentration of microcracks in the central region of the worn surface than at the edges. Two mechanisms may be considered for the occurrence of cracking. First, continuous damage accumulation on the worn surface,2,4 induced by repeated stress cycles, could cause the cracking and subsequent material removal. The mechanics aspects of the initiation and propagation of fatigue cracks in fretted surfaces has already been studied in detail and reported in the existing literature.27,28 Further, careful observation of the worn surface (after fretting in ambient and low humidity conditions) shows that low cycle fatigue in the central region due to high local stress leads to the removal of material. Also, delamination, that is, the initiation and subsequent propagation of subsurface lateral cracks, leads to the detachment of material that needs to be considered. The other possible factor resulting in cracking is stress-induced tetragonal zirconia phase transformation at the tribocontacts. To understand the stress-induced transformation, the nature of the contact stresses needs to be analyzed.29,30 The mechanical stresses at the tribocontact depend on the macroscopic forces, the contact configuration, and the surface roughness of the moving parts.31 The Hertzian compressive stresses are the result of applied normal load and the surface roughness,32 whereas the tangential stresses are determined by the adhesive interaction between the tribosurfaces in the areas of real contact. The tangential stresses are reported to be modified significantly by the chemistry of the contacting surfaces.33 Particularly, the tangential stresses will be reduced in ambient humidity compared with those in dry conditions due to the presence of adsorbed water on the surface and tribochemically modified surfaces with lower shear strength. Following the classical Hertzian theory,32 the mean Hertzian stress field p(r) and the tangential stress field q(r) at any given distance r in the contact area with Hertzian contact radius a can be described as
p(r ) =
3P r2 1 − 2 πa 2 a2
(10.1)
where
3PR a= 4 E*
1/ 3
,
1 1 − ν12 1 − ν22 , = + E* E1 E2
r2 3µP q(r ) = 1− 2 , 2 a 2 πa
(10.2)
â•… 153
10.5 Wear Mechanisms
with P the normal load, μ the COF, and indices 1 and 2 referring to the sphere and the flat, respectively, for elastic modulus E and Poisson’s ratio ν. The Hertzian contact stress for a sphere-on-flat geometry varies with location within the contact circle, with the maximum Hertzian contact stress (3/2 times the mean pressure) being at a small distance beneath the surface at the center (r╯=╯0) of the contact circle. Also, it is clear from Equations 10.1 and 10.2 that the tangential stresses will have a maximum value around the center (at r╯=╯0) of the Hertzian contact area. Furthermore, it may be noted here that the tangential stresses are tri-axial in nature and are also variable with distance at the tribocontact. Assuming a concentrated tangential force, Hamilton and Goodman34 and Hamilton35 formulated the components of the tangential stress at the contact. They described the contact stress in terms of the von Mises effective stress and found that the maximum von Mises stress occurs below the surface. For example, in the absence of a tangential force (COF╯=╯0), the maximum von Mises stress is located at a distance of around 0.5a (a being the Hertzian contact radius) on the loading axis beneath the surface for a material with Poisson’s ratio of 0.3. In case of fretting, friction increases the magnitude of the von Mises stress and shifts it closer to the surface and toward the back edge of the sliding contact. For COF╯>╯0.3, the maximum von Mises stress shifts to the contacting surface (Refs. 24, 25). Also, the deviatoric (shear) component of the von Mises stress occurs at the tribosurface when high friction (COF╯>╯0.3) occurs at the tribocontact. It was also noted that tangential friction increases the compressive stress at the front edge, and intensifies the shear and tensile stresses at the trailing edge of the moving counterbody. A similar situation of contact stress is also noted in the work of Fischer et al.1 Hamilton observed that the maximum in principal tensile stress occurs at the surface.35 Using this theoretical background, Ginnakopoulos and Suresh developed the evolution of the surface and the subsurface stress fields for different levels of interfacial friction and the externally applied stress at the sphere-on-flat contact using three-dimensional finite element modeling (FEM).28 Their analysis reconfirmed that the location of the maximum tangential stress shifts from beneath the fretted surface toward the contacting surface with an increase in friction coefficient. FEM results were further supported by fretting experiments performed on the self-mated engineering steels. In the literature on transformation toughening, it is well recognized that a critical tensile stress is required to activate the stress-induced t-ZrO2 transformation.36 Evans and Cannon commented that the tetragonal zirconia transformation is more likely to be caused by shear stresses than by dilatational-hydrostatic stresses.37 This is more evident as a result of the fact that the shear component of the shape strain involved (∼0.16) in the t-ZrO2 transformation is almost 4 times larger than the dilatational strain (∼0.04). Also, Lambropoulos showed that the shear component of the critical tensile stress plays an important role in nucleating the transformed m-ZrO2, whereas the dilatational strain is assumed to contribute mostly to toughness enhancement.38 Following the preceding discussion, the occurrence of stress (tensile and shear)-induced tetragonal phase transformation on the tribosurfaces is inconsistent with the stress state at the sphere-on-flat contact. Also, the stress-induced t-ZrO2 transformation on the fretted surface would be much favored due to the reduced
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constraint compared with that in the bulk,39 as the constraints around a transforming tetragonal grain strongly influence its transformation behavior in the stress field. Since the t–m transformation is always accompanied by microcracking to accommodate transformation shape strain in the transformation zones,39 the microcracks observed on the worn surfaces of the highly transformable Tio3 ceramic (shown in Figs. 10.4–10.6) can be attributed to the tetragonal zirconia phase transformation induced by tangential stresses. Furthermore, according to the FEM calculations performed by Ginnakopoulos and Suresh, the maximum value of von Mises stress (corresponding to the maximum normal Hertzian compressive stress) should be shifted to the center of the Hertzian contact area on the tribosurfaces as considerably high COF (>0.3) was measured in our fretting experiments. This is also reflected in our experimental observations. First, Raman study indicates that the amount of the transformed m-ZrO2 is clearly much higher than that at the edges of the wear pit (see Fig. 10.7). Additionally, SEM investigation showed the presence of increased microcracking concentration at the center of the worn surface, while no microcracking is observed at the edges of the wear pit. Another interesting point observed in our experiments is that the fretting damage is mostly restricted to the surface, as is evident from the absence of deep cracks (see Fig. 10.6a). An important feature of the stress-induced transformation is the development of compressive stresses around the crack tip. This additional compressive stress prevents the growth of the microcracks at the tribocontacts.
10.6 RELATIONSHIP AMONG MICROSTRUCTURE, TOUGHNESS, AND WEAR It is an established fact that transformation toughening40,41 plays a dominant role in the enhanced toughness of Y-TZP ceramics. The effectiveness of transformation toughening depends to a great extent on the transformability of the t-ZrO2 phase, strongly influenced by grain size, dopant content, and dopant distribution and phase assemblage.21 The yttria-coated powder-based ceramic (Tio3 grade) has a bimodal microstructure consisting of large cubic zirconia grains embedded in a tetragonal matrix with finer grains. The comparatively high toughness of the Tio3 ceramics compared with T3 and D3 ceramics has been found to be caused by the nonuniform distribution of yttria.5 The high COF and volumetric wear loss of the Tio3 ceramic can also be in part related to the presence of cubic zirconia grains, which have inferior wear resistance as observed by Fischer et al.1 The inferior wear resistance of the cubic zirconia materials was reported to be due to brittle fracture. The co-precipitated powder-based ceramics (T3, D3, T2 grades) contain t-ZrO2 grains. Grain size, transformability, and fracture toughness increase with decreasing yttria content. The observed dependence of wear volume on toughness can be attributed to the fact that both the t-ZrO2 transformability and toughness increase in T2 and hence cause significant transformation at the tribocontact. The transformation
â•… 155
10.6 Relationship among Microstructure, Toughness, and Wear
of the tetragonal to monoclinic phase will result in microcracks at the fretting contacts, which will increase spalling of material from the contacting surfaces. Finally, the wear–toughness relationship established in our work shows a different trend in contrast to the experimental results obtained by Fischer et al.1 Results reported from this work have shown that the rate of sliding wear of self-mated yttriastabilized zirconia ceramics decreases with increasing toughness (fourth-power inverse relationship). In our study, however, the loss due to fretting wear of Y-TZP ceramics against WC-Co increases with increasing toughness depending on the microstructures of the TZP flats. This can be explained by two factors. First, the wear experiments performed by Fischer1 and co-workers are in the unidirectional sliding mode with the pin-on-disk setup, whereas the results presented in this chapter are concerned with the sliding of Y-TZP/WC-Co friction couples with the ball-onplate (fretting contact, reciprocatory sliding) configuration. The different test configuration and material combinations have an effect on the contact stress conditions. Second and more important, the zirconia materials investigated by Fischer and coworkers1 contain different yttria stabilization levels (3–6â•›mol%). Depending on the yttria level, the microstructures of the zirconia materials change from fully tetragonal to fully cubic. Fischer et al. studied only one fully tetragonal material (3â•›mol% Y-TZP, toughness 11.6â•›MPaâ•›m1/2 as measured with the diametral fracture toughness tests), which they thought to be representative of this microstructure.1 It is shown here that TZP includes a range of microstructures and properties depending on the preparation route. It is also shown here that the fretting wear resistance of the tetragonal zirconia ceramics critically depends on transformation toughness. The experimental results have quite significant implications for the tribological applications of transformation-toughened zirconia ceramics (see Fig. 10.8). The toughness of sintered Y-TZPs has been found to be very important in determining the fretting wear of Y-TZP/WC-Co friction couples. Since wear is a system property, this relationship between wear and toughness may not be true for other tribosystems containing Y-TZP and a different counterbody. However, the results discussed here strongly suggests that the optimization of toughness is critical to tailoring the wear resistance of Y-TZP ceramics. In fact, a trade-off between transformation toughness and wear resistance needs to be achieved for the successful application of Y-TZP ceramics in the field of tribology. This should be given due importance when designing Y-TZP-based materials for specific tribological applications. Compromise between toughness and wear resistance
Low t-ZrO2 transformability and poor toughness - abrasive wear (no cracking) - low wear rate
High tt-ZrO2 transformability and better toughness - transformation-induced cracking - localized spalling - high wear rate
Figure 10.8â•… Summary and implications of the experimental results presented in this chapter.
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10.7 INFLUENCE OF HUMIDITY ON TRIBOLOGICAL PROPERTIES OF SELF-MATED ZIRCONIA To study the effect of humidity, a different set of testing parameters is used; a schematic of the fretting test setup is shown in Figure 10.9. The influence of humidity on the frictional behavior of self-mated Y-TZP ceramic is shown in Figure 10.10. The material combination behaves similarly under different levels of humidity, whereby the coefficient of friction rises rapidly from a low value during the runningin period (first 100,000 cycles) and then reaches the steady-state value under high humidity conditions, where the friction value reaches the steady-state value at quite an early stage (first 10,000 cycles). Comparing the steady-state friction coefficient values, it is evident that a high coefficient of friction (1.05–1.075) is associated with fretting in both dry (5–8% RH) and ambient humidity (50–55% RH) conditions, whereas a relatively lower friction coefficient (0.6–0.625) is observed under high 8N
ZrO2 ZrO2
200 µm, 10 Hz, 500,000 Cycles Figure 10.9â•… Schematic illustration of the contact configuration with operating parameters.43
1.2 1.1
COF
1.0 0.9
85-90% RH 50-55% RH 5-8% RH
0.8 0.7 0.6 0.5
0
100,000 200,000 300,000 400,000 400,000
No. of cycles Figure 10.10â•… Evolution of the coefficient of friction (COF) in different humidity conditions when a Y-TZP flat is fretted against a commercial TZP ball under a load of 8N with displacement amplitude 200â•›µm and frequency 10â•›Hz.43
10.8 Wear Mechanisms in Different Humidity
Flat wear Ball wear
5e+5
4e+6
4e+5
3e+6
3e+5
2e+6 1e+6 0
5-8% RH
50-55% RH
2e+5
1.4e+4
Negligible ball wear
Wear volume on flat, µm3
5e+6
6e+5
85-90% RH
1e+5
Wear volume on ball, µm3
6e+6
â•… 157
0
Figure 10.11â•… Wear volumes of the self-mated Y-TZP tribocouple after the tests under different humidity conditions.43 Testing conditions are shown in Figure 10.9.
humidity (85–90% RH) conditions. Thus, a drastic reduction in friction of Y-TZP in contact with itself is observed in a moist atmosphere. This is indicative of a possible change in the wear mechanism under high humidity compared with that in the ambient and dry conditions. The wear volumes measured in the self-mated Y-TZP tribosystem after tests under different humidity levels are plotted in Figure 10.11. The error bars in the wear data indicate the scatter in the data (standard deviation) for at least three test results. A similar trend is seen in both the ball wear and the flat wear. Comparing the flat wear data, it is observed that wear volume under both ambient and dry conditions varies within the same order of magnitude as the highest wear under the dry condition, whereas it is reduced by two orders of magnitude under the high humidity condition (see Fig. 10.11). Furthermore, the ball wear data showed one order of magnitude less wear compared with the flat wear under dry and ambient humidity conditions (see Fig. 10.11). In contrast, the wear of the counterbody under the high humidity condition is too small to be studied and is reported as negligible wear in Figure 10.11. Considering the total wear (flat wear╯+╯ball wear) of the tribosystem, it is evident that the fretting wear rate of self-mated Y-TZP is reduced by two orders of magnitude under high humidity compared with that under dry and ambient humidity conditions. It should be noted here that similar behavior was seen in the wear of self-mated silicon nitride ceramics under different atmospheric humidity conditions.2 The wear rate of this material is reduced by three orders of magnitude as the ambient humidity is increased.
10.8 WEAR MECHANISMS IN DIFFERENT HUMIDITY The wear mechanisms as revealed by the morphologies of the worn surfaces were also studied (see Fig. 10.12). Comparison of the topographical features of the wear
(a)
(b)
(c)
Figure 10.12â•… SEM micrographs showing the details of the center of the as-worn surfaces developed under different environmental conditions: (a) 5% RH, (b) 50% RH, and (c) 85% RH. Double-pointed arrows indicate the sliding direction, while some of the microstructural features are indicated by a single-pointed arrow.43
158
10.8 Wear Mechanisms in Different Humidity
â•… 159
scars shows that the mechanisms of wear are similar under both ambient and dry sliding conditions, differing only in the degree of severity. This observation corresponds well with the friction behavior under the corresponding RH conditions, as shown in Figure 10.12. In both cases, the central region of the wear pit appears to be quite rough with the occasional presence of some microcracks (oriented perpendicular to the sliding direction), possibly resulting from high stresses and fatigue process. The occurrence of cracking at the fretting contacts has been studied in detail and reported in an excellent review article.42 The continuous removal of material by spalling, induced by the microcracks, might be one of the major causes of the higher wear rate under the low humidity conditions. A closer look at the regions indicated by the arrow in the microstructures presented in Figure 10.12a,b shows the sign of plastic deformation as previously observed by many researchers.1,10 Careful observation of the worn surface (after testing under ambient and low humidity conditions) shows that the removal of material from the contacting surfaces is mainly due to cracking. In the case of wear process induced by cracking, the location of maximum shear stress in the subsurface will shift downward with time, preventing enough plastic deformation from generating at the same location and thus leading to crack initiation and, subsequently, delamination. Therefore, the surface damage is likely to be more dominant than delamination. Transmission electron microscopy (TEM) studies10 of the worn surface on 3Y-TZP material (after sliding with Mg-PSZ in dry and water-lubricated conditions) provides evidence for dislocation flow in both the stabilized tetragonal and the transformed monoclinic phases. Although a different test configuration (pin-ondisk) is used in the work of Sanchez and Rainforth et al.,10 the loading conditions (10N/pin) are similar to that of our work (8N). As regards the sliding velocity, our tests are carried out at much lower speed (0.004â•›m/s) than the wear tests (sliding speed 0.24â•›m/s) reported in Ref. 10. Thus, in our case there is less likelihood of the ball losing contact with the flat. Considering all the aforementioned conditions combined with our microstructural observations, it may also be reasonable to suggest here that plastic deformation and associated delamination are the other possible mechanisms of material removal in this case. Based on observation of the surface topography of the worn surfaces, two mechanisms can be proposed for the high wear rate under low and ambient humidity conditions. The first mechanism is cracking-induced spalling, and the other is plastic deformation. Both of these mechanisms have high possibility of occurrence. The observation of cracking under dry and ambient conditions on worn surfaces is possibly due to tetragonal zirconia transformation, as discussed earlier in this chapter, supported by information obtained from Raman spectroscopy measurements. It can therefore be concluded that tribomechanical wear with possible material removal by cracking might be the major wear mechanism under dry and ambient humidity conditions.43 On the as-worn surface under high humidity condition (85% RH), a tribofilm is seen to occur all over the wear scar (see Fig. 10.12c). The morphology of the film shows that it is adherent to the worn surface. Some cracks are seen in this tribolayer. The chemical nature of the tribofilm is analyzed by XPS and is discussed in Section 10.9. Since the as-worn surface under 85% RH is fully covered by a tribofilm, the wear scar is ultrasonically cleaned before carrying out further investigation.
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10.9 TRIBOCHEMICAL WEAR IN HIGH HUMIDITY Chemical changes of the worn surface (tribofilm formation) produced in a humid atmosphere (85–90% RH) were analyzed with the help of XPS experiments. The details of the tribofilm are shown in Figure 10.13. Small area XPS (SAXPS) measurements were carried out in the middle of the wear track and at different spots on the wear debris accumulated outside the wear track. Two lens modes were used, with the diameter of the analyzed area being 50â•›µm and 150â•›µm, respectively. However, it cannot be ignored that part of the signal was coming from the native surface next to or under the wear debris even in the 50â•›µm lens mode. The XPS spectra obtained from both the native surface and the wear debris over a wide binding energy (BE) region show characteristic peaks of zirconium, yttrium, oxygen, and carbon. The carbon peak is due to the overlying contamination layer of hydrocarbon inherent in any XPS analysis. Figure 10.14a shows a comparison of the O 1s spectra obtained from the native surface (denoted as A), the wear track (SAXPS 150â•›µm, denoted as B), and two different spots on the wear debris next to it (SAXPS 150â•›µm [denoted as C] and SAXPS 50â•›µm [denoted as D], respectively). Spectrum A shows a peak at 529.7â•›eV and a shoulder at the high-BE side. The low-BE peak possibly corresponds to oxygen bound to metal, whereas the shoulder corresponds to oxygen in a metal–OH bond (about 531.5â•›eV44) and bound H2O (533.2â•›eV44). The main peak of spectrum B shifts to 530.1â•›eV and its high-BE shoulder is much higher than that of A, indicating an increase of the OH content on the worn surface. The shape of O 1s observed at position C is in between that of spectra A and B, whereas spectrum D is similar to
Figure 10.13â•… Nomarski interference micrograph showing the as-worn surface on the flat after fretting test under high humidity condition. XPS spectra have been acquired from different locations on the debris (C and D), the middle of the worn surface (B), and the native surface (A). Fretting conditions are the same as shown in Figure 10.1. The fretting direction is indicated by a double-pointed arrow.43
10.9 Tribochemical Wear in High Humidity
â•… 161
300 O 1s
Intensity (a.u.)
250 200
D
150
B
100
C A
50 536
534
532 530 Binding Energy (eV) (a)
528
9000
Zr 3d
8000 7000
A
6000 CPS
526
5000 4000
B
3000 2000 1000 188
C 186
184 182 Binding Energy (eV) (b)
178
180
550
Y 3d
500
CPS
450
B
A
400 350 C
300 250 165
160 155 Binding Energy (eV) (c)
150
Figure 10.14â•… O 1s (a), Zr 3d (b), and Y 3d (c) XPS spectra recorded on the native surface (A), in the middle of the wear pit (B), and on the wear debris (C and D) produced during the fretting wear of self-mated Y-TZP tribocouple under highly humid (RH 85–90%) atmosphere. The spectra were normalized to the same maximum intensity for ease of comparison.43
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CHAPTER 10â•… Case Study: Transformation-Toughened Zirconia
B. This indicates that there is similarity in the chemical state of oxygen in the wear debris (D) and on the worn surface (B). The shape of spectrum C can be explained assuming that the signal comes from both the wear debris and the native surface below it. Figure 10.14b reveals the Zr 3d spectra obtained at the same spots as the corresponding O 1s lines just discussed (the notation is the same). The Zr 3d spectrum for case D has been omitted for clarity. The Zr 3d5/2 peak of spectrum A is at 181.8â•›eV, which is slightly lower than that cited in Reference 44 for Zr in ZrO2 (182.2â•›eV). The Zr 3d5/2 peak of spectrum B shifts by 0.4â•›eV toward higher BE (182.2â•›eV). Spectrum C exhibits a Zr 3d5/2 peak with a value of BE between those of spectra A and B. This spectrum C has a larger full-width at half maximum (FWHM), while both shape and peak position of spectrum D (not shown) are close to those of B. There is similarity in the chemical state of Zr in the wear debris (D) and on the worn surface (B). The Zr 3d spectra recorded from the debris (case C) can be well fitted by using spectra A and B (see Fig. 10.14b), showing that the signal came both from the debris and from the native surface. The BE of the Zr 3d5/2 electrons in the Zr–OH state is stated to be higher than that in ZrO2 by 0.5â•›eV.45 Taking into consideration the change in the O 1s spectra of the wear track indicating a high amount of OH present there, it may be assumed that hydroxylated Zr species are formed due to the fretting process at high humidity. Another possible reason for the shifting of the Zr 3d peak in the wear track is the tetragonal-to-monoclinic phase transformation of zirconia due to the wear process, as discussed in the previous section. This reasoning might be true, taking into consideration that the number of nearest neighbors surrounding the Zr atoms changed from 8 to 7 as the tetragonal zirconia phase transformed to monoclinic symmetry.24 However, no Zr 3d peak shift was reported46 after a t–m phase transformation of zirconia partially stabilized by Y during annealing in water at 353â•›K, possibly due to insufficient energy resolution. On considering Figure 10.14, we observe that the amount of phase transformation is not significant enough to cause the observed shift in the Zr 3d spectra. Thus, it is well justified to suggest here that formation of a considerable amount of Zr–OH compounds on the worn surface is the major cause for the measured shift in the Zr 3d spectra. Significant changes are seen on comparing the Y 3d spectra acquired from the native surface with those from the wear track and the wear debris (see Fig. 10.14c). For the native surface (case A), the Y 3d5/2 peak is at 156.6â•›eV, which agrees with the data for Y2O3.44–46 The Y 3d peak recorded in the wear track (B) shifts by 1.2â•›eV toward higher BE (157.8â•›eV). Spectrum C is more complex, showing three distinct peak; spectrum D (not shown) is very similar to B. Similar to the Zr 3d peak (case C), a good fit of the Y 3d spectrum from the debris (case C) is obtained by using the Y 3d spectra recorded on the native surface (A) and in the wear track (B). It is thus suggested that the signal recorded in position C is a mixture of A and B (see Fig. 10.14c). The XPS results clearly show that fretting under high humidity causes a significant change in the local chemical bonding of yttrium ions. It is assumed to be due to the formation of Y–OH bonds, as a result of the corresponding change in the O 1s spectrum discussed earlier. Similar observations on the possible formation of
10.10 Closing Remarks
â•… 163
yttrium hydroxide on the surface of water annealed Y–partially stabilized zirconia specimens together with a t–m phase transformation are reported in Reference 46 with a shift of Y 3d peak by ∼2â•›eV toward higher BE from the BE for Y2O3. Summarizing the XPS results, it is seen that the presence of a shoulder and the increase in intensity of this shoulder in the O 1s spectra obtained from the worn surface and wear debris (see Fig. 10.14a) strongly indicate the formation of hydroxides in the fretting pit under high humidity conditions. Furthermore, the evolution of the Y 3d (see Fig. 10.14c) and Zr 3d spectra (see Fig. 10.14b) in the worn surface indicates the formation of hydroxide compounds of Zr and Y atoms. On combining the XPS results with the friction and wear data, it is seen that the possible formation of Zr–OH and Y–OH compounds could serve as a protective lubricating layer during fretting at 85–90% RH condition, thus reducing friction and wear loss. This further suggests that tribochemical wear is the predominant wear mechanism under high humidity conditions. Thus, the beneficial effect of humidity in lowering the friction and wear of the self-mated Y-TZP materials investigated in the work discussed here may be attributed to tribochemical wear. It may be recalled here that a protective lubricating film (SiO2·2H2O) is also formed in self-mated Si3N4 contacts under conditions of high humidity.2
10.10 CLOSING REMARKS The following key points emerge: a. The results presented in this chapter suggest that fracture toughness is an important parameter in determining the tribological behavior of engineering ceramics. Within the range of toughness investigated, volumetric wear loss increases with increasing toughness for the co-precipitated and powder mixture-based Y-TZP ceramics. Wear volume measured for the yttria-coated powder-based ceramics is much higher due to the presence of a significant number of cubic grains. b. Raman spectra obtained from the worn surfaces clearly show that the wear is accompanied by the tetragonal-to-monoclinic phase transformation for highly transformable TZP ceramics. The t–m ZrO2 transformation is observed to cause extensive microcracking on the worn surfaces in both ambient and dry environments. c. Based on the observations of the topography of the worn surfaces, the fretting wear mechanism is proposed. The morphology of the wear pits indicates that a mild abrasion is the major mechanism of wear for low-toughness Y-TZP ceramics. On the other hand, microcracking-assisted spalling is the major wear phenomenon for high-toughness Y-TZP materials. d. Tribological study under different environmental condition with varying levels of humidity indicates that the severity of wear decreases with increasing relative humidity in the surrounding atmosphere. The largest wear damage accompanied by the maximum wear volume is observed with the lowest humidity level, 5╯±â•¯8% RH.
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e. The morphology of the wear scar in both dry and ambient sliding conditions indicates the occurrence of plastic deformation and extensive cracking-induced delamination. Also, tribomechanical wear as a result of transformation-induced cracking might be the major wear mechanism under dry and ambient humidity conditions. f. XPS spectra from the tribofilm formed under high humidity conditions show the possible formation of Zr–OH and Y–OH compounds in the wear debris. The formation of these compounds possibly serves as a protective layer in reducing the friction and wear of 2.5â•›mol% Y-TZP ceramic under sliding against a commercial 3Y-TZP ball in a highly humid atmosphere (85╯±â•¯90% RH). Tribochemical wear is found to be the predominant wear mechanism under high humidity conditions.
REFERENCES ╇ 1â•… T. E. Fischer, M. P. Anderson, S. Jahanmir, and R. Salher. Friction and wear of tough and brittle zirconia in nitrogen, air, water, hexadecane and hexadecane containing stearic acid. Wear 124 (1988), 133. ╇ 2â•… J. Takadoum, H. H. Bennani, and D. Mairey. The wear characteristics of silicon nitride. J. Eur. Ceram. Soc. 18 (1998), 553. ╇ 3â•… M. F. Wani, B. Prakash, P. K. Das, S. S. Raza, and J. Mukerji. Friction and wear of HPSN bearing materials. Am. Ceram. Soc. Bull. 76(8) (1997), 65. ╇ 4â•… N. Wallbridge, D. Dowson, and E. W. Roberts. The wear characteristics of sliding pairs of high density polycrystalline aluminium oxide under dry and wet conditions, in Proceedings of the International Conference on Wear of Materials, K. Ludema (Ed.). American Society of Mechanical Engineers, New York, 1983, 202. ╇ 5â•… B. Basu. Zirconia-titanium boride composites for tribological applications. PhD thesis, Katholieke Universiteit Leuven, Belgium; March, 2001. ╇ 6â•… K. F. Dufrane. Wear performance of ceramics in ring/cylinder applications. J. Am. Ceram. Soc. 72 (1989), 691. ╇ 7â•… J. J. Habeeb, A. G. Blahey, and W. N. Rogers. Ceramic lubrication. Proceedings of 50 Years of Tribology Conference, London, 1987, 555. ╇ 8â•… W. Morales and D. H. Buckley. Concentrated contact sliding friction and wear behaviour of several ceramics lubricated with a perfluoropolyalkylether at 25°C. Wear 123 (1988), 345–354. ╇ 9â•… K. H. Zum Gahr. Sliding wear of ceramic-ceramic, ceramic-steel and steel-steel pairs in lubricated and unlubricated contact. Wear 133 (1989), 1–22. 10â•… J. D. Oscar Barceinas-Sanchez and W. M. Rainforth. Transmission electron microscopy study of a 3Y-TZP worn under dry and water-lubricated sliding conditions. J. Am. Ceram. Soc. 82(6) (1999), 1483–1491. 11â•… M. Kalin, G. Drazic, S. Novak, and J. Vizintin. Wear mechanisms associated with the lubrication of zirconia ceramics in various aqueous solutions. J. Eur. Ceram. Soc. 26 (2006), 223–232. 12â•… S. Novak and M. Kalin. The effect of pH on the wear of water-lubricated alumina and zirconia ceramics. Tribology Lett. 17 (2004), 727–732. 13â•… S. Novak, G. Drazic, and M. Kalin. Structural changes in ZrO2 ceranics during sliding under various environments. Wear 259 (2005), 562–568. 14â•… M. Kalin, S. Novak, and J. Vizintin. Surface changes as a new concept for boundary lubrication of ceramics with water. J. Phys. D Appl. Phys. 39 (2006), 3138–3149. 15â•… I. Birkby, P. Harrison, and R. Stevens. The effect of surface transformation on the wear behaviour of zirconia TZP ceramics. J. Eur. Ceram. Soc. 5 (1989), 37–45.
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16â•… Y. He, L. Winnubst, A. J. Burggraaf, H. Verweij, P. G. Th. Van der Varst, and B. De With. Grain-size dependence of sliding wear in tetragonal zirconia polycrystals. J. Am. Ceram. Soc. 79(12) (1996), 3090–3096. 17â•… Y. He, L. Winnubst, A. J. Burggraaf, H. Verweij, P. G. Th. Van der Varst, and B. De With. Influence of porosity on friction and wear of tetragonal zirconia polycrystals. J. Am. Ceram. Soc. 80(2) (1997), 377–380. 18â•… S. R. Jansen, A. J. A. Winnubst, Y. J. He, H. Verweij, P. G. T. Van der Varst, and G. De With. Effects of grain size and ceria addition on ageing behavior and tribological properties of Y-TZP ceramics. J. Eur. Ceram. Soc. 18 (1998), 557–563. 19â•… T. E. Fischer, M. P. Anderson, and S. Jahanmir. Influence of fracture toughness on the wear resistance of yttria-doped zirconium oxide. J. Am. Ceram. Soc. 72(2) (1989), 252–257. 20â•… B. Basu, J. Vleugels, and O. Van Der Biest. Microstructure-toughness-wear relationship of tetragonal zirconia ceramics. J. Eur. Ceram. Soc. 24(7) (2004), 2031–2040. 21â•… B. Basu, J. Vleugels, and O. Van Der Biest. Toughness tailoring of yttria-doped zirconia ceramics. Mater. Sci. Eng. A 380 (2004), 215–221. 22â•… B. Basu, J. Vleugels, and O. Van Der Biest. Transformation behaviour of tetragonal zirconia: Role of dopant content and distribution. Mater. Sci. Eng. A 366(2) (2004), 338–347. 23â•… (a) B.-K. Kim, J.-I. Hahn, and K. R. Han. Quantitative phase analysis in tetragonal-rich tetragonal/ monoclinic two phase zirconia by Raman spectroscopy. J. Mater. Sci. Lett. 16 (1997), 669–671. (b) D. R. Clark and F. Adar. Measurement of the crystallographically transformed zone produced by fracture in ceramics containing tetragonal zirconia. J. Am. Ceram. Soc. 65 (1982), 284. 24â•… E. H. Kisi and C. J. Howard. Crystal structures of zirconia phases and their inter-relation. Key Eng. Mater. 153–154 (1998), 1. 25â•… H. S. Kong and M. F. Ashby. Friction-heating maps and their applications. MRS Bull. 16(10) (1991), 41. 26â•… S. W. Lee, S. M. Hsu, and M. C. Shen. Ceramic wear maps: Zirconia. J. Am. Ceram. Soc. 76(8) (1993), 1937. 27â•… A. E. Ginnakopoulos, T. C. Lindley, and S. Suresh. Aspects of equivalence between contact mechanics and fracture mechanics: Theoretical connections and a life-prediction methodology for fretting fatigue. Acta Mater. 46(9) (1998), 2955–2968. 28â•… A. E. Ginnakopoulos and S. Suresh. A three-dimensional analysis of fretting fatigue. Acta Mater. 46(1) (1998), 177–192. 29â•… J. Halling. Introduction to Tribology. Wykeham Publications, London and Winchester, 1976. 30â•… O. Vingsbo. Fretting and Contact Fatigue Studied with the Aid of Fretting Maps, Standardization of Fretting Fatigue Test Methods and Equipment, ASTM STP 1159, M. Helmi Attia and R. B. Waterhouse (Eds.). American Society for Testing and Materials, Philadelphia, 1992, 49–59. 31â•… T. E. Fischer. Wear of ceramics and metals. Proceedings of NSF/AFOSR/ASME Workshop on Tribology Issues and Opportunities in MEMS, Eds. B. Bhushan, Columbus, Ohio, USA, Nov. 9–11, 1997, 157–164. 32â•… H. J. Hertz. Uber die berührung fester elastischer körper. J. Reine Angewandte Mathematik. 92 (1882), 156. 33â•… D. A. Hills and D. Nowell. Mechanics of Fretting Fatigue. Kluwer Academic Publishers, London, 1994. 34â•… G. M. Hamilton and L. E. Goodman. The stress field created by a circular sliding contact. J. Appl. Mech. 33 (1966), 371–376. 35â•… G. M. Hamilton. Explicit equations for the stresses beneath a sliding spherical contact. Proc. Inst. Mech. Eng. 197c (1983), 53–59. 36â•… B. Basu. Toughening of Y-stabilized tetragonal zirconia ceramics. Int. Mater. Rev. 50(4) (2005), 239–256. 37â•… A. G. Evans and R. M. Cannon. Toughening of brittle solids by martensitic transformations. Acta Metall. 34(5) (1986), 761–800. 38â•… J. C. Lambropoulos. Effect of nucleation on transformation toughening. J. Am. Ceram. Soc. 69(7) (1986), 218–222. 39â•… R. H. J. Hannink, P. M. Kelly, and B. C. Muddle. Transformation toughening in zirconia-containing ceramics. J. Am. Ceram. Soc. 83(3) (2000), 461–487.
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40â•… M. Rühle and A. G. Evans. High toughness ceramics and ceramic composites. Prog. Mater. Sci. 33 (1989), 85–167. 41â•… F. F. Lange. Transformation-toughened ZrO2: Correlations between grain size control and composition in the system ZrO2-Y2O3. J. Am. Ceram. Soc. 69(3) (1986), 240–242. 42â•… A. E. Ginnakopoulos, T. C. Lindley, and S. Suresh. Aspects of equivalence between contact mechanics and fracture mechanics: Theoretical connections and a life-prediction methodology for fretting fatigue. Acta Mater. 46(9) (1998), 2955. 43â•… B. Basu, R. G. Vitchev, J. Vleugels, J. P. Celis, and O. Van Der Biest. Influence of humidity on the fretting wear of self-mated tetragonal zirconia ceramics. Acta Materialia 48 (2000), 2461–2471. 44â•… J. Chastain and R. C. King, Jr. (Eds.). Handbook of X-Ray Photoelectron Spectroscopy. Physical Electronics, Eden Prairie, MN, 1992, 261. 45â•… Y. S. Li, P. C. Wang, and K. A. R. Mitchell. XPS investigation of the interactions of hydrogen with thin films of zirconium oxide II. Effects of heating a 26 angstrom thick film after treatment with a hydrogen plasma. Appl. Surf. Sci. 89 (1995), 263. 46â•… J. F. Li, R. Watanbe, B.-P. Zhang, K. Asami, and K. Hashimoto. X-ray photoelectron spectroscopy investigation on the low-temperature degradation of 2â•›mol% Y2O3-ZrO2 ceramics. J. Am. Ceram. Soc. 79 (1996), 3109.
CHAPTER
11
CASE STUDY: SIALON CERAMICS In designing wear-resistant ceramics, it is important to understand the influence of microstructure and material properties on the tribological properties. To illustrate this issue for non-oxide ceramics, the results of the wear tests conducted on sialon ceramics are discussed in this chapter. The simultaneous occurrence of different wear mechanisms, such as tribomechanical wear, assisted by deep abrasive grooves, plowing, and grain pullout as well as tribochemical wear, will be shown in the case of several compositionally tailored sialon ceramics. This chapter also presents the tribological properties of a relatively newer variety of sialon ceramics,that is, S-sialon ceramics, with a particular focus on the possible damage mechanisms.
11.1 INTRODUCTION The combination of high hardness and toughness renders sialon ceramics particularly attractive for tribological applications.1–4 Similar to other ceramics, the friction coefficient and wear rate of the Si3N4-based materials depends on material properties such as microstructure, grain size, grain shape, toughness, hardness, counterbody material, and the experimental conditions such as load, sliding speed, and environment (humidity, atmosphere, etc.). Jones et al.5 investigated the wear behavior of Y-α/β composite sialon ceramics and reported higher wear resistance for the ceramics with high α-sialon content under mild wear conditions. More important, similar ceramic compositions (lower α-sialon content) with elongated microstructure demonstrate better wear characteristics compared with the fine equiaxed α-sialon materials. Reis et al.6 reported that the wear of α-sialon matrix composites is caused by adhesion and microabrasion between the rubbing surfaces. Zhao et al.7 reported the relationship between load, sliding speed, and wear properties of Si3N4 and explained the wear of Si3N4 as a result of adhesion and pullout of the Si3N4 grains. Nakamura et al.8 studied the self-mated wear properties of α-Si3N4 reinforced with in-plane aligned rodlike β-Si3N4 grains and grain pullout was reported as the major wear mechanism. Xie and co-workers9 investigated the effect of microstructure on wear behavior of Ca-doped α-sialon ceramics and reported that elongated-grained Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
167
168â•…
CHAPTER 11â•… Case study: SiAlON Ceramics
microstructure exhibited a reduced wear rate in the severe wear regime, compared with the fine equiaxed-grained microstructure. Such behavior was primarily due to a greater resistance to crack propagation arising from the large elongated-grained microstructure. To assess the tribological performance of Si3N4-based materials, a number of friction and wear studies indicated the role of chemical interaction with the environment.8,10–15 The chemical interaction takes place between Si3N4 and H2O in the mild wear regime.15 Such an interaction can occur in addition to chemical reaction of Si3N4 with a mating counterbody of different composition. It is now known that a transition from tribochemical wear at a low load to tribomechanical wear at a higher load occurs for many Si3N4 systems. A general mechanism for mechanical wear is reported to be grain pullout, aided by intergranular fracture. As far as the microstructural aspect is concerned, both the relative amount of the phase assemblage and the grain morphology influence wear resistance.14 In the work by Jones et al.,13 α-sialon material with coarse elongated morphology is identified as a potential material for various structural components requiring a good combination of mechanical and tribological properties over a large range of load conditions (5–90â•›N). This chapter presents the summary of tribological results to reveal the influence of microstructure, fracture toughness, and hardness on the tribological performance of sialon ceramics, and more extensive discussion is published elsewhere.1,16 The selection of sialon materials with a broad range of hardness and toughness properties was useful in understanding the relationship between wear resistance and mechanical properties. The last part of this chapter also reports the tribological properties of a new type of sialon system, that is, Ba-doped S-sialon ceramics.17–19 The microstructure and mechanical properties of such ceramics are reported elsewhere.20,21 While this chapter also presents the tribological properties of this relatively new sialon ceramic, a detailed discussion can be found elsewhere.22
11.2 MATERIALS AND EXPERIMENTS In the present case study, various sialon ceramics are characterized by a distinctly different combination of microstructure (α:β sialon phase ratios, grain size, z values, intergranular phase), hardness, toughness, elastic modulus, and thermal conductivity, as summarized in Table 11.1. As far as the processing is concerned, the densification was carried out either by gas pressure sintering under 2.2â•›MPa nitrogen gas pressure at different temperatures or by pressureless sintering at 1850°C for 1 hour. The postsintering heat treatment was done on one composition (coded SN5) at 1990°C for 5 hours under 2.2â•›MPa nitrogen gas pressure to facilitate grain growth. Figure 11.1 illustrates that SN5 and SN6 ceramics contain elongated β grains, α phase, and some glassy phase. In developing Ba-doped S-sialon ceramics for tribological study, the powder mixture of BaCO3, Si3N4, AlN, and Al2O3 in appropriate ratio, was hot pressed to full density at 30â•›MPa in boron nitride (BN)-coated graphite dies at 1750°C for 2 hours in N2 atmosphere. Scanning electron microscope (SEM)–energy-dispersive
169
1940°C, 2 hours, 22 bar N2
1940°C, 2 hours, 22 bar N2
1940°C, 2 hours, 22 bar N2
1800°C, 1 hours, 22 bar N2 HT: 1990°C, 5 hours, 22 bar N2 1850°C, 1 hours, 22 bar N2
1850°C, 1 hours, 1 bar N2
1850°C, 1 hours, 22 bar N2
Silzot (d50: 1 µm)
Beta (d50: 1 µm)
%50wtA1╯+╯%50wt B1
Beta (d50: 0.5 µm) Seed: B3 (d50: 3 µm) Beta (d50: 0.5 µm)
Beta (d50: 0.5 µm)
UBE (d50: 0.5 µm)
SN2
SN3
SN4
SN5
SN7
SN8
SN6
1940°C, 2 hours, 22 bar N2
Sintering conditions
UBE (d50: 0.5 µm)
Starting powder
SN1
Sample
3311
3312
3333
3330
3361
3332
3356
3358
Density (kg/m3)
17.59╯±â•¯0.19
12.70╯±â•¯0.20
13.22╯±â•¯0.08
12.41╯±â•¯0.18
14.69╯±â•¯0.16
14.22╯±â•¯0.13
15.76╯±â•¯0.15
16.07╯±â•¯0.06
HV10 (GPa)
6.24╯±â•¯0.23
3.01╯±â•¯0.20
4.98╯±â•¯0.40
5.45╯±â•¯0.17
4.15╯±â•¯0.17
3.76╯±â•¯0.15
5.27╯±â•¯0.13
5.61╯±â•¯0.21
KIc (MPa·m1/2)
293.3
263.9
264.8
230.4
312.7
292.5
302.6
316.7
Elastic modulus (GPa)
12.63
14.32
15.78
23.90
18.66
17.46
16.04
16.20
Thermal conductivity (W/m/K) 67β:33α M:0.89 Z: 0.33 76β:24α M:0.71 Z:0.35 92β:8α M:0.24 Z: 0.42 84β:16α M:0.20 Z: 0.20 100β Amorphous Z:0.10 100β Amorphous Z:0.24 100β Amorphous Z:0.35 71β:29α M:0.65 Z:0.59
XRD (polished surface)
Very fine
Very fine
Fine
CoarseBimodal
Coarse
Coarse
Fine-Bimodal
Fine-Bimodal
Microstructure
TABLE 11.1.â•… Summary of the Composition, Mechanical, and Thermal Properties as well as the Phase Assemblage of the Compositionally Tailored SiAlON Ceramics16
170â•…
CHAPTER 11â•… Case study: SiAlON Ceramics
(a)
(b)
Figure 11.1â•… SEM micrographs representing the microstructure of samples (a) SN6 (wear rate: 2.8╯×╯10−6â•›mm3/Nâ•›m) and (b) SN5 (wear rate: 1.4╯×╯10−5â•›mm3/Nâ•›m). The solid arrows show the presence of glassy phase arising from sintering additives and the dashed arrow shows elongated β-SiAlON grains.16
x-ray spectrometry (EDS) analysis reveals the predominant presence of S-sialon (>90%) with residual glass (2–4%) phase, mostly at the grain boundary triple pockets (Fig. 11.2). Based on transmission electron microscope (TEM)–EDS analysis, the S-phase has an average composition of Ba2Si12−xAlxO2+xN16−x (x╯=╯2╯±â•¯0.2). The hardness and indentation fracture toughness of an S1750 sample, as evaluated by indenting at varying loads of 50–300â•›N using a Vickers Hardness Tester (ModelMVM 50), were determined to vary in the range of 12.5–15.3â•›GPa and 3.7–7.9â•›MPa m1/2, respectively. The mechanical properties of the S-sialon flat and counterbody are presented in Table 11.2.
11.2 Materials and Experiments
â•… 171
10 mm (a) Si
S-phase
Al Ba Ba 1
2
3
4
Ba Ba
5 Energy (keV)
(b)
Figure 11.2â•… SEM image (back-scattered electron [BSE] mode) of the polished (unworn) surface of the investigated Ba-S-sialon ceramic (S1750 sample) (a). Various microstructural phases can be distinguished as S-sialon (grey), darker/black:unreacted Si3N4 and/or β-Si3N4 needles. The bright contrast phase is residual glass phase. Typical EDS analysis of the S-phase (b) is also shown.22
The fretting experiments were performed using a fretting wear tester and a ball-on-flat configuration working on the principle of mode I fretting (linear relative tangential displacement at constant normal load) was used. While fretting tests were carried out on various compositionally modified sialon ceramics against a 9.5-mmdiameter β-sialon ball, S-sialon ceramics were are tested against three counterbodies (Al2O3, sialon, steel). For compositionally tailored sialon ceramics, the test parameters were an 8-N load with a displacement of 100â•›µm at a frequency of 6â•›Hz for 45,000 cycles. On the other hand, S-sialon ceramics were tested at a load of 8â•›N, with displacement of 100 µm at a frequency of 8â•›Hz for 105 cycles.
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TABLE 11.2.â•… Composition and Mechanical Properties of the S-Sialon and Counterbody Materials22
Elastic modulus E (GPa)
Hardness HV (GPa)
Fracture toughness KIC (MPa. m1/2)
Thermal conductivity (W/m/K)
7.8
—
Flat material â•… S-Sialon (BaAlSi5O2N7), S1750 grade
212
15.1
Counterbody ball materials â•… Sialon(Si6-zAlzOzN8-z; z-/0.3,grade TCQ) â•… Alumina (99.7% pure, grade 10) â•… Steel (DIN100Cr 6, construction grade)
310
15.5
4.7
22
300 210
19.0 20
4.0 7.8
39 58
11.3 TRIBOLOGICAL PROPERTIES OF COMPOSITIONALLY TAILORED SIALON VERSUS β-SIALON For various compositionally tailored sialon ceramics, the steady-state coefficient of friction (COF) varies over a window of 0.59–0.64. A lower steady-state COF of 0.59 was recorded for the (84β:16α) SiAlON, while a relatively higher COF of 0.64 was measured with (92β:8α) sialon; some representative two-dimensional (2D) surface profiles across the worn surface were obtained using laser surface profilometer and are shown in Figure 11.3. A critical observation reveals the occurrence of differential wear, that is, variation of depth at different locations on the worn surface. Typically, the maximum wear depth varies between 10 and 20â•›µm. The wear rate data are analyzed in the light of variation in mechanical properties of compositionally tailored sialon ceramics (flat samples in the tribological study). Figure 11.4a,b plots the wear rate versus hardness and toughness, respectively. Observing results in Figure 11.4a, two datasets, contained within dashed and dash-dotted ellipse, follow individually a linear relationship. From the observed correlation between wear rate and hardness, an inverse linear relationship between wear rate and hardness emerges. This essentially implies that the harder the material, the less severe will be the plowing-induced material removal and consequently the less will be the wear rate. Such a clear prediction could, however, not be made, when wear rate is plotted against toughness of sialon flat materials. From Figure 11.4b, we can see that the wear rates of SN2, SN3, SN5, and SN8 ceramics follow a linear relationship with the toughness, while no such relationship is observed for the SN1, SN4, SN6, and SN7 ceramics. Figure 11.4b illustrates that SN1 and SN4 ceramics exhibit very high wear rates, suggesting severe wear. In contrast, SN6 and SN7 ceramics experience low wear rate, indicating ultramild wear. Under a given set of operating parameters and environmental conditions, wear characteristics depend largely on the material microstructure, which in turn is strongly dependent on the
11.3 PROPERTIES OF COMPOSITIONALLY TAILORED SIALON VERSUS β-SIALON
â•… 173
20.00 µm 0.00
–20.00 0.00 mm
(a)
1.50 mm
20.00 µm 0.00
–20.00 0.00 mm
(b)
1.50 mm
10.00 µm 0.00
–10.00 0.00 mm
(c)
1.75 mm
Figure 11.3â•… 2D surface profiles, measured using laser surface profilometer of worn surfaces on various investigated materials (a) SN6, (b) SN7, (c) SN8. Testing conditions: 45,000 cycles; frequency, 6â•›Hz; load, 8â•›N; and displacement, 100â•›µm; counterbody, β-SiAlON ball.16
material composition and mechanical properties. From this discussion, the difference in wear rate cannot be approximately explained on the basis of the difference in toughness properties. Nevertheless, a good number of the data-points in Figure 11.4b clearly suggest an inverse linear relationship with toughness. A critical analysis of x-ray diffraction (XRD) spectra obtained from polished surfaces of various compositionally tailored sialon ceramics provides the quantification of phase assemblage in terms of α:β ratio, β-sialon, and so on. As shown in Figure 11.4, SN2 and SN3 have similar wear rates in spite of being different in terms of α:β phase ratios, fracture toughness, and hardness. Therefore, β-SiAlON content can have an effect on wear resistance. In efforts to probe into the influence of βSiAlON content on wear resistance, the wear rate is plotted against β-SiAlON content in Figure 11.5. A careful analysis of the results in Figure 11.5 suggests that two groups of datasets can be conveniently fitted with a linear plot, implying that β-sialon content should be tailored to optimize wear resistance. The characteristic topographical features of worn surfaces of different SiAlON ceramics are shown in Figures 11.6 and 11.7, which reveal tribolayer formation with
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10 9 SN4
7
SN3
5
SN1
6
4
0
12
SN2
SN8
1
SN6
2
SN7
3 SN5
Wear Rate (∗10–5 mm3/N·m)
8
13
14
15 16 Hardness (GPa) (a)
17
18
10 9 SN4
7 6 SN1
5
3
SN2
SN3
4
0
2.5
3.0
3.5
4.0 4.5 5.0 Toughness (MPa.m1/2) (b)
SN5
SN6
1
SN8
2 SN7
Wear Rate (∗10–5 mm3/N·m)
8
5.5
6.0
6.5
Figure 11.4â•… (a) Variation of measured specific wear rate with hardness (Hv10) of the investigated ceramics. (b) Variation of measured specific wear rate with toughness (KIC) of the various samples. Tribological testing conditions: 45,000 cycles; frequency, 6â•›Hz; load, 8â•›N; and displacement, 100â•›µm; counterbody, β-SiAlON ball.16
abrasive grooves. EDS compositional analysis indicated higher oxygen content of the tribolayer compared with the unworn surface, suggesting the oxidation of sialon to form a thick silica-rich layer. In the case of SN2, SN4, SN5, SN6, SN7, and SN8 ceramics, deep abrasive grooves in the central worn region were observed (see Figs. 11.6 and 11.7). This could be due to the extensive deformation as a result of plowing
11.3 PROPERTIES OF COMPOSITIONALLY TAILORED SIALON VERSUS β-SIALON
â•… 175
SN4
10
6
0
60
70
SN5
SN8
2
80 β-SiAlON Content
90
SN6
SN7
SN2
4
SN3
SN1
Wear Rate (∗10–5 mm3/N·m)
8
100
Figure 11.5â•… Variation of measured specific wear rate with β-SiAlON content of the investigated ceramics. Tribological testing conditions: 45,000 cycles; frequency, 6â•›Hz; load, 8â•›N; and displacement, 100â•›µm; counterbody, β-SiAlON ball.16
action of the counterbody ball.23 While tribochemical layer formation takes place dynamically, the progression of delamination cracks cause subsequent removal of tribolayer (see, e.g., Fig. 11.7b). Tribo-oxidation leads to formation of a tribochemical layer and its fragmentation results in wear debris. Figure 11.6a reveals evidence of microfracture and grain pullouts. In summary, the topographical features of characteristic worn surfaces suggest that tribomechanical wear in combination with tribochemical wear are major wear mechanisms. In particular, the material removal due to abrasive wear results from a combination of three factors: plastic deformation, fracture of the tribolayer, and intergranular fracture. This is in contrast to tribochemical wear, widely observed as the dominant wear mechanism for non-oxide ceramics, such as silicon nitride and silicon carbide. The correlation of wear damage with estimated Hertzian contact pressure can provide an insight into the possible influence of the contact stress. The initial Hertzian contact pressure (maximum) can be determined using the following formula24:
Po = (3 / 2)Pm = (6WE *2 /π 3 R 2 )1/ 3
(11.1)
The initial contact diameter can be estimated using
a = (3WR/ 4 E * )1/ 3,
(11.2)
where Po is the maximum contact pressure, Pm is the mean contact pressure, W is applied load, E* is the effective elastic modulus, and R is the radius of the ball counterbody.
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250 µm
Crack
Grain Pull-out Tribolayer
Tribolayer
200 µm
Cracks 100 µm
50 µm
(a)
(b)
Ploughing
150 µm
Crack
Tribolayer
Grain Pull-out 50 µm
(c)
Figure 11.6â•… SEM micrographs of the worn surface of the SiAlON flat samples (a) SN4, (b) SN5, and (c) SN2. Double-pointed arrow indicates the fretting direction. In the insets, the overview of the fretted damage region is presented. Tribological testing conditions: 45,000 cycles; frequency, 6â•›Hz; load, 8â•›N; and displacement, 100â•›µm; counterbody, β-SiAlON ball.16
The effective elastic modulus E* can be determined as
1/E * = (1 − ν12 ) /E1 + (1 − ν22 ) /E2 ,
(11.3)
where E1 and E2 are the elastic modulus of the mating solids, ν1 and ν2 are Poisson’s ratio of the contacting solids, respectively. Table 11.3 summarizes the estimated values of initial Hertzian contact pressure, contact diameter, and the diameter of as-fretted wear scar for various sialon compositions. It is known that, for materials with higher hardness and higher elastic modulus, the elastic deformation will be less and, consequently, Hertzian contact area will also be less. Therefore, at a given load, the contact stress will be higher, if the contact area is less. Therefore, in spite of higher contact stress, the materials with higher hardness can experience better wear resistance. The results presented in Table 11.3 reveal that the contact region is no longer elastic; that is, the deformation causes the contact to be established over a larger contact area.
11.3 PROPERTIES OF COMPOSITIONALLY TAILORED SIALON VERSUS β-SIALON
â•… 177
Crack 200 µm
200 µm
Tribolayer
Crack
Tribolayer
20 µm
100 µm
(a)
(b)
200 µm
Tribolayer Cracks
50 µm
(c)
Figure 11.7â•… SEM micrographs of the fretted zone of the SiAlON flat samples (a) SN6, (b) SN7, and (c) SN8. Double-pointed arrow indicates the fretting direction. In the insets, the overview of the fretted damage region is presented. Tribological conditions: 45,000 cycles; frequency, 6â•›Hz; load, 8â•›N; and displacement, 100â•›µm; counterbody, β-SiAlON ball.16 TABLE 11.3.â•… Summary of Hertzian Contact Pressure as well as the Initial Hertzian Contact Diameter for Various Compositionally Tailored Sialon Ceramics16
Sample designation
SN1 SN2 SN3 SN4 SN5 SN6 SN7 SN8
Po (Maximum Hertzian contact pressure) (initial), MPa
Pm (Mean Hertzian contact pressure) (initial), MPa
Initial Hertzian contact diameter (µm)
Wear scar diameter on SiAlON flat (µm)
Maximum contact temperature rise (K)
1267 1248 1234 1262 1131 1191 1190 1235
845 832 822 841 754 794 793 823
54.9 55.3 55.6 55 58.1 56.6 56.7 55.6
522 483 550 506 477 533 539 495
365.89 363.17 356.99 349.43 333.14 360.67 379.65 388.58
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It can be reiterated here that the silica-rich tribolayer forms during the fretting tests (see Figs. 11.6 and 11.7) and the feasibility of oxide layer formation also depends on the contact temperature. The contact temperatures can be estimated using Archard’s model.25 According to Archard, the contact temperature depends on the combination of thermal conductivity, COF, contact stress, and sliding speed, and the contact temperature rise (ΔTm) can be estimated using the following formula:
∆Tm = [(µ*(π*Pm )1/ 2 *W 1/ 2 *V ) / (8*k )],
(11.4)
where μ is the COF, Pm is the Hertzian contact stress at yield, W is the applied load, V is the sliding velocity, and k is the thermal conductivity of the sialon ceramic. The Hertzian contact stress at yield can be determined using the following expression:
Pm = [Wy / (π*a 2 )],
(11.5)
where Wy is the load to initiate yield, and a is contact diameter at yield. The contact diameter at yield can be calculated using Equation 11.2 with W replaced by Wy. From the established theory of plasticity criterion at a tribological contact, the load (WY) to initiate yield can be estimated using the following equation26:
WY = 21.17*R 2 *Y *(Y/E * )2,
(11.6)
where Y is the yield strength of the flat sample, E* is the effective elastic modulus, and R is the radius of the ball counterbody. Finally, the maximum contact temperature increase (ΔTmax) can be determined using the following formula:
∆Tmax = 1.67 * ∆Tm.
(11.7)
The values for maximum temperature rise (ΔTmax) presented in Table 11.3 reveal that the maximum temperature rise (ΔTm) is very low and in the range of 60–115°C, which is largely due to low speed (2.4╯×╯10−3â•›m/s) as well as low load (8â•›N). At a tribological contact, the combination of temperature and mechanical stress conditions facilitates the tribochemical reactions at a much lower temperature than would occur under static- or stress-free conditions.27,28 To explain the occurrence of tribochemical wear, two possible oxidation reactions can explain the formation of a silica layer:
Si3 N 4 + 3 O 2 = 3 SiO 2 + 2 N 2 2 SiAlON + 2.5 O2 = 2 SiO2 + Al 2 O3 + N 2
(11.8) (11.9)
The thermodynamic feasibility of reaction 11.8 in the temperature range of 333K–388K, assessed from the free-energy calculations using commercial HSC software29 (also Barin et al.,30 data compilation), indicate that reaction 11.8 is feasible and, therefore, the formation of a silica layer can occur via oxidation of silicon nitride phase present in the sialon ceramics. Therefore, it is possible that unreacted α/β–Si3N4 phase undergoes oxidation to form a SiO2-rich layer, as per reaction 11.8. The observation of a thick SiO2 layer is suggestive of nonprotectiveness, and this is
11.4 Tribological Properties of S-Sialon Ceramic
â•… 179
reflected in cracking or spalling around the tribochemical layer. Also, the grain pullout primarily takes place due to cracking in the residual glass phase.
11.4 TRIBOLOGICAL PROPERTIES OF S-SIALON CERAMIC The results summarized here were obtained when S-sialon ceramic was subjected to fretting tests against steel, sialon, and alumina balls in gross-slip conditions (load 8â•›N, with displacement of 100 µm at a frequency of 8â•›Hz for 105 cycles). All the tribocouples experienced a rapid increase in the COF value during the running-in period, followed by a sudden drop and then the steady-state COF in the range of 0.61–0.63. The wear volume was estimated using Klaffke’s equation and such estimation is based on the wear scar diameter, as summarized in Table 11.4. A plot of wear rate for the ball and the flat is provided in Figure 11.8. Overall, the estimated wear rate varies in range of 10−5 to 10−6â•›mm3/Nâ•›m. In the case of an S-sialon/steel tribocouple, the S-sialon flat experienced more wear by one order of magnitude than the steel counterbody. Interestingly, in the case of S-sialon/β-sialon tribocouple, both mating materials exhibited the same magnitude of wear rate, whereas in case of S-sialon/alumina, the Al2O3 ball suffered lower wear than did the flat. It is reported elsewhere that β-sialon material with moderate hardness (15â•›GPa) and toughness (4.7â•›MPa m1/2) can exhibit wear rate (∼10−6â•›mm3/Nâ•›m) at low load (8â•›N).16 The materials developed via composite approach (α-sialon material reinforced with β-sialon fibers) and self-reinforced (Y, Yb)-α-sialon material also experienced wear rate of the same order of magnitude even at higher loads.14
TABLE 11.4.â•… Maximum and Mean Hertzian Contact Pressure as well as the Initial Hertzian Contact Diameter
Tribocouple
S-Sialon versus alumina S-Sialon versus Sialon S-Sialon versus steel
Maximum Hertzian contact pressure (Initial), MPa
Mean Hertzian contact pressure (Initial), MPa
Initial Hertzian contact diameter, µm
Wear scar diameter on S-Sialon flat, µm
Maximum contact temperature rise23 (°C)
1110.2
740.1
117.3
542.9
232
1058.8
705.9
120.2
568.4
223
941.1
627.4
127.5
766.7
221
The wear scar diameter (measured on SEM images, along the fretting direction), after wear testing at the selected conditions (load: 8â•›N, displacement amplitude: 100 microns, frequency: 8â•›Hz, number of cycles: 100,000) are also provided. The theoretically computed contact temperature rise also provided.22
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CHAPTER 11â•… Case study: SiAlON Ceramics
Flat Ball
1.8×10–5
Wear Rate mm 3/Nm
1.6×10
–5
1.4×10–5 1.2×10–5 1.0×10–5 8.0×10–6 6.0×10–6 4.0×10–6 2.0×10–6 0.0
Counterbody (ball):
Steel
Sialon
Alumina
Figure 11.8â•… Wear rate measured with the respective counterbody, after S-sialon (hot pressed at 1750°C) testing against steel, β-sialon, and alumina under identical gross-slip conditions (load, 8â•›N; displacement, 100â•›µm; frequency, 8â•›Hz; for 100,000 cycles).22
To illustrate typical worn-surface topographical characteristics, representative SEM images of the damage zone on an S-sialon/sialon tribocouple are presented in Figure 11.9. The β-sialon ball exhibited the characteristic “sunburst” pattern, that is, “the central locked region with no visible damage, surrounded by many periodically spaced smooth radial grooves.”31 Also, the finer wear debris particles were entrapped in between the grooves and around the damage zone. Since silicon nitride is harder than the tribofilm, the formation of abrasive grooves can be attributed to direct abrasion during the extrusion of the wear debris. In relation to the central locked portion of the worn ball, the center of the worn flat is characterized by a depression surrounded by the frequently chipped tribolayer. The observation of the relatively smooth worn surface region beneath the tribolayer suggests the absence of any mechanical wear. Broadly, the wear mechanisms were the tribochemical interaction with adhesive wear in case of S-sialon/sialon. The tribochemical layer, once formed, protects the tribosurface from severe asperity–asperity interaction, leading to maintaining steady-state COF of 0.61–0.63, irrespective of counterbody. Based on the Hertzian contact analysis (discussed earlier in this chapter) and contact temperature estimation we can analyze the influence of two important factors, that is, contact stress and flash temperature, on the tribological properties of S-sialon ceramics. The analytically computed values of Pm, Po, and a for three different tribocouples are summarized in Table 11.4. A comparison of Hertzian contact pressure and severity of wear damage (scar diameter) suggests that the tribomechanical wear is not a dominant wear mechanism, since smaller scar diameter is measured with a tribocouple (S-sialon/Al2O3) experiencing the highest magnitude of Hertzian contact stress. This also reconfirms the major role of tribochemical wear mechanisms. In view of the fact that wear is a system-dependent property and considering the total wear loss of the tribosystem (flat╯+╯ball), Ba-doped S-sialon/Al2O3 couple exhibits the highest fretting wear
11.4 Tribological Properties of S-Sialon Ceramic
â•… 181
100 µm
50 µm (a)
400 µm
200 µm (b)
Figure 11.9â•… SEM micrographs (SE imaging mode) of the worn surface on the S-Sialon flat (a) and β-Sialon ball (b). Double pointed arrow indicates the sliding direction. In the insets of (a) and (b), the overview of the overall wear damage region for the respective scar is presented.22
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resistance, while S-sialon/steel exhibits the lowest wear resistance. Such observations, in combination with the data presented in Table 11.4, indicate that despite experiencing higher Hertzian contact stress of more than 1â•›GPa, S-sialon/Al2O3 experiences less material removal, compared with other investigated fretting couples. In assessing the influence of contact temperature, the approximate contact temperature rise (asperity–asperity junction) is estimated utilizing the classical Archard’s model.32 At a tribocontact, the dissipation of frictional heat results in an increase in temperature at the contacting asperities, which can be of high order of magnitude but of short duration.27,32–35 The thermal conductivity of S-sialon is assumed to be 30â•›W/mâ•›K, a value commonly reported for sialon materials. Because of low speed as well as low load (8â•›N), the temperature rise at real contact is just above 200°C and no variation is found dependent on counterbody, as the frictional behavior is similar for S-sialon against steel/Al2O3/sialon. However, the combination of high contact stress and moderate contact temperature rise facilitates tribochemical reactions in the present case.27,34
11.5 CONCLUDING REMARKS Based on the tribological results obtained with the compositionally tailored sialon ceramics, the following inferences can be drawn: a. An inverse relationship between wear rate and hardness is observed for most of the sialons, and the wear rates vary in the range of 10−5–10−6â•›mm3/Nâ•›m. The analysis of 2D depth profiles reveals the maximum depth of around 10–20 µm. b. The tribomechanical wear, assisted by deep abrasive grooves, plowing, grain pullout, and microcracking, is identified as one of the major wear mechanisms for self-mated sialons. The wear of the sialon’s counterbody is dominated by severe abrasion. With respect to the wear resistance of Ba-doped S-sialon ceramic, it is observed that despite having lower hardness (15â•›GPa) compared with the aforementioned sialon ceramics, the S-sialon ceramics exhibit similar wear rate, which can possibly be attributed to two factors: a. The S-sialon material is characterized by elongated grain morphology, and such microstructural features result in considerable crack bridging and deflection. Therefore, brittle-fracture-dominated mechanical wear is reduced for Ba-doped S-sialon ceramic in the severe wear region. In addition, an interlocking S-sialon grain network is beneficial in arresting large crack propagation at tribological contacts. Such an aspect is observed in the case of SiC ceramics.36 b. Independent of mating material (Al2O3/steel/sialon), the wear mechanism of S-sialon is largely dominated by tribochemical wear. Therefore it is the chemistry, that is, the chemical composition of grain boundary glassy phase or the matrix composition that determines the composition and properties of the tribochemical layer.
REFERENCES
Phase assemblage - a/b SiAlON - S-phase SiAlON - Residual glass
â•… 183
Microstructure - Elongated vs. equiaxed grain morphology
Tribological properties of SiAlON ceramic Mechanical properties - Hardness - Fracture toughness
Tribological environment - Humidity - Mating material
Figure 11.10â•… Summary of various factors influencing friction and wear of SiAlON ceramics.
Based on the discussion in this chapter, a summary of various factors influencing the tribological properties of sialon ceramics is provided in Figure 11.10. Like other non-oxide ceramics, the environment has a remarkable influence on the friction and wear of sialon ceramics. Although such results are not shown in this chapter, a number of literature references in this chapter report such influences. Among various factors, the microstructure of the sialon ceramic can be tailored. Depending on the composition and sintering conditions, both the equiaxed and elongated grains in sintered microstructure can be obtained, as is also discussed in the previous chapter. The influence of hardness and fracture toughness are also shown in the case of a number of compositionally tailored sialon ceramics. While the elongated microstructure is useful in obtaining beneficial fracture toughness, such S-sialon ceramics with elongated grain morphology also have good wear resistance comparable with other competing sialon materials.
REFERENCES ╇ 1â•… N. Calis Acikbas, R. Kumar, F. Kara, H. Mandal, and B. Basu. Influence of β-Si3N4 particle size and heat treatment on microstructural evolution of α:β-SiAlON ceramics. J. Eur. Ceram. Soc. 31 (2011), 629–635. ╇ 2â•… I. W. Chen and A. Rosenflanz. A tough Sialon ceramic based on α-Si3N4 with a whisker-like microstructure. Nature 389 (1997), 701–704. ╇ 3â•… E. Y. Sun, P. F. Becher, K. P. Plucknett, C. H. Hsueh, K. B. Alexander, S. B. Waters, K. Hirao, and M. E. Brito. Microstructural design of silicon nitride with improved fracture toughness: II Effects of yttria and alumina additive. J. Am. Ceram. Soc. 81(11) (1998), 2831–2840. ╇ 4â•… M. Zenotchkine, R. Shuba, and I.-W. Chen. Effect of seeding on the microstructure and mechanical properties of alpha-SiAlON: III comparison of modifying cations. J. Am. Ceram. Soc. 86 (2003), 1168–1175. ╇ 5â•… M. I. Jones, H. Kiyoshi, H. Hideki, Y. Yukihiko, and K. Shuzo. Wear properties of Y-α/β composite sialon ceramics. J. Eur. Ceram. Soc. 23 (2003), 1743–1750. ╇ 6â•… P. Reis, J. P. Davim, X. Xu, and J. M. F. Ferreira. Tribological behavior of colloidally processed SiAlON ceramics sliding against steel under dry conditions. Tribology Lett. 18(3) (2005), 295–301. ╇ 7â•… X. Zhao, J. Liu, B. Zhu, H. Miao, and Z. Luo. Wear behavior of Si3N4 ceramic cutting tool material against stainless steel in dry and water-lubricated conditions. Ceram. Int. 25 (1999), 309–315.
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╇ 8â•… M. Nakamura, K. Hirao, Y. Yukihiko, and S. Kanzaki. Wear behavior of α-Si3N4 ceramics reinforced by rod-like β-Si3N4 grains. Wear 254 (2003), 94–102. ╇ 9â•… Z. H. Xie, M. Hoffman, R. J. Moon, P. R. Munroe, and Y. B. Cheng. Sliding wear of calcium αSiAlON ceramics. Wear 260 (2006), 387–400. 10â•… S. Novak, G. Drazic, Z. Samardija, M. Kalin, and J. Vizintin. Wear of silicon nitride ceramics under fretting conditions. Mater. Sci. Eng. A215 (1996), 125–133. 11â•… M. Kalin, J. Vizintin, S. Novak, and G. Drazic. Wear mechanisms in oil-lubricated and dry fretting of silicon nitride against bearing steel contacts. Wear 210 (1997), 27–38. 12â•… M. I. Jones, H. Hyugab, K. Hirao, and Y. Yamauchi. Wear properties under dry sliding of Sialons with in situ reinforced microstructures. J. Eur. Ceram. Soc. 24 (2004), 3581–3589. 13â•… M. I. Jones, K. Hirao, H. Hyuga, Y. Yamauchi, Z. Shen, and M. Nygren. Wear properties of selfreinforced Sialon ceramics produced by spark plasma sintering. Wear 257 (2004), 292–296. 14â•… B. Prakash, S. Bandyopadhyaya, and J. Mukerji. Friction coefficient of Sialon composite against a steel and dense Silicon Nitride tribopair. J. Am. Ceram. Soc. 82 (1999), 2255–2256. 15â•… B. Basu, J. Vleugels, M. Kalin, and O. Van Der Biest. Friction and wear behavior of Sialon ceramics under fretting contacts. Mater. Eng. A359 (2003), 228–236. 16â•… R. Kumar, N. Ackibas, F. Kara, H. Mandal, and B. Basu. Microstructure-mechanical property-wear resistance relationship for Sialon ceramics. Metall. Mater. Trans. A 40(10) (2009), 2319–2332. 17â•… C. J. Hwang, D. W. Susintzky, and D. R. Beaman. Preparation of multication α-Sialon containing strontium. J. Am. Ceram. Soc. 78(3) (1995), 588–592. 18â•… S. Esmaeilzadeh, J. Grins, Z. Shen, M. Eden, and M. Thiaux. Study of Sialon S-phases M2AlxSi12−xN16−xO2+x, M: Ba and Ba0.9Eu0.1, by x-ray single crystal diffraction, x-ray powder diffraction, and solid-state nuclear magnetic resonance. Chem. Mater. 16(11) (2004), 2113–2120. 19â•… Z. Shen, J. Grins, S. Esmaeilzadeh, and H. Ehrenberg. Preparation and crystal structure of a new Sr containing Sialon phase Sr2AlxSi12−xN16−xO2−x (x∼2). J. Mater. Chem. 9 (1999), 1019–1022. 20â•… B. Basu, M. H. Lewis, M. E. Smith, M. Bunyard, and T. Kemp. Microstructure development and properties of novel Ba-doped S-phase Sialon ceramics. J. Eur. Ceram. Soc. 26 (2006), 3919–3924. 21â•… B. Basu, Manisha, and N. K. Mukhopadhyay. Understanding the mechanical properties of hot pressed Ba-doped S-Phase sialon Ceramics. J. Eur. Ceram. Soc. 29 (2009), 801–811. 22â•… Manisha and B. Basu. Tribological Properties of a hot pressed Ba-doped S-Phase Sialon Ceramic. J. Am. Ceram. Soc. 90(6) (2007), 1858–1865. 23â•… S. Venkatachalam and S. Y. Liang. Effects of ploughing forces and friction coefficient in microscale machining. J. Manuf. Sci. Eng. 129(2) (2007), 274–280. 24â•… B. Bhushan. Principles and Applications of Tribology. A Wiley-Interscience Publication, John Wiley & Sons, New York, 1999. 25â•… J. F. Archard. The temperature of rubbing surfaces. Wear 2(6) (1958), 438–455. 26â•… Y. G. Ko, D. H. Shin, K. T. Park, and C. S. Lee. An analysis of the strain hardening behavior of ultra-fine grain pure titanium. Scr. Mater. 54 (2006), 1785–1789. 27â•… M. Kalin and J. Vizintin. High temperature phase transformations under fretting conditions. Wear 249 (2001), 172–181. 28â•… Z. R. Zhou, E. Sauger, J. J. Liu, and L. Vincent. Nucleation and early growth of tribologically transformed structure (TrS) induced by fretting. Wear 212 (1997), 50–58. 29â•… HSC Chemistry® Version 5.1. Available at www.outotec.com. 30â•… I. Barin, O. Knacke, and O. Kubaschewski. Thermodynamic Properties of Inorganic Substances. Springer-Verlag, Berlin, 1973, Supplement 1977. 31â•… S. Jahanmir. Advanced Ceramics in Tribological Applications. Marcel Dekker, New York, 1994, 3–12. 32â•… J. F. Archard. The temperature of rubbing surfaces. Wear 2 (1958–1959), 438–455. 33â•… D. Guha and S. K. Choudhri. The effect of surface roughness on the temperature at the contact between sliding bodies. Wear 197 (1996), 63–73. 34â•… I. V. Kragelsky. Friction Wear Lubrication, Tribology Handbook. Mir Publisher, Moscow, 1981. 35â•… M. Kalin and J. Vižintin. A tentative explanation for the tribochemical effects in fretting wear. Wear 250(1/12) (2001), 681–689. 36â•… O. B. Lopez, A. L. Ortiz, F. Guiberteau, and N. P. Padture. Sliding-wear-resistant liquid-phasesintered SiC processed using α-SiC starting powder. J. Am. Ceram. Soc. 90(2) (2007), 541–545.
CHAPTER
12
CASE STUDY: MAX PHASE—TI3SIC2 Among various materials, ceramics are popular for two contrasting properties: high hardness and poor fracture toughness (i.e., lack of ductility). In this chapter, the friction and wear properties of Ti3SiC2 will be discussed to illustrate the friction and wear of a ceramic having low hardness (slightly harder than fully hardened steel) and good toughness. The microscopic analysis of worn surfaces will be presented along with Raman spectroscopy results to discuss the role of tribochemical wear. An important observation is that “metal-like plasticity” behavior is observed at higher load (>6â•›N) and a probable explanation for the transition in friction and wear with load is also presented.
12.1 BACKGROUND In the last two decades, ternary carbides1 have attracted wider attention in the materials community. Since the discovery of a vast number of ternary carbides and nitrides by Nowotny and Jeitschko,1 Ti3SiC2 has been investigated2–12 extensively by several authors. Barsoum2 identified that these phases represent a new class of solids that can be described as thermodynamically stable laminates. Interestingly, ternary carbides exhibit both metallic and ceramic features.2 This class of materials is now known as MAX phases, because of their chemistry: ternary layered hexagonal carbides and nitrides, with the general formula Mn+1AXn, where n╯=╯1,…, 3, M is an early transition metal, A belongs to the III–VIA group of the periodic table, and X is C and/or N. Among MAX phases, Ti3SiC2 has some outstanding properties, including excellent resistance to oxidation up to 1400°C, high thermal shock resistance, high Young’s modulus (325â•›GPa), relatively low hardness (4–5â•›GPa), high fracture toughness (7–9â•›MPa m1/2), and good machinability with conventional tools.6,7 Like graphite, Ti3SiC2 has a hexagonal structure, as shown in Figure 12.1; its crystal structure can be described as a planar stacking sequence along the c-axis, consisting of double layers of Ti–C edge-sharing octahedra, sandwiched between sheets of square-planar coordinated Si atoms.1 The electrical (2–15╯×╯106â•›(Ω·m)−1) and thermal conductivity Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
185
186â•…
CHAPTER 12â•… CASE STUDY: MAX PHASE—TI3SIC2
a3 a2
a1 c Si
Ti
Ti(1) Ti(2)
17.67 Å C
Ti(3)
Si C
3.0665 Å
Figure 12.1â•… Unit cell of Ti3SiC2.1
(20–50â•›W/m·K) of MAX phases can exceed those of Ti metal. Being readily machinable, Ti3SiC2 has higher elastic (E) modulus (∼320â•›GPa), that is, it is three times stiffer than Ti metal; however, it has the same density as Ti (i.e., ∼4.5â•›g/cm3).8 Ti3SiC2 is also known as a damage-tolerant material: at 1300°C, Ti3SiC2 exhibits “yield” points at 100 and 500â•›MPa in flexure and compression loading, respectively.9 The large-grained, polycrystalline Ti3SiC2, loaded in compression at room temperature (RT), can deform plastically by a combination of shear and kink-band formation.3,4 From the perspective of tribological applications, the frictional and wear resistance of Ti3SiC2 were investigated by several researchers. Myhra et al.,10 using lateral force microscopy, recorded a low kinetic friction coefficient (μ) of 0.002– 0.005 at 1000–20,000â•›nN normal force on basal planes of Ti3SiC2. They also measured a low coefficient of friction (COF) of 0.12 for polycrystalline Ti3SiC2 against stainless steel under a load of 0.15–0.9â•›N. In contrast, Raghy et al. reported a high COF of 0.8 for a Ti3SiC2/steel tribocouple at 5-N load.11 Importantly, the frictional characteristics, measured using a pin-on-disk configuration, were found to be independent of the microstructure (grain size of Ti3SiC2 varied between 5 to 100╛µm). Zhang et al.12 investigated the friction and wear behavior of self-mated Ti3SiC2 and a Ti3SiC2/diamond pair using a pin-on-flat tribometer. While measurements of the COF of the former varied in the range 1.16–1.43, those of the latter were below 0.1 for varying loads of 0.49–9.8â•›N. The low COF of Ti3SiC2 against diamond was due to the formation of a lubricating film on the Ti3SiC2 tribosurface. Tang et al.13 reported excellent wear resistance of a laser-melted ternary metal silicide, Cr13Ni5Si2 alloy, under sliding wear conditions for loads of 98–196â•›N. In more recent work, the
12.1 Background
â•… 187
TiC
5 µm
Figure 12.2â•… Optical microstructure of Ti3SiC2, etched in a 1:1:1 by volume HF:HNO3:H2O solution. Note the presence of large elongated Ti3SiC2 grains and the small amount of TiC (bright contrast).14
TABLE 12.1.â•… Density and Mechanical Properties of the Tribocouple14
Specimen
Flat Ball (Counterbody)
Material
Density ρ (gm/cc)
Hardness (GPa)
E (GPa)
Fracture toughness KIc (MPa m1/2)
Ti3SiC2 (T1) Steel (SAE 52100 Grade)
4.5 7.8
4.7╯±â•¯0.3 ∼7.0
316 210
8.9╯±â•¯0.1 —
fretting wear properties of hot pressed Ti3SiC2 have been reported and this chapter summarizes important results of such studies.14,15 To obtain dense Ti3SiC2, the green compacts containing stoichiometric amounts of TiC0.67 and Si are hot pressed in a graphite die at 1420°C for 90 minutes in an Ar atmosphere. The Ti3SiC2 microstructure is characterized by platelike elongated grains (∼50–200╛µm) with an aspect ratio of about 8 (see Fig. 12.2). The hot pressed polycrystalline Ti3SiC2 exhibited a combination of moderate hardness (∼5â•›GPa), lower than hardened steel (∼7â•›GPa), and good fracture toughness (∼9â•›MPaâ•›m1/2), better than many structural ceramics (2–5â•›MPaâ•›m1/2), as provided in Table 12.1. One of the important aspects of the mechanical properties of this ceramic is its resistancecurve (R-curve) behavior, that is, its enhanced crack growth resistance with crack length.16 Since tribomechanical wear in ceramics is dominated by cracking-induced brittle fracture, such R-curve behavior can be a useful property that can be utilized for tribological applications.
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CHAPTER 12â•… CASE STUDY: MAX PHASE—TI3SIC2
Wear Rate (×10−5 mm3/N·m)
36
0.6
27
0.5 18
0.4
9
Wear Rate
Coefficient of Friction (µ)
0.7
45
Coefficient of Friction
0.3
0
0
2
4 6 Load (N)
8
10
Figure 12.3â•… The measured wear rate and coefficient of friction (μ), obtained for Ti3SiC2 sliding against steel under varying loads for 100,000 cycles.14
12.2 FRICTIONAL BEHAVIOR In Figure 12.3, the steady-state COF of Ti3SiC2 fretted against steel are plotted. At high load (8â•›N), the steady-state COF is measured to be around 0.5. Also, the COF of Ti3SiC2/steel, as measured by pin-on-disk tribometer, lies around 0.83 at 5-N load.11 An interesting observation is that Ti3SiC2, despite having a characteristic chainlike structure, exhibits higher COF (0.5–0.6) compared with other layered structures. In lamellar solids such as graphite, the crystal structure is characterized by layers or sheets, within which the bonding between atoms is covalent and strong. These layers are held together by weak van der Waals-type bonding. For example in MoS2, distance between the planes of Mo and S atoms is 1.58╛Šand adjacent planes of S atoms are 3.01╛Šapart. When loaded with a relatively small force, displacement of the layers by easy slippage occurs, leading to low COF. For example, graphite and MoS2 exhibit low COF (∼0.2) at RT.17 Similarly, h-BN has low COF (0.2), which is maintained up to 850°C. As shown in Figure 12.3, Ti3SiC2 has higher COF and this indicates that a similar lubrication mechanism does not operate. This can be attributed to the inherent bond structure, as shown in Figure 12.1. It has been reported in the literature18 that the interatomic bond length in Ti(1)–Si is around 2.69â•›Å, which is lower than that in graphite (3.40â•›Å). Because of the smaller bond length, the bond strength is higher in Ti3SiC2 than in other lamellar solids (MoS2, graphite), and this results in difficulty of slippage of the Ti-C-Ti-C-Ti-Si network.
12.3 WEAR RESISTANCE AND WEAR MECHANISMS Wear volume, computed based on observations using a laser surface profilometer, is normalized with respect to normal load and total sliding distance to obtain the
â•… 189
12.3 Wear Resistance and Wear Mechanisms
1N
4N
(a)
10 µm
(b)
10 µm
6N 10 mm
10 µm (c)
Figure 12.4â•… SEM images showing deeper abrasive scratches of groove width around 2–3╛µm (a), spalling of nonprotective tribochemical layer due to propagation of cracks (b,c) on Ti3SiC2 worn surface after testing against bearing steel for 100,000 cycles under varying loads, as mentioned against individual micrograph.14
specific wear rate, and the fretting wear rate of Ti3SiC2 against steel is plotted against load in Figure 12.3. While the wear rate varies over the same order of magnitude (i.e., 10−5â•›mm3/Nâ•›m), the wear rate increases with load from 1 to 2â•›N, but does not vary much (20–25╯×╯10−5â•›mm3/Nâ•›m) for a fretting load of 2–6â•›N. A significant increase in wear rate is observed at 8- and 10-N loads, with a maximum of ∼37╯×╯10−5â•›mm3/Nâ•›m being measured at 8-N load. It has been widely reported that the measured wear rate of hard ceramics such as Al2O3, SiC, and Si3N4 is on the order of 10−6â•›mm3/Nâ•›m (the lowest down to 10−9â•›mm3/Nâ•›m).19–21 The higher wear rate measured with Ti3SiC2 is presumably due to lower hardness (∼4–5â•›GPa). Furthermore, a much higher wear rate (10−3â•›mm3/Nâ•›m) has been recorded for Ti3SiC2 against steel using pin-on-disk tribometer.11 Detailed topographical features of the Ti3SiC2 worn surfaces are presented in Figure 12.4. The significant formation of a tribochemical layer is noticed at 4-N load, as revealed in Figure 12.4b. However, such a tribochemical layer is largely nonprotective and spalls off due to propagation of cracks. Also, at 6-N load, a plastically deformed layer can be observed (Fig. 12.4c). Energy-dispersive x-ray spectrometry (EDS) analysis reveals the transfer of a smaller amount from the steel ball.
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CHAPTER 12â•… CASE STUDY: MAX PHASE—TI3SIC2
8N
(a)
(b)
(c)
Figure 12.5â•… SEM images showing extensive plastic deformation (a,c), cracks propagating along the edges of elongated grains (c) on the fretted surface of Ti3SiC2 after testing against steel, for 100,000 cycles at high loads of 8 and 10â•›N, as indicated in the individual SEM images. The details of the deformed tribolayer are also shown in the insets (a,c). The pointed arrow indicates the sliding direction.14
A clear change in worn-surface topography is noticed at higher loads (8 and 10â•›N; see Fig. 12.5). The evidence of extensive plastic deformation as well as other topographical features is more like that of the worn surface on Ti3SiC2 after fretting at 8-N load, resembling metallic materials. The fracture of elongated Ti3SiC2 grains, as well as the propagation of cracks along the edges of elongated grains, is observed. A closer observation of Figure 12.5b reveals that such a fracture process leads to formation of coarser debris particles with sizes of 5–30╛µm. At 10-N load, the contribution of plastic deformation is equally significant and cracks are also observed (Fig. 12.5c).
12.4 RAMAN SPECTROSCOPY AND ATOMIC FORCE MICROSCOPY ANALYSIS Summarizing the preceding discussion of fretting results, it is evident that a distinct transition in the friction and wear of Ti3SiC2 occurs at loads between 6 and 8â•›N. Also,
â•… 191
200
2
Ti O
Intensity (Arb. Units)
Fe
2O 3
+
Fe
2O 3
Ti O
2
+
Si O
2
12.4 Raman Spectroscopy and Atomic Force Microscopy Analysis
8N
6N
300
400
500
600
Raman Shift (cm-1)
Figure 12.6â•… Raman spectra obtained from the worn surface on Ti3SiC2, worn against steel for 100,000 cycles for different loads, as indicated against the individual spectra.14
the plastic-deformation-dominated damage is more severe at 8-N load compared with material removal induced by abrasion and tribochemical reactions at 6-N load. Because of this important observation, the experimental results that were used to characterize the chemistry of the tribolayer and to quantify worn-surface topography after fretting at 6- and 8-N loads were obtained using Raman spectroscopy (Fig. 12.6) and atomic force microscopy (AFM). After fretting at 6-N load, analysis of the Raman spectra, when compared with literature reports,22,23 reveals evidence of the formation of Fe2O3, TiO2, and SiO2. The residual TiC in hot pressed Ti3SiC2 is presumably oxidized to TiO2 during fretting, and Ti3SiC2 can also be oxidized to form TiO2 and SiO2. Therefore, the tribochemical oxidation taking place during fretting can be described by the following reactions:
2 Fe + (3 / 2)O2 = Fe 2 O3 TiC + 2O2 = TiO2 + CO2 (g) Ti3SiC2 + 6O2 = 3TiO2 + SiO2 + 2CO2 (g)
(12.1) (12.2) (12.3)
A literature report24 confirms that reaction 12.3 initiates at 900°C during oxidation in air, which ultimately leads to the formation of distinct rutile and silica layers. Prior to discussion of the AFM results, the concept of bearing area and various important associated parameters need to be defined. The bearing area curve provides the material or bearing ratio, which is the ratio of the material-filled length to the evaluation length at the given profile section level.25 The bearing area curve, measured using AFM, is plotted in Figure 12.7 and the corresponding roughness parameters are summarized in Table 12.2. A closer observation of the roughness data presented in Table 12.2 reveals that all the characteristic roughness parameters—Sa, Sq, Sy, Spk, Sk, and Svk, except Ssc—increase about 5 to 8 times after fretting at 6- and
192â•…
CHAPTER 12â•… CASE STUDY: MAX PHASE—TI3SIC2
100
Unworn Surface of Ti3SiC2 Worn Surface of Ti3SiC2 (6 N) Worn Surface of Ti3SiC2 (8 N)
Bearing Ratio (%)
80
60
40
20
0 0
200
400
600 800 Height (nm)
1000
1200
Figure 12.7â•… Bearing area of unworn and worn Ti3SiC2 surface, as analyzed by AFM.14 TABLE 12.2.â•… Roughness Parameters, as Measured Using an AFM of Unworn and Worn Surfaces after Sliding at 6 and 8â•›N14
Specimen
Unworn surface
Ratio of values measured (Worn surface at 6â•›Nâ•›:â•›Unworn surface)
Ratio of values measured (Worn surface at 8â•›Nâ•›:â•›Unworn surface)
Sa (Roughness average) Sq (Root mean square) Ssk (Surface skewness) Sku (Surface kurtosis) Sy (Peak–peak) Ssc (Mean summit curvature) Spk (Reduced peak height) Sk (Core roughness depth) Svk (Reduced valley height)
28â•›nm 32â•›nm −0.0169 0.446 153â•›nm 18.94╯×╯105â•›nm 16.74â•›nm 98â•›nm 21â•›nm
4.96 5.19 5.30 5.00 5.10 0.20 4.86 5.10 5.14
6.86 7.31 5.60 5.34 7.59 0.05 8.60 6.54 9.43
8-N load, when compared with those of unworn surface. The average peak-to-peak distance (Sy) also increases with normal load, indicating an increase in severity of wear due to deformation or removal of surface asperities. Also, the core roughness data (Sk), an alternative measure of surface roughness for Sa and Sku, are measured to increase by 5 and 6 times as load increases from 1â•›N to 6 and 8â•›N, respectively. The reduced peak height (Spk) of the worn surface increases to 5 and 8 times with similar variation in fretting load. The reduced valley height (Svk) also increases with load, implying the entrapment of wear debris particles at the sliding interface. The
â•… 193
12.5 Transition in Wear Mechanisms
(a)
(b)
(c)
(d)
Figure 12.8â•… 2D-AFM image of worn surface topography of Ti3SiC2 when tested against steel ball: (a) 6â•›N and (c) 8â•›N. The line profile of Ti3SiC2 worn surface: (b) 6â•›N and (d) 8â•›N.14
characteristics of two-dimensional AFM images of the selected sections are presented in Figure12.8, which reveals severe abrasion, interestingly having grooves of different depths. Also, the topographical features of Ti3SiC2 resemble those of the wear of conventional metallic materials. Therefore, the combined effect of severe abrasion and deformation-induced damage results in material removal by layers. The depth of the abrasive grooves is measured to be around 800 and 1200â•›nm for the fretting loads of 6 and 8â•›N, respectively (Fig. 12.8b,d).
12.5 TRANSITION IN WEAR MECHANISMS In the following, the mechanism of material removal for Ti3SiC2 is summarized. Three major mechanisms contributing to friction and wear of Ti3SiC2 are (1) abrasion, (2) tribochemical layer formation, and (3) plastic deformation. The steady-state COF of a Ti3SiC2/steel couple increases from 0.55 to 0.6 as load is increased from 1 to 6â•›N, respectively. However, a decrease in COF from 0.62 to 0.5 is measured when load is increased from 6 to 8â•›N, and COF remains constant (∼0.5) at 10-N load (Fig. 12.3). At higher load (>6â•›N), the formation of tribochemical
194â•…
CHAPTER 12â•… CASE STUDY: MAX PHASE—TI3SIC2
reaction products and wear debris takes place to a larger extent. These debris particles (third body), while being entrapped between Ti3SiC2 and steel (first two bodies), tend to roll during sliding to decrease the friction. Hence, a transition from two-body to three-body abrasion occurs at higher load (>6â•›N), as also observed in Figure 12.3. Concerning wear mechanisms, severe abrasion as well as tribolayer formation is observed at low loads (1–6â•›N). The observation of severe abrasion on Ti3SiC2, even at the lowest load (1â•›N), can be rationalized from the difference in hardness between mating counterfaces. At intermediate load (4 and 6â•›N), the cracking of the tribolayer increases the wear rate of Ti3SiC2. A change in wear mechanism is critically recorded at high load (8–10â•›N). Although the tribochemical wear remains an active wear mechanism at loads greater than 6â•›N, plastic deformation is identified as a significant wear mechanism. The deformation of Ti3SiC2, as explained in the existing literature,4,5,9 can be used to explain the deformation-induced wear of Ti3SiC2. In Ti3SiC2, two adjacent covalent bond chains of Ti-C-Ti-C-Ti-Si form a chain couple and the chains are bonded together with strong metallic Ti layers (see Fig. 12.1). Therefore, the deformation of Ti3SiC2 on worn surfaces appears to be due to metallic-like bonds. Additionally, the polar character of the directional bonding indicates ionic bonding in Ti–C and Ti–Si interactions and such anisotropy of metallic–covalent–ionic bonding appears to be responsible for the Ti3SiC2 plasticity. It is worthwhile to note here that though the elastic modulus (E) of polycrystalline Ti3SiC2 is quite high (316â•›GPa), the ratio of modulus to hardness (∼63) lies in the range of ductile materials.9 According to Barsoum and co-workers,4,9 the plastic deformation of Ti3SiC2 could be due to delamination and kink-band formation at RT, if grains are oriented. Also, the observation that plastic deformation occurs only at higher load (8N) during fretting of Ti3SiC2 indicates that the plasticity of the chainlike structure requires a critical contact pressure. The Hertzian contact pressure at 8-N load is calculated to be around 800â•›MPa. Also, the considerable fraction of frictional energy, dissipated as heat energy, is partitioned between two mating solids. Therefore, the combined effect of high contact pressure (at load╯≥╯8N) and high contact temperature results in observed plasticity on the worn surfaces of Ti3SiC2.
12.6 SUMMARY The steady-state COF and wear rate of Ti3SiC2 are summarized along with the literature data in Table 12.3. As expected, the friction coefficient varies over a wide range depending on the test configuration and operating parameters (load). The important point to be noted is that Ti3SiC2 ceramic exhibits high wear rate (10−5â•›mm3/Nâ•›m) compared with many other structural ceramics (10−6 to 10−7â•›mm3/Nâ•›m), as discussed in this book. This can be attributed to the relatively low hardness of Ti3SiC2. This indicates that it is the combination of hardness, E-modulus, and toughness that needs to be optimized for better wear resistance in engineering applications. From a materials development perspective, the results presented in this chapter indicate that Ti3SiC2-based composites with higher hardness and modulus without compromising much on toughness need to be pursued in the context of tribological applications.
â•… 195
REFERENCES
TABLE 12.3.â•… Summary of Some of the Available Friction and Wear Data Obtained with Ti3SiC2 under Varying Tribological Test Conditions15
Specimen
Counterbody
Contact
Load
COF
Wear rate (10−5â•›mm3/Nâ•›m)
Ti3SiC2
Steel
Ball-on-disk
— Steel Ti3SiC2
LFM Pin-on-disk Pin-on-flat
Pin-on-flat
0.55╯±â•¯0.05 0.56╯±â•¯0.05 0.59╯±â•¯0.05 0.62╯±â•¯0.05 0.50╯±â•¯0.05 0.50╯±â•¯0.05 0.002a 0.83 1.43 1.29 1.34 1.23 1.16 0.09 0.08 0.08 0.07 0.06
11╯±â•¯1 23╯±â•¯2 23╯±â•¯4 23╯±â•¯6 37╯±â•¯4 37╯±â•¯5 — 134–425 —
Diamond
1â•›N 2â•›N 4â•›N 6â•›N 8â•›N 10â•›N 25â•›nN 5â•›N 0.98â•›N 1.96â•›N 2.94â•›N 4.9â•›N 9.8â•›N 0.98â•›N 1.96â•›N 2.94â•›N 4.9â•›N 9.8â•›N
Ti3SiC2 Ti3SiC2 Ti3SiC2
a
0.002, along the basal plane of Ti3SiC2. LFM, lateral force microscopy.
REFERENCES ╇ 1â•… W. Jeitschko and H. Nowotny. Die Kristallstruktur von Ti3SiC2—Ein NeuerKomplexcarbid-Typ. Monatsh. fur Chem. 98 (1967), 329–337. ╇ 2â•… M. W. Barsoum. The MN+1AXN phases: A new class of solids Thermodynamically stable nanolaminates. Solid State Chem. 28(1) (2000), 201–281. ╇ 3â•… A. Murugaiah, A. Souchet, T. El-Raghy, M. Radovic, M. Sundberg, and M. W. Barsoum. Tape casting, pressureless sintering, and grain growth in Ti3SiC2 compacts. J. Am. Ceram. Soc. 87(4) (2004), 550. ╇ 4â•… M. W. Barsoum and T. E. Raghy. Room temperature ductile carbides. Metall. Mater. Trans. 30A (1999), 363–369. ╇ 5â•… Y. Zhou and Z. Sun. Electronic structure and bonding properties in layered ternary carbide Ti3SiC2. J. Phys. Condens. Matter 12(28) (2000), L457–L462. ╇ 6â•… M. W. Barsoum and T. E. Raghy. Synthesis and characterization of a remarkable ceramic: Ti3SiC2. J. Am. Ceram. Soc. 79(7) (1996), 1953–1956. ╇ 7â•… M. W. Barsoum, T. E. Raghy, C. Rawn, W. Porter, H. Wang, A. Payzant, and C. Hubbard. Thermal properties of Ti3SiC2. J. Phys. Chem. Solids 60 (1999), 429–439. ╇ 8â•… T. E. Raghy, A. Zavaliangos, M. W. Barsoum, and S. Kalidindi. Damage mechanisms around hardness indentations in Ti3SiC2. J. Am. Ceram. Soc. 80 (1997), 513. ╇ 9â•… T. E. Raghy, M. W. Barsoum, A. Zavaliangos, and S. Kalidindi. Processing and mechanical properties of Ti3SiC2, Part II: Mechanical properties. J. Am. Ceram. Soc. 82 (1999), 2855–2859.
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10â•… S. Myhra and J. W. B. Summers. Kisi EH. Ti3SiC2—A layered ceramic exhibiting ultra-low friction. Mater. Lett. 39 (1999), 6–11. 11â•… T. E. Raghy, P. Blau, and M. W. Barsoum. Effect of grain size on friction and wear behavior of Ti3SiC2. Wear 238 (2000), 125–130. 12â•… Y. Zhang, G. P. Ding, Y. C. Zhou, and B. C. Cai. Ti3SiC2—A self-lubricating ceramic. Mater. Lett. 55 (2002), 285–289. 13â•… H. B. Tang, Y. L. Fang, and H. M. Wang. Microstructure and dry sliding wear resistance of a Cr13Ni5Si2 ternary metal silicide alloy. Acta Mater. 52(7) (2004), 1773. 14â•… D. Sarkar, B. V. Manoj Kumar, and B. Basu. Understanding the fretting wear of Ti3SiC2. J. Eur. Ceram. Soc. 26 (2006), 2441–2452. 15â•… D. Sarkar, S. J. Cho, M. C. Chu, S. S. Hwang, S. W. Park, and B. Basu. Tribological properties of Ti3SiC2. J. Am. Ceram. Soc. 88(11) (2005), 3245–3248. 16â•… D. Sarkar, B. Basu, M. C. Chu, and S. J. Cho. R-curve behavior of Ti3SiC2. Ceram. Int. 33 (2007), 789–793. 17â•… R. Deacon. Lubrication by lamellar solids. Proc. R. Soc. 243A (1957), 464. 18â•… N. I. Medvedeva, D. L. Novikov, A. L. Ivanovsky, M. V. Kuznetsov, and A. J. Freeman. Electronic properties of Ti3SiC2-based solid solutions. Phys. Rev. B 58 (1998), 16042–16050. 19â•… M. Chen, K. Kato, and K. Adachi. Friction and wear of selfmated SiC and Si3N4 sliding in water. Wear 250 (2001), 246–255. 20â•… M. C. Jeng and L. Y. Yan. Environmental effects on wear behavior of Al2O3. Wear 161 (1993), 11–16. 21â•… S. M. Hsu and M. C. Shen. Ceramic wear maps. Wear 200 (1996), 154–175. 22â•… U. Serincan, G. Kartopu, A. Guennes, T. G. Finstad, R. Turan, Y. Ekinci, and S. C. Bayliss. Characterization of Ge nanocrystals embedded in SiO2 by Raman spectroscopy. Semicond. Sci. Technol. 19 (2004), 247–251. 23â•… S. C. Tjong. Electron microscope and Raman characterization of the surface oxides formed on the Fe–Cr alloys at 400–850°C. Mater. Char. 26 (1991), 29–44. 24â•… M. W. Barsoum, T. E. Raghy, and L. Ogbuji. Oxidation of Ti3SiC2 in air. J. Electrochem. Soc. 144 (1997), 2508–2516. 25â•… B. Bhushan. Principles and Applications of Tribology. John Wiley & Sons, New York, 1999, 413.
CHAPTER
13
CASE STUDY: TITANIUM DIBORIDE CERAMICS AND COMPOSITES Among various non-oxide ceramics, titanium diboride (TiB2)-based materials have attracted wider attention in the materials community because of their high hardness, elastic modulus, and superior thermal and electrical conductivity. To overcome the problems associated with poor sinterability of monolithic TiB2, several nonmetallic sinter-additives, such as TiSi2 and MoSi2, are added. This chapter addresses how the compositional design aspect (sinter-additive addition of up to 10â•›wt%) to obtain optimally sintered ceramics would have an influence on the tribological properties of borides. To address this aspect, the results of a series of fretting experiments on TiB2–MoSi2 and TiB2–TiSi2 against WC–(6â•›wt%)Co cermet will be discussed. The MoSi2 addition degrades neither the wear-resistance properties nor the frictional properties of TiB2, and microcracking-induced spalling was the dominant mechanism. In case of the TiB2–TiSi2 system, wear volume exhibits a linear dependence on the abrasion parameter, confirming the role of abrasive wear in material removal and damage.
13.1 INTRODUCTION As mentioned in the overview of structural ceramics given in Chapter 9, superior mechanical properties, in combination with the thermal and chemical properties, can potentially lead to better tribological properties of TiB2-based ceramics. Despite useful properties and wider potential applications, challenges involved in sintering monolithic TiB2 ceramics, in addition to their poor fracture toughness (∼4–5â•›MPaâ•›m1/2), limit the widespread application of such materials. Hence, considerable research efforts are invested in obtaining dense TiB2-based materials via various sintering techniques and using different metallic as well as nonmetallic additives.1–8 Extensive research has been undertaken to investigate the tribological behavior of hard materials such as ceramics and cermets9–19 with limited tribological work on TiB2.10,20–24
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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The tribological behavior of monolithic TiB2 as well as TiB2-containing cermets have been investigated using a fretting wear tester.25–27 The abrasive scratches and adhesion of tribochemical reaction products were observed after fretting against steel.25 The tribochemical reaction was influenced by the solubility of the binder phase in steel, in addition to the chemical dissolution of TiB2.25 Wasche et al.28 reported the influence of relative humidity on friction and wear behavior of pressureless sintered (PS) TiB2 ceramic against SiC and Al2O3. TiB2 exhibited low wear rates (<1╯×╯10−6â•›mm3/N m) when slid against SiC and Al2O3. Interestingly, the coefficient of friction (COF) showed a clear dependence on relative humidity, but not the wear, confirming experimentally that friction and wear are not necessarily correlated. In a different study, a comparative fretting wear study was made on TiB2 and TiB2– (5â•›wt%)MoSi2 against steel and WC-Co counterbody materials.29 The COF and specific wear rate results indicated the effectiveness of TiB2 in improving wear resistance in bulk materials and coatings.30–33 The tribological properties of TiB2 coatings were also studied to a limited extent, and the use of hard TiB2 coating is reported to increase wear resistance by a factor of 10 more than that of the bulk TiB2.34–39 It is well recognized that the tribological properties of composites are influenced by the physical, chemical, and mechanical properties and by the amount and distribution of the reinforcement phases. In addition, the counterbody material and operating parameters (sliding speed, load, humidity, etc.) also cast significant influence on the tribological response.40 Importantly, improved bulk properties, resulting from the incorporation of a secondary phase, may or may not lead to considerable improvement in friction and wear properties.41–46 For example, the addition of TiB2 to a SiC matrix enhances fracture toughness, and the incorporation of SiC into a Si3N4 matrix improves bend strength as well as fracture toughness, but in neither case does addition of SiC result in improved wear performance.41 However, the addition of either TiC or TiN to Si3N4 matrix is reported to improve wear performance substantially under tribological conditions where tribochemical wear is dominant.41 From this perspective, the development of any new material via compositional design calls for evaluation of the tribological properties. TiB2–TiSi2 and TiB2–MoSi2 with low silicide content (≤10â•›wt%) have been developed using a hot pressing route.47–49 In this chapter, the tribological properties of TiB2–TiSi2 and TiB2–MoSi2 ceramics are summarized, while more details can be found elsewhere.50,51
13.2 MATERIALS AND EXPERIMENTS As part of materials processing, appropriate amounts of TiB2 and MoSi2 powders were ground and mixed using a WC grinder to process TiB2–(X wt%)MoSi2 compositions (X╯=╯0, 2.5, and 10); more details of preparation of the starting powders are reported elsewhere.49 The hot pressing experiments were performed at selected temperatures of 1700°C and 1800°C (heating rate, 15°C/min) in vacuum (10−5â•›Pa) with a maximum applied pressure of 30â•›MPa for 1 hour. It has been reported47 that an optimal combination of density, hardness, and toughness can be obtained in TiB2–(5â•›wt%)TiSi2 and, accordingly, this chapter briefly describes the tribological
13.2 Materials and Experiments
TiB2
â•… 199
TiB2
MoSi2
TiB2 MoSi2
0.5 µm
Figure 13.1â•… STEM bright field image showing the phase assemblage and grain structure of TiB2–(10â•›wt%)MoSi2, hot pressed at 1700°C for 1 hour.49
properties of the same ceramic composition. TiB2–TiSi2 materials were hot pressed at a temperature of 1650°C for 1 hour, with an applied pressure of 30â•›MPa, in a flowing argon atmosphere from commercially available TiB2 (Grade F, H.C. Starck GmbH and Co., Goslar, Germany) and TiSi2 (Goodfellow Cambridge Limited, England) powders. The bright field scanning transmission electron microscope (STEM) image reveals the presence of polygonal TiB2 grains and the MoSi2 grains (gray contrast), revealing dislocation activity (see Fig. 13.1). In addition, grain boundary phases at the triple junctions can also be clearly observed. In Figure 13.2, a bright field conventional TEM image of TiB2–(10â•›wt%)TiSi2 consisting of various microstructural phases is shown. The Ti5Si3 phase is located at the triple pocket and surrounded by the TiB2 and TiSi2 grains. The morphology and size of Ti5Si3 grains appears to be a signature of liquid phase sintering. For the tribological tests, samples of 3â•›mm╯×╯4â•›mm╯×╯15â•›mm were machined from the hot pressed discs, using electrodischarge machining (EDM). The mechanical properties of as–hot-pressed ceramics are summarized in Table 13.1. The wear experiments were performed using a computer-controlled fretting testing machine with ceramic flat samples (TiB2/TiB2– MoSi2/TiB2-TiSi2) against cemented carbide (WC–(6â•›wt%)Co cermets) balls with a diameter of 10â•›mm. According to the commercial supplier’s data, the cemented carbide balls have hardness ∼ 16â•›GPa and toughness ∼ 14â•›MPaâ•›m1/2. All the experiments were performed at an oscillating frequency of 4â•›Hz and a linear stroke 100â•›µm, for a duration of 100,000 cycles. After the fretting tests, detailed characterization of the worn surfaces of the TiB2 samples and counterbody balls were performed using a laser surface profilometer and SEM energy-dispersive x-ray spectrometry (SEM-EDS).
200â•…
CHAPTER 13â•… Case Study: Titanium Diboride Ceramics and Composites
Ti5Si3 TiSi2
Ti5Si3
TiB2
TiB2
0.5 µm
Figure 13.2â•… Representative conventional TEM bright field image showing the grain morphology of the various constituent phases in hot pressed TiB2–TiSi2.48
TABLE 13.1.â•… Summary of the Densification and Mechanical Properties of the TiB2 Ceramics, Which Were Hot Pressed for 1 Hour in Vacuum47–50
Material
Monolithic TiB2 TiB2–(2.5â•›wt%)MoSi2 TiB2–(10â•›wt%)MoSi2 TiB2–(5â•›wt%)TiSi2 WC–(6â•›wt%)Co balla
Hot pressing temperature (°C)
Relative density (% ρth)
Vickers hardness, HV5 (GPa)
Elastic modulus (GPa)
Indentation toughness (MPa m1/2)
1800 1700 1700 1650 —
96.1 99.1 96.7 99.6 100
28.6╯±â•¯0.53 29.9╯±â•¯0.31 22.9╯±â•¯0.44 25.2╯±â•¯0.6 16
496.0╯±â•¯11 503.9╯±â•¯9 464.1╯±â•¯14 517.9╯±â•¯11 630.0
4.6╯±â•¯0.45 5.7╯±â•¯0.35 4.6╯±â•¯0.46 5.8╯±â•¯0.5 14
a
Data supplied by the commercial supplier.
13.3 TRIBOLOGICAL PROPERTIES OF TiB2–MoSi2 CERAMICS 13.3.1â•… Friction and Wear In this section, the tribological properties of TiB2–MoSi2 ceramics, as reported elsewhere,50 are summarized. The average COF (0.51–0.53) of the TiB2 composite was observed to vary insignificantly with load, and no significant change in COF of TiB2 with the addition of MoSi2 (up to 10â•›wt%) could be recorded. The wear depth profiles after testing at various loads are summarized in Figure 13.3; relatively maximum wear depth with wider area is observable at the highest load (10â•›N). Figure 13.4 plots the effect of load on the wear rate of the TiB2–MoSi2 ceramics sliding against
13.3 TRIBOLOGICAL PROPERTIES OF TiB2–MoSi2 CERAMICS
â•… 201
Results Profile 10.00
2N
µm 0.00
−10.00 0.00 mm
1.20 mm
10.00
5N
µm 0.00
−10.00 0.00 mm
1.20 mm
10.00
10 N
µm 0.00
−10.00 0.00 mm
1.20 mm
Figure 13.3â•… The maximum wear depth profiles obtained at various loading conditions of TiB2. Testing conditions: frequency, 4â•›Hz; 100,000 cycles; and stroke length, 100â•›µm.50
cemented carbide. A lower wear rate is measured at higher load (5 or 10â•›N). The wear of the TiB2 samples is also plotted against dissipated energy and abrasion parameter (Fig. 13.5) and the implication is discussed later. According to Wasche and Klaffke, TiB2 exhibited low wear rates (<1╯×╯10−6â•›mm3/Nâ•›m) against SiC and Al2O3.32 Yang et al. also reported the sliding friction and wear of monolithic TiB2 against SiC, Al2O3, and mullite.23,39 The COF and wear rate varied between 0.63–0.77 and 11.2–31.1╯×╯10−5â•›mm3/Nâ•›m, respectively, while low friction and wear were measured with a TiB2/SiC tribocouple. Tribofilm formation, oxide wear debris, and microcracking were observed to be the wear mechanisms. The combination of abrasion, adhesion, and tribochemical wear resulted in a higher wear rate (∼10−5â•›mm3/Nâ•›m) for TiB2 and TiB2–(20â•›wt%)MoSi2 after fretting against steel.29 These observations therefore indicate that friction and wear properties of the developed TiB2–MoSi2 ceramics against cemented carbide counterbody are better than or comparable with the earlier reported TiB2 tribosystems.
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4.5 TiB2 TiB2-2.5 MoSi2 TiB2-10 MoSi2
Wear rate (× 10−6 mm3/Nm)
4.0 3.5 3.0 2.5 2.0 1.5 1.0 0.5 0.0 2
5 Load (N)
10
Figure 13.4â•… Variation of wear rate of the TiB2–MoSi2 ceramics as a function of load after fretting against WC–(6â•›wt%)Co counterbody. Testing conditions: load, 2–10N; frequency, 4â•›Hz; 100,000 cycles; and stroke length, 100â•›µm.50
13.3.2â•… Wear and Dissipated Energy Wear is generally defined as the material removal process when materials are subjected to relative motion. The wear rate of TiB2 materials decreases with load, with maximum wear rate being measured at a load of 2â•›N (Fig. 13.4). It is evident that the wear volume and wear rates of the TiB2 ceramics are largely influenced by load and are not affected much either by the material compositions or by their mechanical properties. Material removal during sliding can be correlated with dissipated frictional energy at the contact.52 The friction between two sliding bodies increases the contact temperature and generates energy that is subsequently consumed in deformation, cracking, or tribochemical reactions. For a ball-on-flat fretting test configuration, the cumulative dissipated energy (Ed) can be determined from the area of the tangential friction (FT) versus displacement (S) loop50:
Ed =
∑ F S. T
(13.1)
The dissipated energy can be computed using the following relation:
Ed = µWvt ,
(13.2)
where μ is the average COF, W is the normal load, v is the fretting velocity, and t is the total testing duration. The variation in wear volume against dissipated energy for the investigated TiB2-based ceramics is plotted in Fig. 13.5a, and it is observed that the wear volume increased linearly with dissipated energy. Such a linear relationship indicates that the major wear mechanism is not changed considerably with the amount of MoSi2 sinter-additive. It also implies that a similar wear mechanism
13.3 TRIBOLOGICAL PROPERTIES OF TiB2–MoSi2 CERAMICS
â•… 203
Wear Volume (× 10−4 mm3)
2.6 2.4 2.2 2.0 1.8 1.6 TiB2 TiB2-2.5 MoSi2 TiB2-10 MoSi2
1.4 1.2 20
40
60 80 Dissipated Energy (J) (a)
100
120
Wear Volume (× 10−4 mm3)
2.6 2.4 2.2 2.0 1.8 1.6 TiB2 TiB2-2.5 MoSi2 TiB2-10 MoSi2
1.4 1.2 0.1
0.2
0.3
0.4
0.5
0.6
0.7
0.8
Abrasion Parameter ([W9/8/(KIc1/2 × H5/8)])
(b)
Figure 13.5â•… The variation of wear volume with the dissipated energy (a) and abrasion parameter (b) for the investigated TiB2-based ceramics against WC–(6â•›wt%)Co counterbody.50
is operative for different TiB2 compositions. The slope of the linear fit provides the proportionality factor (wear rate) and is estimated to be ∼1.3╯×╯10−6â•›mm3/N m.
13.3.3â•… Wear and Abrasion Parameter The wear resistance of brittle materials can also be analyzed by the combined influence of the material parameters, such as hardness and toughness, as well as operating parameters, such as load. All these parameters can be integrated into a single parameter: the abrasion parameter. The measured wear volume for TiB2-based ceramics is
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plotted in Fig. 13.5b against the abrasion parameter (α), which can be explained as follows53; W 9/8 α = 1/ 2 5 / 8 (13.3) K IC H where W is the applied load, KIc is the fracture toughness, and H is the hardness. If abrasion plays a major role in the material damage and removal, the wear volume (V) varies with the abrasion parameter (α) as
V ∝α
(13.4)
From Fig. 13.5b, the linear dependence of wear volume with the abrasion parameter confirms that the wear of TiB2 ceramics obeys Equation 13.4. At a given load, a ceramic with a larger value of (K I1C/ 2 H 5 / 8) therefore will have a smaller abrasion parameter and will exhibit higher wear resistance.
13.3.4â•…Material Removal Mechanisms To understand the wear mechanisms, the topographical observations of the worn surfaces were analyzed using SEM-EDS (Figs. 13.6–13.8). Figure 13.6 presents representative SEM images of monolithic TiB2 after sliding against various loads of 2, 5, and 10â•›N. From EDS analysis, it is observed that both the worn surface and wear debris consist of strong Ti and weak O peaks, which are indicative of the mild oxidation of TiB2. The worn surfaces on monolithic TiB2 are characterized by mild abrasive scratches, grain pullouts, and microcracking (Fig. 13.6). Importantly, the microcracks are observed perpendicular to the sliding direction and mainly caused the material removal (see Fig. 13.6b,d,f). At a number of regions within the wear scar, the intersection of microcracks facilitated the localized spalling of material. Figure 13.7 shows representative SEM images of the worn surface on TiB2–(10â•›wt%)MoSi2 ceramics. The worn surfaces are characterized by shallow abrasive grooves and microcracking.
13.4 TRIBOLOGICAL PROPERTIES OF TiB2–TiSi2 CERAMICS In a research paper published in 2010, the tribological properties of TiB2–TiSi2 against different mating materials were reported.51 In the case of TiB2–(5â•›wt%)TiSi2/ WC-Co, the COF was reported to lie between 0.49 and 0.56 and the wear rate was measured to be (1.2–4.1)╯×╯10−6â•›mm3/Nâ•›m, when fretted against a WC-Co counterbody. Figure 13.8 compares the worn surfaces on both the TiB2–(5â•›wt%)TiSi2 and WC-Co ball after fretting at a load of 10â•›N. SEM observation of the worn surface of the WC-Co ball reveals the occurrence of mild abrasion wear along with grain pullout. No evidence of transferred Ti can be found on an EDS spectrum obtained from the abraded area of cemented carbide ball (inset of Fig. 13.8d). Hence, the primary wear mechanism for the TiB2/WC-Co tribosystem is abrasion with limited tribo-oxidation.
13.4 TRIBOLOGICAL PROPERTIES OF TiB2–TiSi2 CERAMICS
â•… 205
Wear debris
250 µm
10 µm
(a)
(b)
Wear debris
10 µm
250 µm (c)
(d)
Ti
Ti
O
O
Ti
Ti
Wear debris
10 µm
250 µm (e)
(f )
Figure 13.6â•… Overview and details of as-fretted surfaces (SEM images) of monolithic TiB2 after 2-N (a,b), 5-N (c,d), and 10-N loads (e,f). Parts a, c, and e represent the topography of worn surfaces along with the wear debris; parts b, d, and f indicate the magnified view of worn surfaces. The EDS analysis of wear debris (inset e) and worn surface (inset f) are also shown. The single-pointed arrows in (b,d,f) indicate microcracks, while the dotted regions in (b,d,f) reveal regions of microcracking-induced localized spalling. Fretting conditions: frequency, 4â•›Hz; stroke length, 100 µm; and 100,000 cycles; counterbody, WC–(6â•›wt%)Co. (Double-headed arrows indicate sliding direction.).50
206â•…
CHAPTER 13â•… Case Study: Titanium Diboride Ceramics and Composites
100 µm (a)
10 µm (b)
Figure 13.7â•… (a) Overview of the worn surfaces of TiB2–(10â•›wt%)MoSi2 and (b) the magnified image of the worn surface. The single-headed arrows indicate microcracks. Testing conditions: load, 10N; frequency, 4â•›Hz; stroke length, 100 µm; and 100,000 cycles. (Double-headed arrows indicate the sliding direction.).50
13.5 CLOSING REMARKS An important observation has been the lower wear rate, on the order of 10−6â•›mm3/Nâ•›m, measured with hot pressed TiB2–MoSi2 and TiB2–TiSi2 ceramics after fretting against a harder (WC-Co) counterbody. The addition of silicides does not degrade the wearresistance property of TiB2, irrespective of load. Also, small amounts of silicides (up
13.5 Closing Remarks
â•… 207
Ti B Si O
10 µm
100 µm (a)
Ti
(b) W
O
5 µm
100 µm (c)
CoW
(d)
Figure 13.8â•… (a) Worn surfaces of TiB2–(5â•›wt%)TiSi2; (b) the magnified image of the worn surface along with the EDS; (c) the corresponding worn surface of WC–(6â•›wt%)Co ball; and (d) the magnified image of the worn surface along with the EDS. Testing conditions: load, 10N; frequency, 4â•›Hz; stroke length, 100 µm; and 100,000 cycles. (Arrows indicate sliding direction.).51
to 10â•›wt%) can be used as a sinter-aid to obtain a good combination of hardness and indentation toughness, without compromising on the frictional property or wear resistance. Also, a linear relationship between frictional dissipated energy and wear volume was established for TiB2 ceramics. Similarly, a linear relationship between wear volume and the abrasion parameter has been observed. The experimental results discussed in this chapter establish that the tailoring of both hardness and toughness is important to achieve better wear resistance. As a concluding note, various factors influencing the wear-resistance property of TiB2-based materials are summarized in Figure 13.9. From the materialsdevelopment perspective, the type and amount of sinter-additive as well as sintering conditions determine the microstructure and mechanical properties of the borides. Finer microstructure and better mechanical properties are ideally desired. Based on
208â•…
CHAPTER 13â•… Case Study: Titanium Diboride Ceramics and Composites
Better densification at lower sintering temperature and finer grain sizes
Sintering additive for TiB2
Tribochemical wear
(Type and amount)
Wear resistance of TiB2-based ceramics
Mechanical property enhancement and resistance to tribomechanical wear
Fig. 13.9â•… Summary of various factors that influence the wear-resistance properties of TiB2-based materials.
the results presented in this chapter, it can be said that enhancement of the mechanical properties can lead to better resistance against abrasive and brittle-fracture dominated wear behavior of borides. Also, the type and composition of the sinteradditive would be critical in the scenario in which the wear of borides is dominated by tribochemical wear.
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10â•… A. H. Jones, R. S. Dobedoe, and M. H. Lewis. Mechanical properties and tribology of Si3N4-TiB2 ceramic composites produced by hot pressing and hot isostatic pressing. J. Eur. Ceram. Soc. 21 (2001), 969–980. 11â•… B. Basu, D. Sarkar, and T. Venkateswaran. Pressureless sintering and Tribological properties of WC-ZrO2 composites. J. Eur. Ceram. Soc. 25 (2005), 1603–1610. 12â•… B. Basu, R. G. Vitchev, J. Vleugels, J. P. Celis, and O. Van Der Biest. Influence of humidity on the fretting wear of self-mated tetragonal zirconia ceramics. Acta Materialia 48 (2000), 2461–2471. 13â•… W. M. Rainforth. The sliding wear of ceramics. Ceram. Int. 22 (1996), 365–372. 14â•… J. Mukerji and B. Prakash. Wear of nitrogen ceramics and composites in contact with bearing steel under oscillating sliding condition. Ceram. Int. 24 (1998), 19–24. 15â•… J. Larsen-Basse. Binder extrusion in sliding wear in WC-Co alloys. Wear 105 (1985), 247–256. 16â•… M. Z. Huq and J. P. Celis. Expressing wear rate in sliding contacts based on dissipated energy. Wear 252 (2002), 375–383. 17â•… J. Pirso, S. Letunovits, and M. Viljus. Friction and wear behaviour of cemented carbides. Wear 257 (2004), 257–265. 18â•… J. L. Ortiz-Merino and R. I. Todd. Relationship between wear rate, surface pullout and microstructure during abrasive wear of alumina and alumina/SiC nanocomposites. Acta Materialia 53 (2005), 3345–3357. 19â•… B. Basu, J. Vleugels, and O. Vanderbiest. Fretting wear behavior of TiB2-based materials against bearing steel under water and oil lubrication. Wear 250 (2001), 631–641. 20â•… S. Torizuka and T. Kishi. Effect of SiC and ZrO2 on sinterability and mechanical properties of titanium nitride, titanium carbonitride and titanium diboride. Mater. Trans. JIM 37(4) (1996), 782–787. 21â•… B. Basu, J. Vleugels, and O. Vanderbiest. Unlubricated tribological performance of advanced ceramics and composites at fretting contacts with alumina. J. Mater. Res. 18(6) (2003), 1314–1324. 22â•… T. S. R. Ch. Murthy, B. Basu, A. Srivastava, R. Balasubramaniam, and A. K. Suri. Tribological properties of TiB2 and TiB2–MoSi2 ceramic composites. J. Eur. Ceram. Soc. 26 (2006), 1293–1300. 23â•… Q. Yang, T. Senda, N. Kotani, and A. Hirose. Sliding wear behavior and tribofilm formation of ceramics at high temperatures. Surf. Coatings Technol. 184 (2004), 270–277. 24â•… R. Wasche, D. Klaffke, and T. Troczynski. Tribological performance of SiC and TiB2 against SiC and Al2O3 at low sliding speeds. Wear 256 (2004), 695–704. 25â•… J. Vleugels, B. Basu, K. C. Kumar, R. G. Vitchev, and O. Vanderbiest. Unlubricated fretting wear of TiB2 containing composites against bearing steel. Metall. Mater. Trans. A 33(12) (2002), 3847–3859. 26â•… B. Basu, J. Vleugels, and O. Vanderbiest. Fretting wear behavior of TiB2-based materials against bearing steel under water and oil lubrication. Wear 250 (2001), 631–641. 27â•… T. S. R. C. Murthy, B. Basu, A. Srivastava, R. Balasubramaniam, and A. K. Suri. Tribological properties of TiB2 and TiB2–MoSi2 ceramic composites. J. Eur. Ceram. Soc. 26 (2006), 1293–1300. 28â•… R. Wasche, D. Klaffke, and T. Troczynski. Tribological performance of SiC and TiB2 against SiC and Al2O3 at low sliding speeds. Wear 256 (2004), 695–704. 29â•… A. Mukhopadhyay, G. B. Raju, and B. Basu. Understanding influence of MoSi2 addition (5 weight percent) on tribological properties of TiB2. Metall. Mater. Trans. A 39 (2008), 2998–3013. 30â•… S. C. Tjong and K. C. Lau. Abrasion resistance of stainless-steel composites reinforced with hard TiB2 particles. Comp. Sci. Technol. 60 (2000), 1141–1146. 31â•… A. V. Smith and D. D. L. Chung. Titanium diboride particle-reinforced aluminum with high wear resistance. J. Mater. Sci. 31 (1996), 5961–5973. 32â•… R. Wasche and D. Klaffke. Ceramic particulate composites in the system SiC-TiC-TiB2 sliding against SiC and Al2O3 under water. Tribology Int. 32 (1999), 197–206. 33â•… B. Prakash, E. Richter, H. Pattyn, and J. P. Celis. Ti–B and Ti–B–C coatings deposited by plasma immersion ion implantation and their fretting behavior. Surface and Coatings Technology 173 (2–3) (2003), 150–160. 34â•… M. Burger and S. Hogmark. Tribological properties of selected PVD coatings when slid against ductile materials. Wear 252 (2002), 557–565.
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35â•… M. Burger and S. Hogmark. Evaluation of TiB2 coatings in sliding contact against aluminium. Surface and coatings. Technology 149 (2002), 14–20. 36â•… M. Burger, M. Larsson, and S. Hogmark. Evaluation of magnetron-sputtered TiB2 intended for tribological applications. Surf. Coatings Technol. 124 (2000), 253–261. 37â•… M. Berger, L. Karlsson, M. Larsson, and S. Hogmark. Low stress coatings with improved tribological properties. Thin Solid Films 401 (2001), 179–186. 38â•… B. Prakash, C. Ftikos, and J. P. Celis. Fretting wear behavior of PVD TiB2 coatings. Surf. Coatings Technol. 124 (2000), 253–261. 39â•… Y. Yang, Z. Zheng, X. Wang, X. Liu, J. G. Han, and J. S. Yoon. Microstructure and tribology of TiB2 and TiB2/TiN double-layer coatings. Surf. Coatings Technol. 84 (1996), 404–408. 40â•… B. Bhushan. Principles and Applications of Tribology. John Wiley & Sons, New York, 1999, 487. 41â•… A. H. Jones, R. S. Dobedoe, and M. H. Lewis. Mechanical properties and tribology of Si3N4–TiB2 ceramic composites produced by hot pressing and hot isostatic pressing. J. Eur. Ceram. Soc. 21 (2001), 969–980. 42â•… T. Venkateswaran, D. Sarkar, and B. Basu. Tribological properties of WC-ZrO2 nanocomposites. J. Am. Ceram. Soc. 88(3) (2005), 691–697. 43â•… S. D. Bakshi, B. Basu, and S. K. Mishra. Fretting wear properties of Sinter-HIPed ZrO2-ZrB2 composites. Composites A 37 (2006), 1652–1659. 44â•… T. S. R. Ch. Murthy, B. Basu, A. Srivastava, R. Balasubramaniam, and A. K. Suri. Tribological properties of TiB2 and TiB2–MoSi2 ceramic composites. J. Eur. Ceram. Soc. 26 (2006), 1293–1300. 45â•… E. M. Jayasingh, P. S. Tantri, T. A. Bhaskaran, S. K. Biswas, and S. K. Ramasesha. Performance of monolithic and TiB2 reinforced MoSi2 in dry sliding contact with steel. Mater. Lett. 53 (2002), 379–383. 46â•… B. Basu, J. Vleugels, and O. Vanderbiest. Fretting wear behaviour of TiB2-based materials against bearing steel under water and oil lubrication. Wear 250 (2001), 631–641. 47â•… G. B. Raju and B. Basu. Densification, sintering reactions and properties of titanium diboride with titanium disilicide as a sintering aid. J. Am. Ceram. Soc. 90(11) (2007), 3415–3423. 48â•… G. Brahma Raju, K. Biswas, and B. Basu. Microstructural characterization and isothermal oxidation behavior of hot-pressed TiB2-10 wt% TiSi2 composite. Scr. Mater. 61 (2009), 104–107. 49â•… G. Brahma Raju, K. Biswas, A. Mukhopadhyay, and B. Basu. Densification and high temperature mechanical properties of hot pressed TiB2-(0–10 wt. %) MoSi2 composites. Scr. Mater. 61 (2009), 674–677. 50â•… a) G. Brahma Raju; Processing-Microstructure-Property Correlation of High Temperature TiB2Silicide Ceramics; PhD thesis, Indian Institute of Technology Kanpur, India, February, 2009. b) G. Brahma Raju and B. Basu. Influence of MoSi2 addition on load-dependent fretting wear properties of TiB2 against cemented carbide. J. Am. Ceram. Soc. 92(9) (2009), 2059–2066. 51â•… G. Brahma Raju and B. Basu. Wear mechanisms of TiB2 and TiB2-TiSi2 at fretting contacts with steel and WC-6 wt. % Co. Int. J. Appl. Ceram. Technol. 7(1) (2010), 89–103. 52â•… M. Z. Huq and J. P. Celis. Expressing wear rate in sliding contacts based on dissipated energy. Wear 252 (2002), 375–383. 53â•… A. G. Evans and D. B. Marshall. Wear Mechanisms in Ceramics: Fundamentals of Friction and Wear, D. A. Rigney (Eds.). Am. Soc. Metals, Metals Park, OH, 1981, 439.
SECTION
III
FRICTION AND WEAR OF BIOCERAMICS AND BIOCOMPOSITES
CHAPTER
14
OVERVIEW: BIOCERAMICS AND BIOCOMPOSITES This section of the book discusses some case studies to illustrate how wear takes place in some of the potential biomaterials—in particular, bioceramics or ceramiccontaining biocomposites. In this overview chapter, some generic definitions related to biomaterials and their fundamental property requirements, as well as the importance of friction and wear in vitro to the biological performance of biomaterials, are discussed. In the last few decades, impressive progress has been recorded in terms of developing new materials to obtain better performance in biomedical applications. The success of such efforts clearly demands better understanding of various concepts, for example biocompatibility, host response, and cell–biomaterial interaction, as well as in vitro and in vivo evaluation of properties, including tribological properties. In this section, research results obtained with Al2O3/glass-infiltrated Al2O3 and polymer-ceramic biocomposites are summarized. The physical as well as tribological properties of polyethylene (PE)-based hybrid biocomposites will be discussed to illustrate the concept of how the physical and wear properties can be enhanced along with biocompatibility due to combined addition of bioinert and bioactive ceramics to a bioinert polymeric matrix.
14.1 INTRODUCTION The lack of durability of biomaterials in vivo and the potential genotoxic effect due to the release of fine wear debris particles are recognized as some of the potential problems, which among many other factors often necessitate revision surgery. It is therefore important to evaluate the tribological properties so that the durability and mechanism of wear resistance of biomaterials can be evaluated. The importance of biomaterials can be well realized from their economical aspect, that is, in terms of an estimate of total health care expenditure around the world. In the most developed country of the world, the United States, total health care expenditure in the year 2000 was approximately $14 billion. It was also reported that the U.S. market for biomaterials in 2000 was $9 billion. It can be further noted that
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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Biological Performance of Materials: Host Response Material Response Structure/Composition/ Function Relationships in Manufactured and Natural Materials Foundation Disciplines
Interactions
NonLiving Materials
Engineering Physical Sciences
Living Materials (Patient)
Medicine Biological Sciences
Figure 14.1â•… Concept triangle illustrating the synergistic interaction of engineering and biological science disciplines involved in designing biomaterials. The schematic also demonstrates the multidisciplinary approach of the science and technology of biomaterials.1
the respective annual expenses in other countries are typically around 2–3 times the U.S. expenses.1 To this end, the development of biomaterials and related devices is important. The design of biomaterials requires the synergistic interaction of materials science, biological science, chemical science, medical science, and mechanical science. Such interaction is illustrated in Figure 14.1, which also shows the need to develop a cross-disciplinary approach to designing new biomaterials. A summary of a number of existing issues, which are to be dealt with while developing new biomaterials and, in particular, bioceramic implant materials, is provided in Figure 14.2. Despite significant research on biomaterials, it is difficult for any synthetic material to mimic the extremely complex structure of bone in all aspects and that the fundamental drawback of synthetic materials is that they cannot repair themselves as living bone does. Among different kinds of biomaterials, metals and metallic alloys are used in orthopedics, dentistry, and other load-bearing applications; ceramics are used with emphasis on either their chemically inert nature or their high bioactivity; polymers are used for soft-tissue replacement and many other nonstructural biomedical applications. To achieve better biological properties and mechanical strength, composite materials of metals, ceramics, and polymers are being developed and clinically assessed to a limited extent. The modulus and strength properties exhibited by natural tissues, such as cortical/cancellus as well as enamel and dentine, are summarized in Table 14.1. Broadly, all biomaterials are being developed to maintain a balance between the mechanical properties of the replaced tissues and the biochemical effects of the material on the tissue. However, in most (if not all) biological systems, a range of properties is required: for example, biological activity, mechanical strength, and chemical durability. Therefore, often a clinical need can only be fulfilled by a designed material that exhibits a complex combination of some of those properties.
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14.2 Some Useful Definitions and Their Implications
Revision surgery due to prosthetic infection, aseptic loosening
Slow host response and longer healing time (≥6 weeks)
Existing issues with synthetic implants Lack of fracture resistance
HAp-inherently insulator (Natural bone: piezoelectric in vivo)
Figure 14.2â•… Schematic illustration of various issues concerning biomaterials and, in particular, bioceramic implants.
TABLE 14.1.â•… Mechanical Properties of Different Hard Tissues of Human Osseous System (Summarized from Reference 1)
Tissues Cortical bone Cancellous bone Enamel Dentin
Elastic modulus (GPa)
Tensile strength (MPa)
17.7 12.8 0.4 11.0
133 52 7.4 39.3
14.2 SOME USEFUL DEFINITIONS AND THEIR IMPLICATIONS 14.2.1â•… Biomaterials Broadly, biomaterials can be defined as synthetic materials that can be designed to induce a specific biological activity.2 The major difference of biomaterials from structural materials is their ability to remain in a biological environment without damaging the surroundings and without getting damaged in the process.3 It must be emphasized here that the properties and response of a material in a physiological environment are, by far, the most important consideration for selecting and defining biomaterials. From the health care perspective, it is desirable that a biocompatible material interrupts normal body functions as little as possible.
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14.2.2â•… Biocompatibility The fundamental requirement of any biomaterial concerns the ability of the material to perform effectively with an appropriate host response for the desired application; that is, the material and the tissue environment of the body should coexist without having any undesirable or inappropriate effect on each other. Such a requirement is broadly described by a concept known as biocompatibility.4 Broadly, biocompatibility is defined as “the ability of a material to perform with an appropriate host response in a specific application.” A more recent definition was proposed by Williams5: “Biocompatibility refers to the ability of a biomaterial to perform its desired function with respect to a medical therapy, without eliciting any undesirable local or systemic effects in the recipient or beneficiary of that therapy, but generating the most appropriate beneficial cellular or tissue response in that specific situation, and optimizing the clinically relevant performance of that therapy.” Biocompatibility arises from the acceptability of nonliving materials (synthetic biomaterial) in a living body (mammals, humans). There are three important aspects of biocompatibility, which a candidate biomaterial should achieve in diverse environments, such as bone, blood vessels, and the eye. In the first place, biomaterials must be biochemically compatible, nontoxic, nonirritative, nonallergenic, and noncarcinogenic; second, they must be biomechanically compatible with surrounding tissues; and third, a bioadhesive contact must be established between the materials and living tissues. It needs to be emphasized here that biocompatibility depends on where a material is to be used; for example, a specific material could be biocompatible in bone replacement application, but the same material may not be biocompatible in a direct blood contact application. As is discussed later, a range of in vitro and in vivo tests are suggested to completely describe the biocompatibility of a material.
14.2.3â•… Host Response To develop new materials, it is desirable to understand the in vivo host response to various biomaterials. Ideally, biomaterials should not induce any change or provoke any undesired reaction in the neighboring or distant tissues. An additional objective for an implant would be the formation of a structural and biological bond between the material and host tissues. When biocompatibility is lacking, materials cause tissue reactions, which may be systemic or local. According to International Organization for Standardization (ISO) 10993-11 standards, a systemic effect can be categorized on the basis of the severity of postimplantation time-dependent response in the osseous system: (1) acute response, if recorded within 24 hours of implantation; (2) subacute, if observed within 14–28 days; (3) subchronic, if observed within up to 90 days or within 10% of the animal’s life span; and (4) chronic, if observed after more than 90 days or more than 10% of the animal’s life span. In general, various biomaterials, on the basis of biocompatibility and host response, can be classified into following categories6: a. Bioinert or biotolerant: Bioinert materials cannot induce any interfacial biological bond with bone. Many bioinert materials (e.g., Al2O3, ZrO2) have better physical properties. It must be mentioned that bioinertness does not
14.3 Experimental Evaluation of Biocompatibility
â•… 217
necessarily mean that biological cells will not adhere on material surfaces and, in fact, cell adhesion is experimentally observed on these material surfaces without any cytotoxic effect. b. Bioactive: Bioactive materials can attach directly with osseous tissues and form biological bonds during the early stage of implantation. Examples are 45S5 bioglass and calcium phosphates (e.g., hydroxyapatite [HA or HAp]). c. Bioresorbable: Bioresorable materials are gradually resorbed during the postimplantation period and finally are replaced by new tissues, in vivo. Examples are tricalcium phosphate (TCP) and bone cement.
14.3 EXPERIMENTAL EVALUATION OF BIOCOMPATIBILITY It is recommended that any research program involving development of new biomaterials must include a range of in vitro and in vivo tests, and the detailed protocols are stated by various standard agencies, for example, the ISO. As far as biological test protocols are concerned, it is important to mention the difference between in vitro and in vivo tests. In vitro tests are laboratory-scale simulated experiments, which are rapid and are a must as initial screening tests. However, most in vitro experiments use a single cell line, which does not simulate the interaction of material with actual tissue (which comprises multiple cell types) in vivo. In addition, in vitro experiments provide very poor representation of physiological conditions. These tests, nevertheless, are effective as the first step of biocompatibility evaluations. On the other hand, in vivo experiments clearly produce a better approximation to a human environment. Here, the materials are placed in direct contact with different cell types, and the effect of hormonal factors can be analyzed. Also, in vivo tests provide important information on interactions of materials with extracellular matrix, blood-borne cells, protein, and molecules. These experiments can be regarded as the second step prior to clinical use. To harmonize the existing biological test guidelines, ISO prepared a document to provide guidance on test selection, “Biological Evaluation of Medical Devices—Part 1: Evaluation and Testing within a Risk Management Process” (ISO 10933-1), which contains all the national and international documents. According to ISO 10993 requirements for “long-lasting tissue/ bone implant,” the following biological tests are to be followed: • Cell culture: This is an in vitro test and is widely used to assess the cellular proliferation and adhesion on material substrates. Cell culture experiments are performed in the laboratory using a cell line that is relevant for a specific application and the cells are seeded on the materials. The first stage involves sterilization of the samples, typically in an autoclave at 15â•›psi, 121°C for 20 minutes. Ultraviolet (UV) sterilization is also adopted, if the material degrades in a steam autoclave. For most of the cells that form bone (osteoblast) or connective tissue (fibroblast), the culture medium is Dulbecco’s modified Eagles’ medium (DMEM), containing 10% serum, 1% antibiotic cocktail. The samples are incubated for various timescales (1–7 days or more) at 37°C (human body
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temperature). Following this, the cells are fixed in glutaraldehyde/formaldehyde and the cell proliferation or adhesion is observed using fluorescence microscopy or scanning electron microscopy (SEM). In Figure 14.3, the experimental steps followed in culturing mammalian cells on material substrates are summarized. • Biochemical assays for cellular functionality: The MTT assay is a widely used colorimetric assay (an assay that measures changes in color) for quantifying the mitochondrially active biological cells. MTT [3-(4,5-dimethylthiazol2-yl)-2,5-diphenyltetrazolium bromide, a tetrazole], a biochemical reagent, forms the dark blue formazan product by reduction of the tetrazolium ring of MTT by succinate dehydrogenase, a mitochondrial enzyme.3 The absorbance of the treated solution can be determined by measuring the optical density (OD) at a certain wavelength (usually between 500 and 600â•›nm) by a spectrophotometer. This reduction is associated with the activity of mitochondrial reductase enzymes and, therefore, the conversion is directly related to the number of viable (living) cells. Another biochemical assay, the alkaline phosphatase (ALP) assay, is considered as an early-stage differentiation marker for bone cell differentiation, such as osteogenesis. It is known that ALP enzyme is bound to the cell membrane of osteoblasts and functions to promote osteogenesis by degrading pyrophosphates.3 The osteocalcin (OC) assay is useful to assess the late-stage osteoblast differentiation ability of a biomaterial.3 • Genotoxicity: In this in vitro experiment, the DNA-damaging capability of biomaterial eluate is assessed using single cell gel electrophoresis (SCGE; also called the comet assay) and the micronucleus assay. As part of genotoxicity testing, the comet assay is widely used and the protocol proposed by Tice et al. is generally followed.7 After treatment of cells with finer particles, the results are quantified using comet parameters, obtained from the analysis of fluorescence microscope images as well as using an image analysis system (Komet 5.0, Kinnetic Imaging, Liverpool, U.K.). The evaluation of cellular functionality and genotoxicity of wear debris particles is important to ensure safety during their long-term use in biomedical applications. The finer debris particles have a combination of unique physiochemical properties, including large ratio of surface area to volume and high chemical activity. There has been a general consensus that implants, such as those made of metallic, polymeric, and ceramic materials, generate a large number of de novo particles, which cause aseptic loosening followed by inflammation by macrophages in vivo. It is therefore important to study the effects of de novo generated wear particles at the cellular and genetic level and to get an insight into the mechanism and effect of nano- and micronsize wear particles.8,9 Sample preparation, stability, preparation medium, test method, and reference material are important considerations in evaluating cytotoxicity and genotoxicity of de novo generated particles.10–12 The
219
ELISA reader
ALP Kit Osteocalcin Kit
on
esi
dh
ll a
Ce
Observation in fluorescence microscope
Staining nucleus with HOECHST dye in blue color
SEM observation
Staining cytoskeleton with Phalloidin dye in green color
Drying with HMDS (Hexamethyl disilazane)
Dehydration
Cell seeding
Figure 14.3â•… Schematic illustration of various experimental steps involved in cell culture protocol to observe cellular adhesion on biomaterials, as well as various biochemical assays to quantify cell viability and early- and late-stage osteoblast differentiation (adapted from Reference 3).
MTT Kit
Incubation
Sterilization of the samples
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commonly observed size range of wear debris particles is from nanometric to submicron size.13,14 The purpose of genotoxicity tests is to screen implant materials for their toxic effect at the genetic (DNA damage) level. Such biological assays are necessary to comply with the Organization for Economic Co-operation and Development (OECD) and ISO guidelines “in vitro tests for genotoxicity” (OECD 473, ISO 10993-3).15 The selection of the combination of cytotoxicity and genotoxicity test methods, concentration, and exposure time as well as the type of particulate materials (ceramics, polymers, metals, and composites) are also equally important. For example, Papageorgiou et al. studied the DNA damage and chromosome aberrations that occur after human cells are treated with cobalt-chrome alloy particles, using SCGE assay.16 It has also been reported that single- and double-strand breaks frequently were caused by metal ions as well as by particles, whereas ceramic nanoparticles induced singlestrand breaks and DNA damage at alkaline labile sites.17,18 • Implantation and histopathology investigation: Implantation in animal bones is an important in vivo experiment for hard tissue and bone replacement materials, where the sample of a predefined shape is placed in the bone defect of a mammal (rabbit, rat, and mouse). After the desired time period, the samples and surrounding tissues are examined histopathologically to investigate the in vivo or host response to the materials. In general, short-term implantation tests are performed for up to 12 weeks and long-term tests for up to 78 weeks. It is important to mention that the number of animals should be limited to a minimum from an animal-welfare point of view. For implantation experiments, control samples, for example HAp or ultra-high-molecular-weight polyethylene (UHMWPE), are used. An example of how bioceramic implants are implanted in a rabbit femur is shown in Figure 14.4. During the operation under a clean and sterilized environment, cylindrical holes are normally drilled on the femur and test samples are implanted in the holes; subsequently, the wound is closed using stitches. The implantation sites are macroscopically examined using x-ray radiography for any evidence of host response. Following standard protocols, thin sections of implant with natural bone are stained with Stevenels’ blue with counterstaining with Van Gieson’s Picro-Fuschin to observe cellular activity as well as neobone formation around the implant, using fluorescence microscopy, SEM, and atomic force microscopy (AFM). The entire procedure for preparation of samples for histopathological analysis is shown in Figure 14.5. Sufficient knowledge about the tribological properties of biomaterials is essential in the evaluation of local or systemic effects, which can be caused to patients. For example, the wear of orthopedic alloys causes the release of wear debris particles in the human tissues and has to be assessed with respect to amount, size, and shape as well as kinetics and chemical state of the ions. Currently, there are no standard practices, methods, or guidelines for the evaluation of wear of orthopedic alloys, and of the products formed from wear. In vitro methods have been developed to test changes occurring in an environment simulating the in vivo conditions. Typically, the in vitro tests involve
14.4 Wear of Implants
(a)
(b)
(c)
(d)
â•… 221
Figure 14.4â•… Digital camera images showing (a) the pin-shaped (2-mm diameter and 6-mm length) bioceramic implants, prior to implantation, (b) rabbit’s femur before implantation, showing the implantation sites (holes), (c) control samples inserted inside the holes, and (d) HAp–(20â•›wt%)mullite samples, implanted inside the holes.3
0.9% (w/v) NaCl solution or Hank’s balanced salt solution. Also, tests have been carried out in immersing the materials in solutions containing proteins, enzymes, lactic acid, and so on, or in physiological fluids such as blood and sweat. For some material systems, such as various glass-ceramics, the biomineralization (i.e., formation of Ca-P-rich layer) in vitro is also assessed.
14.4 WEAR OF IMPLANTS In view of better physical properties (strength, hardness, ductility), some biocompatible metals and their alloys are often used for joint and bone implants. The applications of metals include bone replacement, bone repair, metal plates for fractures, dental implants (fillings and posts), screws, staples, and parts of other devices such as artificial hearts, pacemakers, and catheters. The major issues for metallic implants are biocompatibility, wear resistance, and corrosion resistance in body fluid. In the following, such issues will be discussed first and this will be followed by a discussion on the properties and applications of various implant metals and alloys in reference to earlier reported literature.
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(d)
(c)
(b) (a)
(e)
Dehydration in alcohol
Embedding in PMMA
Embedded bone
Implants Sections cut by diamond saw
(i) Portions cut by diamond saw
(f)
Rabbits’ femur after fixation
(h)
Staining
Stevenels’ blue followed by Van Gieson’s Picrofuschin
Optical microscopy observation
(g)
Thin section
Diamond polishing
Figure 14.5â•… Schematic illustration of the major steps involved during the sample preparation for histopathology analysis: (a) long femur of the rabbit after fixation, where the red circles indicate implants; (b) dehydration of bone pieces along with implants; (c) embedding in poly(methyl methacrylate) (PMMA) polymer; (d) embedded bone piece in PMMA, after removing from the bottle; (e) cutting of thin sections by diamond saw, where thin dotted lines indicate the cutting path; (f) thin section showing the bone (yellow) and bone marrow (pink) embedded in PMMA matrix; (g) polishing of thin section using diamond paste; (g) staining in Stevenels’ blue to observe various cells; and (i) fluorescence microscopy observation.3
For implant applications, wear of metallic implants is a serious concern for various biomaterials. The examples are different joints of human body, where two similar or dissimilar materials come in contact. In a typical hip prosthesis, a metallic stent is attached to a ceramic ball and the ceramic ball moves inside the polymeric acetabular cup. At the joint of metal–ceramic interface, fretting fatigue could be responsible for the implant’s loosening. On the other hand, the ceramic ball–polymer cup interface experiences sliding wear. The presence of body fluid and different types of proteins may trigger the wear rate in vivo. Pazzaglia et al.19 investigated the reason behind the loosening of metal–plastic total hip prostheses. Metal-on-plastic total hip prostheses liberate metal particles due
14.5 Coating on Metals
0.6
CP Titaniun
0.5
Ti-13Nb-13Zr Ti-6Al-4V
0.4 COF
â•… 223
Co-28Cr-6Mo
0.3 Ti-5Al-2.5Fe
0.2 0.1 0.0
0
2000
4000 6000 8000 NO. OF CYCLES
10,000
Figure 14.6â•… Comparative plot, showing the evolution of coefficient of friction (COF) with number of fretting cycles for potential orthopedic implant materials against steel at 10-N load, 10-Hz frequency, and 80-µm displacement stroke.20
to wear of the femoral stem. Such relative movement between two surfaces presumably aids the passage of metal particles from the cement–metal interface to the cement-tissue interface in the absence of direct contact via fissures in the cement. Metallic wear debris may also be shed from the femoral head articulating in the cup if the latter is contaminated with an abrasive, such as bone cement; however, this was not noticed. Irrespective of the mechanism of their production and release, the particles of stainless steel and Co–Cr–Mo alloy at the bone–cement interface encourage macrophage-related bone resorption.19 It was suggested that this aspect represents a contributing, and in some cases not inconsiderable, factor in loosening of metal–plastic total hip prostheses. In a study on tribological behavior of Ti-based alloys, Choubey et al.20 conducted fretting wear experiments on a number of Ti-alloys, in a simulated body fluid (SBF) environment. Their results revealed that the COF of Ti–6Al–4V alloy lies between 0.46 and 0.50 and the COF of Ti–5Al–2.5Fe alloys is 0.3 (see Fig. 14.6). The major wear mechanism was found to be tribomechanical abrasion, transfer-layer formation, and cracking.
14.5 COATING ON METALS An impetus for developing various coatings is that the properties of coated bioimplants will combine advantageous properties of both coating and substrate materials. For example, bioactive ceramic coatings on metallic substrates combine the good strength of metal with the good bioactivity of the ceramic coating. The coatingsubstrate adhesion or deposition route also influences the physical properties of coatings. For example, HAp-containing glass coating on Ti dental implants is reported to have better adhesion than flame-sprayed HAp coatings. An in vitro study
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to investigate the biological response of HAp/Ti–6Al–4V composite21 coatings in SBF solutions revealed that the coatings can undergo two-stage biointegration processes, that is, dissolution during the initial 4 weeks’ soaking in SBF and the subsequent bonelike apatite crystal precipitation. With the interposition of a composite bond coat composed of 50â•›vol% HAp and 50â•›vol% TiO2, a composite coating on titanium substrate22 was processed using plasma spraying, and no chemical reaction was observed between HAp and TiO2. Godley and co-workers23 demonstrated that chemical treatment can be used to create a calcium phosphate (CaP) surface layer, which might provide the alkalitreated Nb metal with bone-bonding capability. The formation of a similar CaP layer upon implantation of alkali-treated Nb into the human body can promote the bonding of the implant to the surrounding bone. This bone-bonding capability could make Nb metal an attractive material for hard tissue replacements and this requires further investigation. Pajamaki et al.24 studied the effect of glass-ceramic coating on titanium implants. They exhibited the results with uncoated Ti and showed that after 52 weeks the coated metal showed 78% bone ingrowth, whereas in the case of uncoated metal, the bone coverage was only 37%. Munting25 compared the in vivo biocompatibility of implants without coating, with a pure titanium coating, and with a HAp coating. It was shown that HAp coating suffers from limited strength and poor fatigue strength.25 Also, HAp can be dissolved in vivo in an acidic environment, created by macrophages. Wang et al.26 compared the in vivo bone apposition on plasma-sprayed and electrochemically deposited HAp coatings on Ti–6Al–4V alloy with uncoated Ti–6Al–4V alloy. It was revealed that plasma-sprayed HAp coatings can induce higher bone apposition ratios than those exhibited by uncoated Ti–6Al–4V and electrochemically deposited HAp coatings after 7 days. However, after 14 days of implantation, both the coated materials exhibited similar bone apposition ratios, much higher than that for uncoated Ti–6Al–4V. In a different work, Choubey et al.27 investigated the tribological behavior, under fretting contacts, of diamondlike carbon (DLC) coated Co–Cr–Mo alloys, wherein the coating was synthesized via chemical vapor deposition (CVD). The wear mechanism was mainly governed by the wear with no significant change in surface morphology in DLC-coated flat material. The uncoated Co–Cr–Mo exhibited extensive plastic deformation and delamination-induced grain pullout at higher load (10N) after 10,000 cycles (see Fig. 14.7).
14.6 GLASS-CERAMICS Glass-ceramics are fabricated by prolonged heat treatment of glass at elevated temperature (above the glass transition temperature Tg). This is an alternative route for synthesizing polycrystalline ceramics. The amount of crystalline ceramic phase may vary between 50 and 99â•›vol%. Since 2000, a number of glass-ceramic systems are being researched for their biomedical potential, in particular, dental restoration applications. It can be recalled here that human teeth act as a mechanical device during masticatory processes such as cutting, tearing, and grinding of food particles.
14.6 Glass-Ceramics
93 µm
44 µm
(a)
â•… 225
(b)
Figure 14.7â•… Optical microscopy images of the worn surface on DLC-coated Co–Cr–Mo surface illustrating (a) the mild wear after fretting at 10-N load for 10,000 cycles and (b) the severe wear after fretting at 10-N load after 100,000 cycles. The counterbody is bearing steel. DLC coating is worn through as seen by the white contrast in (b).27
Teeth get damaged or worn away with age and, therefore, partial or total replacement of human teeth with a suitable biocompatible material is required. Among various glass-ceramic materials, 45S5 glass, originally invented by Clupper and co-workers, is known to be the best bioactive glass (BG) material.28 Their typical composition29 (all in wt%) is SiO2, 46.1; P2O5, 2.6; CaO, 26.9; Na2O, 24.4. However, the poor mechanical properties and lack of machinability of 45S5 glass have been a major concern. Therefore, these materials are unsuitable for dental applications or any biomedical applications requiring a complex shape. Besides 45S5 glass, Bioverit®I base glass has wider clinical applications, including the possibility of using Bioverit II base glass as matrix for Ti-particle-reinforced composite30 coatings. Mechanical characterization showed a good adherence of the coatings to the substrate. In another study, Bioverit® III base glass and composites of glassceramic matrix with Ti particles were prepared using pressureless sintering.31 The cell growth of fibroblasts on the surface of the glass-ceramic-matrix composites confirms their biocompatibility. In a different system, complex reactions between Ti and HAp occurred during the sintering of Ti/HAp/BG composites.32 Composites of a Bioverit®III glass-ceramic33 matrix with yttria–partially stabilized zirconia (YPSZ) particles, successfully prepared using pressureless sintering, also confirmed the toughening effect of the Y-PSZ particles. Highly dense (>98% of relative density) Si3N4–bioglass composite34 has the potential advantages of each constituent, that is, the high fracture toughness of Si3N4 with the bioactivity of a bioglass. The most significant result is the improvement in fracture toughness (4.4â•›MPaâ•›m1/2) and bending strength (383╯±â•¯47â•›MPa) with respect to currently used bioceramics, glasses, and glass-ceramics for load-bearing applications. Barrors et al.35 investigated the in vivo bone tissue response of another fluoride-containing canasite glassceramic (0.47K2O, 0.94Na2O, 1.42CaO, 5.67SiO2, 1.5CaF2). The canasite formulation evaluated was not osteoconductive and appeared to degrade in the biological environment.
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14.7 BIOCOMPATIBLE CERAMICS 14.7.1â•… Bioinert Ceramics One potential bioinert ceramic, Al2O3, showed good performance in vivo, although low fracture toughness (3–4â•›MPa m1/2) restricts its use in demanding applications. On the other hand, tetragonal zirconia ceramic (8–11â•›MPa m1/2) has an edge over alumina. However, various attempts have been made to toughen Al2O3 by adding monoclinic ZrO2, partially stabilized ZrO2 (PSZ), and so on. This leads to the development of zirconia-toughened alumina (ZTA).36 Dense ZTA has considerably better toughness37 and wear resistance38,39 than monolithic alumina. Although ceramic composite materials have a potential use in load-bearing orthopedic applications, very few materials have been tested clinically so far. In an important investigation, Hayashi et al.40 studied the in vivo response of bioinert ceramics such as alumina ceramic (99.5% purity Al2O3), zirconia ceramic (5â•›wt% Y2O3-stabilized ZrO2), and SUS316L stainless steel (Fe, 65%; Cr, 18%; Ni. 13%; MO, 2%; Mn, 2%). The results were also compared with dense sintered HAp. The push-out test result revealed the bone–implant interface shear strength for alumina, zirconia, steel, and HAp as ∼0.8, 0.9, 0.5, 12.1â•›MPa, respectively. It was concluded that the bioinert ceramics should not be used as a bone-bonding material, but as the material for the articulating surface. Colon et al.41 described the function of osteoblast and Staphylococcus epidermidis on nanophase ZnO and TiO2 inert bioceramics. It was evident from their result that nanophase ceramics (ZnO, TiO2) could decrease S. epidermidis adhesion and increased osteoblast adhesion compared with macropahse materials. A more recent investigation42,43 of Ca- and MgO-doped zirconia revealed that 8â•›mol% CaO-doped PSZ ceramics possess 97.5% ρth and 16â•›mol% CaO-doped fully stabilized zirconia (FSZ) possess only 91.6% ρth, whereas more than 95% ρth in both Mg-PSZ and Mg-FSZ can be obtained, when all are microwave (MW) sintered at 1585°C for 1 hour. The optimized microstructure of both Mg-PSZ and Mg-FSZ samples are characterized by the presence of coarser grains with size in the range of 5–10â•›µm. Apart from microstructural investigation and measurement of mechanical properties, tribological properties were also evaluated, under dry and SBF conditions. For both Mg- and Ca-doped ZrO2, a steady-state COF of ∼ 0.5 against a bearing steel ball was measured under dry conditions, and it showed lower values (∼0.35–0.4) in SBF lubrication contact. The wear mechanism was dominated by the formation of FexOy-rich tribochemical layer. The investigated CaO/MgO-doped ZrO2 materials experienced wear rates on the order of 10−5â•›mm3/Nâ•›m (in air) and 10−6 to 10−7â•›mm3/Nâ•›m (in SBF), with the lowest wear rate recorded with FSZ materials in SBF solution. However, further experiments to study cytotoxicity and clinical trials need to be conducted on these materials to assess their potential for biomedical applications.
14.7.2â•… Calcium Phosphate-Based Biomaterials HAp and its composites are potential implant materials, particularly for hard tissue replacement applications.44 The most popular bioactive calcium phosphate material
14.7 Biocompatible Ceramics
â•… 227
is HA or HAp [with chemical composition Ca10(PO4)6(OH)2] having a similar mineral composition to bone and teeth. A number of compounds with varying Ca/P ratio, belonging to the CaP family, are relevant to biomedical applications. These include octacalcium phosphate (OCP, Ca/P╯=╯1.33), tri-calcium phosphate (TCP) (Ca/P╯=╯1.5), HAp (Ca/P╯=╯1.67), and tetracalcium phosphate (TTCP, Ca/P╯=╯2). TCP can also exist in two polymorphs: α-TCP and β-TCP. While TCP and HA are the commonly reported phases in CaP-based materials, in vitro or in vivo formation of OCP or other phases have also been reported to a limited extent.45,46 Also, a Ca/P ratio of less than 1.0 is not biomedically important. It needs to be pointed out here that several reports in the literature have emphasized that nonstoichiometric HAp promotes better osteoconduction.47 To date, some attempts are being made to develop HAp-based bioceramic composites in various systems, which include HAp–alumina,48,49 HAp–zirconia,50,51 HAp–bioglass, HAp–HApw, and HAp–TiO2 composites, and so on. In most of the cases, two major common phenomena occur: (1) dissociation of HAp to TCP (α/β) and (2) interfacial reactions between HAp and the reinforcement ceramic phase (e.g., CaZrO3 in the case of HAp–ZrO252). Besides HAp-based ceramic–ceramic composites, limited attempts were made to use metallic reinforcements, such as Ti, to fabricate biocomposites. Ti and TiO2 both are biocompatible and could be useful as the reinforcing phase in CaP-based composites. During pressureless sintering, HAp starts dissociating to other CaP phases (mainly TCP) at temperatures higher than 900°C, depending on the deficiency of calcium53. In a study involving heat treatment of HAp–Ti powder mixtures (10–20â•›wt% Ti) at a sintering temperature of 1100°C in vacuum, it was reported that HAp transforms to a number of CaP phases, such as α-TCP, TTCP54. In addition, a reaction product with a composition of Ca2Ti2O5 was formed. In a different study,55 a specific sintering reaction was reported to take place when a powder compact of HAp, Ti, and TiO2 was sintered at 1100°C for 2 hours in air. It was found that a sintering reaction between TiO2 and HAp took place upon heating above 960°C, and such a reaction led to the formation of the products, such as Ca3PO4 and CaTiO355. In another attempt56 to fabricate HA–(20â•›wt%)Ti composites by the hot pressing route, the existence of Ti metal phase was found to promote the dehydroxylation and decomposition of the HAp ceramic phase into more stable calcium phosphate phases, such as α-TCP and Ca4O(PO4)2, at high temperatures. Zirconia, in contrast, did not induce the decomposition of HAp matrix in the hot pressed HAp–ZrO2 composites.57 In a separate study, Hall and Clifford58 reported that HAp/Al2O3 composites of functionally graded structure can be fabricated by the underwater-shock compaction technique. The composite showed a continuous compositional variation after heat treatment up to 1200°C. Hill and Clifford observed that HAp decomposes at 950°C58 and, accordingly, sintering experiments were carried out using fluorohydroxyapatite (FHA), while replacing HAp. It was found that the biocompatibility of FHA is comparable with pure HAp. Another study59 revealed that HAp in HAp–ZrO2–Al2O3 nanocomposites formed biphasic calcium phosphate (BCP) when hot pressed at 1400°C. BCP is reported to have high biocompatibility. Recently, BCP ceramics have received increased attention as an ideal bone substitute due to their controlled degradability.60 The literature report indicates that it is possible to alter the HAp-to-TCP ratio to
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form BCPs with desired properties.60 In one of the earlier studies, it was observed that 10â•›wt% phosphate glass additions to HAp composites lead to partial dissociation of HAp to TCP61. By varying the glass addition, the resorption of TCP can be controlled. Suchanek and co-workers62 developed HAp–HApw (whisker) composite, which was hot pressed at 1000–1100°C. Despite using whiskers as a toughening agent, the toughness of HAp–HApw (10% addition) could be increased to 1.1â•›MPaâ•›m1/2 under optimal processing conditions. It is, however, possible that toughness could be further improved with increased whisker addition; such experiments have not yet been conducted. Xin et al.63 compared the CaP formation behavior of a few bioceramics (bioglass, glass-ceramics, HAp, α-TCP, and β-TCP) in SBF and in vivo physiological environments. The presence of OCP was observed on all types of bioceramic surfaces in vitro and in vivo, except on β-TCP. One major observation was that CaP formation on bioceramic surfaces is more difficult in vivo than in vitro. Blacik et al.64 performed the weight-bearing study of porous HA/TCP (60/40) ceramics, implanted as intramedullary fixation in segmental bone defects in rabbit bone. They found that the use of these ceramics is limited in their ability to treat load-bearing segmental bone defects, but they did not fail at the early stages of implantation. However, additional internal fixation should be used when immediate mobilization and load bearing are required.
14.8 OUTLOOK In closing, the required property combination for hard tissue replacement biomaterials (ceramics, metals, and polymers and their composites) is summarized in a schematic diagram shown in Figure 14.8. Among various properties, tribological
Processing
In vitro Biomineralization
In vivo Biocompatibility
Microstructure
Bone Analogue Materials
Antimicrobial Properties
Surface Properties
Physical Properties
In vitro Biocompatibility
Figure 14.8â•… Schematic illustration of various physical and biological properties that need to be considered while developing an implant material.
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SBF composition (ionic concentration, serum protein)
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Reduction in COF and less severity in friction
Wear of biomaterials in simulated body fluid (SBF)
Surface property (hydrophilicity/ hydrophobicity) and biological reactivity in SBF
Size, shape, composition, and amount of wear debris particles
Figure 14.9â•… Various factors and aspects of wear of biomaterials in SBF environment.
properties are important and, as has been explained in the early chapters of this book, tribological properties are influenced by the physical properties of the mating materials as well as the tribological environment. It is suggested that SBF or artificial saliva be used in simulating the physiological environment in tribological testing of biomaterials. In this context, the next few chapters discuss the tribological properties of some important biomaterials. It can be reiterated here that, while various in vitro biological test protocols are to be evaluated to assess the potential of various biomaterials prior to clinical applications, the tribological properties in vivo determine long-term performance. Apart from the general factors, such as operating parameters and materials-related factors, some specific factors and aspects that are to be considered while evaluating in vitro tribological properties are schematically illustrated in Figure 14.9. The surface properties as well as the biological reactivity of biomaterials in SBF solution are important; to this end, the SBF composition and the amount of protein dissolved play an important role. Many bioceramic materials, particularly those based on calcium phosphate, exhibit the formation of an apatite layer at the tribological contact and, therefore, the friction and wear behavior would be dependent on the properties as well as the adhesion of such an apatite layer. As mentioned in this overview chapter, biomaterials should not release wear debris particles to a significant extent and this demands extremely high wear resistance. Also, the COF values for biomaterials in articulating surfaces ideally should be low, as higher friction will definitely cause trauma to the patients. The issue related to wear debris particles is a major concern because, depending on the size and composition of the debris particles, they can cause cytotoxicity and genotoxicity of the biological cells in vivo.
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46â•… W. E. Brown, N. Eidelman, and B. Tomazic. Octacalcium phosphate as a precursor in biomineral formation. Adv. Dent. Res. l(2) (1987), 306–313. 47â•… F. Betts, N. C. Blumenthal, and A. S. Posner. Bone mineralization. J. Cryst. Growth 53 (1981), 63–73. 48â•… S. Gautier, E. Champion, and D. B. Assollant. Processing, microstructure and toughness of Al2O3 platelet-reinforced hydroxyapatite. J. Eur. Ceram. Soc. 17 (1997), 1361. 49â•… J. Li, B. Fartash, and L. Hermansson. Hydroxyapatite–alumina composites and bone bonding. Biomaterials 16 (1995), 417–422. 50â•… R. R. Rao and T. S. Kannan. Synthesis and sintering of hydroxyapatite–zirconia composites. Mater. Sci. Eng. C 20 (2002), 187. 51â•… V. V. Silva, F. S. Lameiras, and R. Z. Domínguez. Microstructural and mechanical study of zirconia– hydroxyapatite (ZH) composite ceramics for biomedical applications. Comp. Sci. Tech. 61 (2001), 301–310. 52â•… H. W. Kim, Y. J. Noh, Y. H. Koh, H. E. Kim, and H. M. Kim. Effect of CaF2 on densification and properties of hydroxyapatite–zirconia composites for biomedical applications. Biomaterials 23 (2002), 4113–4121. 53â•… N. Y. Mostafa. Characterization, thermal stability and sintering of hydraxyapatite powders prepared by different routes. Mater. Chem. Phys. 94 (2005), 333. 54â•… Y. Yanga, K. H. Kima, C. M. Agrawal, and J. L. Onga. Interaction of hydroxyapatite–titanium at elevated temperature in vacuum environment. Biomaterials 25 (2004), 2927–2932. 55â•… C. Ergun and R. H. Doremus. Thermal stability of hydroxylapatite-titanium and hydroxylapatitetitania composites. Turkish J. Eng. Env. Sci. 27 (2003), 423–429. 56â•… C. Chu, P. Lin, Y. Dong, X. Xue, J. Zhu, and Z. Yin. Mechanical and biological properties of hydroxyapatite reinforced with 40 vol. % titanium particles for use as hard tissue replacement. J. Mater. Sci. Mater. Med. 13 (2002), 985. 57â•… A. Rapacz-Kmita, A. Slosarczyk, and Z. Paszkiewicz. Mechanical properties of HAp–ZrO2 composites. J. Eur. Ceram. Soc. 26(8) (2006), 1481–1488. 58â•… R. Hill and A. Clifford. Apatite-mullite glass-ceramics. J. Non-Cryst. Sol. 196 (1996), 346–351. 59â•… Y. M. Kong, C. J. Bae, S. H. Lee, H. W. Kim, and H. E. Kim. Improvement in biocompatibility of ZrO2–Al2O3 nano-composite by addition of HA. Biomaterials 26(5) (2005), 509–517. 60â•… I. Manjubala and M. Sivakumar. In-situ synthesis of biphasic calcium phosphate ceramics using microwave irradiation. Mater. Chem. Phys. 71 (2001), 272–278. 61â•… D. C. Tancred, B. A. O. McCormack, and A. J. Carr. A quantitative study of the sintering and mechanical properties of hydroxyapatite/phosphate glass composites. Biomaterials 19(19) (1998), 1735–1743. 62â•… W. Suchanek, M. Yashima, M. Kakihana, and M. Yoshimura. Hydroxyapatite/hydroxyapatite-Whisker composites without sintering additives: Mechanical properties and microstructural evolution. J. Am. Ceram. Soc. 80(11) (1997), 2805–2813. 63â•… R. Xin, Y. Leng, J. Chen, and Q. Zhang. A comparative study of calcium phosphate formation on bioceramics in vitro and in vivo. Biomaterials 26 (2005), 6477–6486. 64â•… C. Balcik, T. Tokdemir, A. Senkoylu, N. Koc, M. Timucin, S. Akin, P. Korkusuz, and F. Korkusuz. Early weight bearing of porous HA/TCP (60/40) ceramics in vivo: A longitudinal study in a segmental bone defect model of rabbit. Acta Biomaterialia 3 (2007), 985–996.
CHAPTER
15
CASE STUDY: POLYMERCERAMIC BIOCOMPOSITES In the context of tribological applications, high-density polyethylene (HDPE) exhibits an extremely low coefficient of friction (COF) of less than 0.1; therefore, it is of interest to evaluate whether ceramic filler additions can improve wear resistance without compromising much on frictional properties. In particular, the friction and wear properties of compression-molded HDPE-based composites with varying ceramic filler content (up to 40â•›vol%) against Al2O3 and ZrO2 under ambient conditions as well as in simulated body fluid (SBF) are summarized in this chapter. In discussing the experimental results, some issues are addressed: (1) whether the improvement in physical properties (hardness, E-modulus) will lead to corresponding improvement in friction and wear properties; (2) whether wear in SBF will provide sufficient lubrication in order to considerably enhance the tribological properties, compared with those under ambient conditions; and (3) whether the generation of wear debris particles will be reduced for various compositionally modified polymer composites, compared with unreinforced HDPE. Wear mechanisms in terms of deformation of the polymer matrix, tribolayer formation, and wear debris generation will also be discussed.
15.1 INTRODUCTION In the overview given in Chapter 14, the general importance of biomaterials is discussed. In particular, musculoskeletal disorders are considered as among the most significant human health problems, costing society more than an estimated $250 billion every year.1 One of the major reasons for the failure of implants is the release of wear debris from the implant surface into surrounding tissue resulting in bone resorption, which ultimately leads to implant loosening. As an illustration of wearinduced damage of a total hip joint replacement, the presence of wear debris particles generated at the acetabular socket–femoral ball head is shown in Figure 15.1. The wear of implants necessitates the patient to undergo revision surgery, for which the success rate is less compared with the first implantation, besides the fact that it is also expensive; 10–20% of orthopedic joints are replaced within 15–20 years and Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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Polymer acetabular cup
Reciprocatory motion of small amplitude Ceramic femoral ball
Wear debris particles
Cytotoxicity or genotoxicity i i off b bone cells ll THR implant
Stem
Figure 15.1â•… Wear of an implant material (total hip joint replacement [THR]) inside the human body (adapted from Ref. 33). The polymeric cup is typically made of ultra-highmolecular-weight polyethylene (UHMWPE). However, in view of the fact that UHMWPE suffers large wear loss and produces coarser wear debris particles, the research effort discussed here investigated the tribological performance of HDPE-based polymer biocompsites against two ceramic materials (alumina, zirconia), which are used as femoral ball head materials.
aseptic loosening accounts for approximately 80% of the revisions. Postmortem studies of the patients who have received total hip or knee replacements reported the accumulation of wear particles in the liver, spleen, or abdominal lymph nodes. The widely used hip joint comprises a femoral head articulating against an ultrahigh-molecular-weight polyethylene (UHMWPE) acetabular cup. From the implant retrieval studies of femoral heads of Co–Cr–Mo alloys, 316L stainless steel (SS), and Ti–6Al–4V alloy, Ti-alloy femoral heads were reported to have the maximum wear, averaging 74.3% against the UHMWPE acetabular component. The enhancement of wear resistance of orthopedic implants is a major concern for the medical community. Such an important aspect of human health care has been a major driving force for research on developing new biomaterials and assessing their physical, tribological, and biological properties. Among various biomaterials, the low coefficient of friction, easy moldability, low density, ability to undergo large inelastic deformation, good biocompatibility, and low cost make polymers candidate materials for various biomedical applications.1–7 However, low stiffness (E-modulus), high wear rate, coarser wear debris particles, and low hardness limit their use in various demanding tribological applications. Hence, a new approach has been to develop polymer-ceramic composite materials.6 In several years of research, the following composite materials are being developed8: hydroxyapatite/HDPE (HAp/HDPE); HAp/UHMWPE; SiO2/silicone rubber; HAp/ ethylene vinyl acetate (HAp/EVA); biphasic calcium phosphate/poly(methyl meth-
15.2 Materials and Experiments
â•… 235
acrylate) (BCP/PMMA); HAp/polylactic acid (HAp/PLA); HAp/polyether ether ketone (HAp/PEEK); bioactive glass/PMMA; and nano-HAp/Poly (hexamethylene adipamide). UHMWPE has been used as a bearing material in total-joint-replacement prostheses for more than 30 years.9 However, the main concern about the use of UHMWPE is regarding adverse biological tissue response, developed by its wear debris.10,11 The debris particles or lumps, formed due to wear of UHMWPE, lead to osteolysis, resulting in implant loosening. This necessitates the revision surgery as well as trauma for patients. Fruh et al.12 developed carbon-fiber-reinforced plastics (CFRPs) and successfully investigated their tribological properties against a hard alumina counterbody. In several studies, a number of researchers proposed that HDPE-based composites could be a suitable alternative to pure UHMWPE materials.13–15 Xue et al.16 successfully developed a composite by reinforcing HDPE/UHMWPE blend with multiwall carbon nanotubes (MWCNTs), and an extremely low wear rate could be achieved (∼12╯×╯10−8â•›mm3/Nâ•›m) when fretted against an austenitic steel counterbody. In another study, considerable improvement in wear resistance has been achieved with UHMWPE/HDPE biocomposite. In another attempt to improve mechanical and tribological properties, HDPE-based biocomposites, reinforced by carbon and Kevlar fibers, were developed and an acetabular cup was prepared by compression molding.17 Substantial research efforts have been invested to develop bioactive composites based on a bioinert HDPE matrix with bioactive HAp ceramic particulates.18–20 Bonfield and co-workers21–24 were the first to develop HAp-reinforced (up to 40â•›vol%) HDPE biomaterial for skeletal applications and coined a trade name, HAPEX™. Wang and co-workers also studied the tribological behavior of HAp/HDPE composite against duplex stainless steel under dry and lubricated conditions.25 However, the results showed that, beyond 10â•›vol% HAp reinforcement, wear properties deteriorate in the presence of low-viscosity liquid. To minimize the generation of wear debris as well as wear loss, a new generation of HDPE–HAp–Al2O3 biocomposites was developed by simultaneous addition of bioinert (Al2O3) and bioactive (HAp) ceramic fillers (up to 40â•›vol%) into a biocompatible HDPE matrix. The processing and the tribological and biocompatibility properties of these composites have been reported elsewhere.26–28 In the development of new composite systems, such as HDPE–HAp–Al2O3 in the present case, it is important to evaluate the performance-limiting properties (e.g., friction and wear) against various mating materials with different chemistry and material properties; this chapter summarizes the friction and wear properties of these composites against two important counterbodies (alumina and zirconia). The results presented here have relevance as far as the application of the HDPE-based composites as acetabular socket material in total hip joint replacement is concerned (see Fig. 15.1).
15.2 MATERIALS AND EXPERIMENTS In developing polymeric biocomposites, HDPE (0.95â•›g/cm3) has been used as a matrix. One of the ceramic fillers, HAp powder (d50╯=╯1.9â•›µm) was synthesized by a widely reported suspension–precipitation route.29 Besides HAp, commercially
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TABLE 15.1.â•… Maximum and Mean Hertzian Contact Pressure at the Time of Initial Contact with Initial Contact Diameter for Different HDPE-Based Composites (Flat Samples) against 9-mm Alumina Ball Counterbody27
Material composition
HDPE HDPE–(20â•›vol%) HAp HDPE–(20â•›vol%) Al2O3 HDPE–(20â•›vol%) HAp –(20â•›vol%) Al2O3 HDPE–(40â•›vol%) Al2O3 HDPE–(40â•›vol%) HAp
Elastic modulus, E (GPa)
Hardness, Hv (MPa)
Maximum Hertzian contact pressure (Initial), MPa
Mean Hertzian contact pressure (Final), MPa
Initial Hertzian contact diameter, µm
0.8╯±â•¯0.03 2.4╯±â•¯0.45
56.5╯±â•¯0.3 129.5╯±â•¯1.1
47.7 97.1
31.8 64.8
633.1 443.5
2.7╯±â•¯0.45
131.1╯±â•¯5.6
105.0
70.0
426.6
6.2╯±â•¯0.91
226╯±â•¯21.5
177.3
118.2
328.3
7.1╯±â•¯0.50
252╯±â•¯9.6
193.7
129.1
314.1
4.2╯±â•¯0.30
176╯±â•¯7.3
137.3
91.5
373.1
available alumina (α-Al2O3, average size 4.8â•›µm, 99.5% pure) was used as another ceramic filler in our composites. Spherical alumina balls (9-mm diameter) and yttria-stabilized tetragonal zirconia polycrystal (Y-TZP) balls (10-mm diameter), both with mirror-finished surfaces (Ra╯∼╯0.05â•›µm), were used as counterbody material. In wear experiments, the compression molded biocomposite pellets (see Table 15.1) were used as flat materials (moving) and the alumina or zirconia balls as the counterbody material. All the experiments were carried out at 35°C with relative humidity (RH) 45╯±â•¯5%. The testing parameters include a normal load of 10â•›N, relative displacement of 80â•›µm, frequency of 10â•›Hz, and testing duration of 100,000 cycles. It can be noted here that the selection of these fretting parameters establishes gross slip contact, that is, the entire contact area undergoes deformation. Based on well-known Hertzian contact mechanics theory, the maximum stress at the start of a fretting test has been calculated and is presented in Table 15.1. The data in Table 15.1 indicate that the tests were conducted with variation in contact stress of ∼48– 194â•›MPa, which is more than the upper end of the commonly experienced contact pressure at various orthopedic joints (hip joint, maximum contact pressure ∼25â•›MPa30). To simulate human physiological environment, the experiments were conducted both in SBF (Hank’s balanced salt solution [composition in wt%]31—8 NaCl, 0.4 KCl, 0.14 CaCl2, 0.2 MgSO4·7H2O, 0.06 KH2PO4, 0.06 Na2HPO4·2H2O, 0.35 NaHCO3, 1.00 glucose [values in grams per liter]) and in an ambient atmosphere.
â•… 237
15.3 Frictional Behavior
15.3 FRICTIONAL BEHAVIOR To illustrate the influence of ceramic fillers (Al2O3/HAp) on frictional properties of HDPE-based composites, Figure 15.2 reveals that against ZrO2, all the biocomposites exhibit COF of less than 0.1 in SBF at 10â•›N load. Similar frictional response has also been observed when the same biocomposites were worn against an Al2O3 ball and a summary of steady-state COF values are summarized in Table 15.2. Overall, the steady-state COF under ambient conditions for the HDPE-based biocomposites are higher than that in SBF. In the present case, the two mating surfaces, that is, hard ceramic ball and soft polymer composite flat, possess completely different mechanical properties. The harder asperities of the alumina ball continuously encounter softer asperities of the polymeric composite and deformation causes lower COF values. Phenomenologically, COF depends on the initial roughness of the harder counterbody. Since the roughness of the alumina ball was very low (Ra╯=╯0.05â•›µm), the contribution of deformation to friction is negligible. Therefore, the observed frictional force originates from the adhesion of two counterbodies. The adhesion, responsible for HDPE friction, originates primarily from the weak bonding forces (hydrogen bonds and van der Waals bonds). During the wear experiments, polymer layers can be transferred onto the hard alumina surface and form a transfer layer on the alumina ball. To illustrate the evidence of the transfer layer on the Al2O3 counterbody, Figure 15.3a,b illustrate the topographical features of worn Al2O3 and ZrO2 balls, respectively, after wear under ambient condition. Energy-dispersive x-ray spectrometry (EDS) compositional analysis (see inset of Fig. 15.3) confirms that the transfer layer on ceramic balls is a polymer-rich film. Therefore, it should be clear that, after a few initial
0.18
0.18
0.15
0.15
0.12
0.12
HDPE HDPE-20 HAp HDPE-20 Al2O3
COF
HDPE-20 HAp-20 Al2O3
0.09
HDPE-40 Al2O3 HDPE-40 HAp
0.09 0.06
0.06
0.03
0.03 0
20,000 40,000 60,000 80,000 100,000
No. of Cycles (a)
0
20,000 40,000 60,000 80,000 100,000
(b)
Figure 15.2â•… Illustration of representative frictional behavior of HDPE-based biocomposites: coefficient of friction (COF) versus number of cycles in air (a) and SBF solution (b). Test conditions: duration, 100,000 cycles; frequency, 10â•›Hz; load, 10â•›N; stroke length, 80â•›µm; counterbody, zirconia ball.28
238
HDPE HDPE–HAp (20â•›vol%) HDPE–Al2O3 (20â•›vol%) HDPE–HAp(20â•›vol%)–Al2O3(20â•›vol%) HDPE–Al2O3 (40â•›vol%) HDPE–HAp (40â•›vol%)
Environment
Counterbodies
Sample designation
0.05 0.08 0.12 0.11 0.15 0.09
Ambient
Zirconia
0.04 0.07 0.06 0.05 0.07 0.07
SBF
Alumina
0.07 0.08 0.13 0.12 0.16 0.09
Ambient
Steady-state COF
0.05 0.05 0.07 0.11 0.12 0.11
SBF 13.8╯±â•¯0.5 7.1╯±â•¯0.4 4.3╯±â•¯0.6 1.8╯±â•¯0.4 0.9╯±â•¯0.5 3.9╯±â•¯0.3
Ambient
SBF 8.1╯±â•¯0.2 4.9╯±â•¯0.4 2.3╯±â•¯0.3 1.1╯±â•¯0.7 0.2╯±â•¯0.4 3.4╯±â•¯0.6
Zirconia
8.9╯±â•¯0.3 4.9╯±â•¯0.2 4.2╯±â•¯0.2 0.7╯±â•¯0.1 0.3╯±â•¯0.1 2.5╯±â•¯0.2
Ambient
SBF 8.3╯±â•¯0.3 4.6╯±â•¯0.2 3.3╯±â•¯0.2 0.6╯±â•¯0.1 0.1╯±â•¯0.1 3.0╯±â•¯0.2
Alumina
Wear rate (×10−6â•›mm3/Nâ•›m)
TABLE 15.2.â•… Summary of Steady-State Friction and Wear Rate of HDPE-Based Biocomposites, Compression Molded at 130°C, 0.5 Hour26
â•… 239
15.3 Frictional Behavior Full Frame C
20 µm (a)
Zr C O
Ca
Al
20 µm (b)
Figure 15.3â•… Representative scanning electron microscopy (SEM) image of worn surface on alumina ball (a) and zirconia ball (b) after wear against HDPE–(40â•›vol%)alumina countersurface. The adhesion of the polymer transfer layer (see also EDS spectra in inset) on the abraded surface can be noticed. EDS compositional analysis (full frame) is shown in inset. Double-pointed arrows indicate the sliding direction.27,28
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fretting cycles, the real tribocontact is established between polymer surface and polymeric transfer layer. Such three-body abrasive wear situations eventually lead to lower COF values in steady-state conditions (Fig. 15.2a). In SBF medium (Fig. 15.2b), COF is lowered in general terms. For composites containing alumina filler, surface energy is increased compared with pure HDPE and, therefore, it can hold a sufficient amount of liquid film on its surface. This may result in considerable decrease of COF values. In SBF medium, the alumina ball surface gets hydrated in the presence of water and an Al(OH)3 layer is formed on its surface. Also, the OH radicals of Al(OH)3 and HAp can easily make hydrogen bonds and adhesion could be increased. As a result, COF increases for HApcontaining composites (40â•›vol% HAp), in spite of SBF lubrication. For composite samples with 20â•›vol% HAp, this adhesion effect is minimized.
15.4 WEAR-RESISTANCE PROPERTIES To quantify the wear damage, Figure 15.4 illustrates representative 2D profilometry scans of the fretted wear surfaces after testing against Al2O3. In general, 2D profiles reveal the smooth nature of a worn surface. The smoothness of the polymer biocom-
10.00
HDPE
ambient
10.00
µm
µm
0.00
0.00
10.00
HDPE-20HAp
10.00
µm
µm
0.00
0.00
−10.00
HDPE-20Al2O3
−10.00
µm
µm
0.00
0.00
10.00
HDPE-20HAp-20Al2O3
10.00
µm
µm
0.00
0.00
10.00
HDPE-40Al2O3
10.00
µm
µm
0.00
0.00
−10.00
HDPE-40HAp
−10.00
µm
µm
0.00
0.00
−10.00
−10.00
HDPE
simulated body fluid
HDPE-20HAp
HDPE-20Al2O3
HDPE-20HAp-20Al2O3
HDPE-40Al2O3
HDPE-40HAp
Figure 15.4â•… Typical 2D surface profiles, recorded using laser surface profilometer of worn surfaces on various HDPE-based biocomposites after testing against alumina ball in air and SBF environments. Test conditions: duration, 100,000 cycles; frequency, 10â•›Hz; load, 10â•›N; stroke length, 80â•›µm; counterbody, 9-mm-diameter alumina ball.27
15.4 Wear-Resistance Properties 8
â•… 241
SBF ambient
Wear depth (µm)
7 6 5 4 3 2
250
300
HDPE-40Al2O3
200
HDPE-20HAp -20Al2O3
150
HDPE-40HAp
100
HDPE-20HAp
50
HDPE
0
HDPE-20Al2O3
1
(a) SBF ambient
Wear depth (µm)
8 7 6 5 4 3 2 1
HDPE-40HAp
250
300
HDPE-40Al2O3
200
HDPE-20HAp -20Al2O3
150
HDPE-20Al2O3
100
HDPE-20HAp
50
HDPE
0
(b)
Figure 15.5â•… Plot of the wear depth of various polymeric biocomposites after wear testing against (a) Al2O3 and (b) ZrO2. Less than 10% variation around the wear depth was experimentally measured. The maximum depth of wear scar is measured using laser surface profilometer. Test conditions: duration, 100,000 cycles; frequency, 10â•›Hz; load, 10â•›N; stroke length 80â•›µm (plotted with data from Refs. 27 and 28).
posite worn surfaces indicates the absence of any severe damage or deep abrasion scratches. This is useful for long-term applications, such as articulating surfaces. Based on wear rate data summarized in Table 15.2 and wear depth data plotted in Figure 15.5, the wear-resistance properties can be discussed. Overall, the wear depth varies in the range of 2–9â•›µm for all the biocomposites, irrespective of
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CHAPTER 15â•… Case Study: Polymer-Ceramic Biocomposites
counterbody. For the Al2O3 counterbody, the wear depth recorded in SBF does not show any noticeable difference from that under ambient conditions. However, a noticeable difference is observed when the ZrO2 counterbody is used. In both cases, however, the wear depth is considerably reduced from 8â•›µm (pure HDPE) to around 4â•›µm or less in the case of HDPE–(20â•›vol%)Al2O3–(20â•›vol%)HAp, irrespective of environment—SBF or ambient conditions. Although any systematic trend could not be observed, the general trend is that the wear depth decreases with hardness. The wear rate data summarized in Table 15.2 were obtained using the same tribometer under the identical combination of operating parameters. The addition of ceramic fillers to HDPE results in a hardness increase, and a systematic decrease in wear rate is recorded for biocomposites (see Table 15.1 and Fig. 15.5). Although the wear rate varies within the same order of magnitude (10−6â•›mm3/Nâ•›m), the highest wear resistance is measured for HDPE–(40â•›vol%)Al2O3 biocomposite, under both ambient and SBF conditions (see Table 15.2). The differences in wear rate among HDPE-based biocomposites will be discussed later on the basis of contact pressure, wear mechanisms, E-modulus, hardness, and so on.
15.5 WEAR MECHANISMS It needs to be pointed out that HDPE–(20â•›vol%)HAp–(20â•›vol%)Al2O3 biocomposite does not show any cytotoxic effect, as tested using L929 mouse fibroblast cells.26 Preliminary short-term in vivo implantation experiments in rabbit femur also reveal a good osseointegration property for HDPE–(20â•›vol%)HAp–(20â•›vol%)Al2O3. In view of its favorable biocompatibility property, the major focus in this chapter is given to discussion of wear mechanisms of HDPE–(20â•›vol%)Al2O3–(20â•›vol%)HAp compared with unreinforced HDPE. The topographical features of the worn surface of unfilled HDPE, fretted under ambient and SBF conditions, are shown in Figure 15.6. Figure 15.6a,c and Figure 15.6b,d display the details of the as-fretted HDPE polymer surface under ambient and SBF conditions, respectively. It is believed that, during repeated loading, plastic strain is accumulated, which results in a ripple type of fatigue cracking, as observed in Figure 15.6. Overall, three types of wear mechanism can be summarized for unreinforced HDPE: (1) severe abrasive wear and formation of deep grooves parallel to the sliding direction; (2) large deformation of the polymer by the contact stress; and (3) presence of fatigue cracks normal to the sliding distances. Figure 15.7 illustrates the topographical features on the worn HDPE–(20â•›vol%) HAp–(20â•›vol%)Al2O3 composite, after wear under ambient and SBF conditions. Importantly, limited wear debris generation and smooth as-fretted surfaces are commonly observed. Figure 15.8b,d reveal that homogeneously dispersed ceramic particles remain embedded and protect the polymer matrix from severe asperity-induced abrasive wear. During fretting wear, the sharp asperities of extremely hard alumina (Hv╯∼╯19â•›GPa) counterbody easily abrade and plow the soft polymers (Hv╯∼╯56â•›MPa), resulting in deep scratches over the fretted surfaces. It may be recalled here that polymers have the unique ability to undergo viscoelastic deformation in a loaded
15.5 Wear Mechanisms
10 µm (a)
â•… 243
20 µm (b) Full Frame C Na
10 µm (c)
Cl
5 µm (d)
Figure 15.6â•… SEM images revealing the overview as well as details of as-fretted/worn surface of pure HDPE in two different environments: air (a,c) and SBF (b,d). Test conditions: duration, 100,000 cycles; frequency, 10â•›Hz; load, 10â•›N; stroke length, 80â•›µm; counterbody, alumina ball. EDS composition analysis of worn surface under SBF conditions is shown in inset. Dotted lines indicate the boundary of wear damage region; double-pointed arrows indicate the sliding direction.27
contact. Additionally, the viscoelastic deformation is also expected to be enhanced at tribocontacts, in the presence of complex stress conditions. Reviewing the fundamentals of contact mechanics,32 the stress field for sphere-on-flat contact contains the nonzero components of shear stress and compressive stress (surface and subsurface) as well as tensile stress at the surface (near the edge of contact). This results in bulk deformation of polymers and causes wavy worn tracks due to variation of stress at the tribocontact. Based on classical Hertzian contact mechanics theory,32 the maximum Hertzian contact pressure for unreinforced HDPE is estimated to be about 47.7â•›MPa (Table 15.1). In contrast, a maximum contact pressure of 193.7â•›MPa for HDPE–(40â•›vol%)Al2O3 is calculated (Table 15.1). Interestingly, in the case of pure polymer, SBF condition does not cause a significant change in wear-resistance properties (see Table 15.2). HDPE, due to its
244â•…
CHAPTER 15â•… Case Study: Polymer-Ceramic Biocomposites P
Full Frame Ca
C O Na
Cl
20 µm (a) Al
Ca
20 µm (b)
Full Frame
P O C
20 µm (c)
20 µm (d)
Figure 15.7â•… SEM images revealing the overview of as-fretted/worn surface of HDPE– (20â•›vol%)HAp–(20â•›vol%)Al2O3 biocomposite in two different environments: air (a) and SBF (b). Also shown are worn surfaces on HDPE–(40â•›vol%)HAp after tribological testing in air (c) and SBF (d). Test conditions: duration, 100,000 cycles; frequency, 10â•›Hz; load, 10â•›N; stroke length, 80â•›µm; counterbody, alumina ball. Full frame EDS spectrum is shown as inset (b,c). Dotted lines indicate the boundary of wear damage region; double-pointed arrows indicate the sliding direction.27
lower surface tension, makes a higher contact angle with water, resulting in a nonadherent water film at the contact interface. This explains the very small effect of lubricating medium on wear properties. Additionally, the experimental observations suggest that the generation of fatigue cracks is also limited under SBF conditions. Overall, mild abrasive wear and limited plastic deformation are major wear mechanisms, irrespective of the lubricating condition. Similar to unreinforced HDPE, lower wear rate and COF are measured in the case of composites under SBF lubricating conditions. Importantly, the results summarized in this chapter indicate that the addition of alumina particles results in greater enhancement of wear properties of composites, compared with HAp addition. The wear data summarized in Table 15.2 demonstrate that the enhancement of the hardness and elastic modulus of bulk composites is reflected in the extremely
15.5 Wear Mechanisms
â•… 245
Full Frame
C
Na
50 µm
20 µm (a)
Cl
(b)
C
5 µm (c)
5 µm (d)
Figure 15.8â•… SEM images revealing the details of as-worn surface of unfilled HDPE in SBF (a,b). The details of the worn debris particles (c) and fatigue crack formation (d) are also presented. Test conditions: duration, 100,000 cycles; frequency, 10â•›Hz; load, 10â•›N; stroke length, 80â•›µm; counterbody, zirconia ball. EDS composition analysis of debris particles is shown in inset. Dotted lines indicate the boundary of wear damage region; double-pointed arrows indicate the sliding direction.28
low wear rate of the bulk composites (approximately of the order of 10−7â•›mm3/Nâ•›m), irrespective of the lubricating medium. Interestingly, maximum hardness (252â•›MPa) and elastic modulus (7.1â•›GPa) were recorded in the case of 40â•›vol% aluminareinforced composites. The mechanical property data correlate well with the lowest wear rate (∼3.1╯×╯10−7 and 1.3╯×╯10−7â•›mm3/Nâ•›m under ambient and SBF conditions, respectively) of HDPE-40 vol% Al2O3 biocomposite. As an illustration of wear damage against ZrO2, Figure 15.8 reveals severe abrasive wear and deformation of the polymer matrix as well as deep grooves along the sliding direction (Fig. 15.8a,b). The size and amount of wear debris particles provide important information on wear of polymers. After testing under dry and ambient conditions, large amounts of flaky and often agglomerated wear debris particles were observed along the rim as well as on the worn surfaces. It is expected
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S4
S4
ambient
SBF
200 µm
200 µm (a)
(c) Al
C
P
Ca
Al
Full Frame
O
O
C
P Ca
5 µm (b)
5 µm (d)
Figure 15.9â•… SEM images revealing the overview of as-worn surface of HDPE–(20â•›vol%) HAp–(20â•›vol%)Al2O3 biocomposite in two different environments: air (a) and SBF (c). Closer and detailed observations of wear tracks are also represented in (b) and (d). Test conditions: duration, 100,000 cycles; frequency, 10â•›Hz; load, 10â•›N; stroke length, 80â•›µm; counterbody, zirconia ball. EDS results of the noticeable tribological features are depicted as insets in the corresponding micrographs. Dotted lines indicate the boundary of wear damage region; double-pointed arrows indicate the sliding direction.28
that, although a similar type of wear mechanism is prevalent under lubricated conditions, the wear debris is mostly washed out from the worn surface. Another common observation of polymer worn surfaces is the generation of parallel arrays of fatigue cracks perpendicular to the sliding direction, irrespective of the tribological environment (see Fig. 15.8d). Figure 15.9a,c illustrates the topographical features on the worn HDPE– (20â•›vol%)HAp–(20â•›vol%)Al2O3 composite after testing against ZrO2 counterbody under ambient and SBF conditions, respectively. The presence of microcracks (often connected) was observed on the composite surface (Fig. 15.9b). Figure 15.9d shows
15.6 CORRELATION AMONG WEAR RESISTANCE, WEAR MECHANISMS
â•… 247
that asperity-induced scratches could not propagate as the stiff alumina particles often restrict the deformation of the bulk polymer matrix during sliding motion. It can be recalled here that the combined addition of 20â•›vol% Al2O3 along with 20â•›vol% HAp results in improved hardness (∼226â•›MPa) and elastic modulus (∼6.2â•›GPa), as well as low wear rate (1.78╯×╯10−6â•›mm3/Nâ•›m under ambient and 1.08╯×╯10−6â•›mm3/Nâ•›m in SBF).
15.6 CORRELATION AMONG WEAR RESISTANCE, WEAR MECHANISMS, MATERIAL PROPERTIES, AND CONTACT PRESSURE From the preceding discussion, severe abrasive wear, formation of large grooves, and generation of fatigue cracks can be identified as the important wear mechanisms for unreinforced HDPE. Due to repeated loading and unloading experienced at the tribocontact, the polymer surface experiences a large number of stress cycles. Subsequently, the accumulation of large viscoelastic strain resulted in parallel arrays of cracks, perpendicular to the sliding direction (Fig. 15.8d). However, the occurrence of fatigue wear and cracking was not observed in any of the composites due to improved mechanical properties. It is clear from Table 15.1 that, with increasing addition of ceramic fillers (Al2O3/HAp), the initial Hertzian contact pressure is increased and that wear rate decreases. It can be pointed out here that, for material with higher hardness and higher elastic modulus, the initial Hertzian contact area will be less due to less elastic deformation. If the contact area is less, then for a given load, the contact stress will be higher. Therefore, in spite of having higher contact stress, the materials exhibit better wear resistance properties due to higher hardness values. Softer materials, such as HDPE, try to deform elastically or plastically (depending upon surface roughness, etc.) to reduce the contact stress. As far as the influence of ceramic filler addition is concerned, a reduction in wear rate with increasing amount of filler volume fraction (up to 40â•›vol%) is commonly observed. Importantly, the investigation discussed here revealed much less wear debris formation in the case of HDPE-reinforced composites compared with unfilled HDPE. Additionally, in terms of wear rate, a substantially lower wear rate is measured for the composites. Such improvements in wear resistance can be correlated with an increase in the elastic modulus and/or hardness. Overall, less abrasive wear and limited plastic deformation are the major wear mechanisms in the composites. The simultaneous addition of 20â•›vol% HAp and 20â•›vol% Al2O3 to the HDPE matrix significantly improves its mechanical and tribological properties. In the presence of the harder ceramic phase, the deformation of the polymer matrix is restricted. Observing the wear data measured for zirconia counterbody, the composites reinforced with 40â•›vol% ceramic fillers (20â•›vol% HAp and 20â•›vol% Al2O3) as well as the composite containing 40â•›vol% alumina exhibit lower specific wear rate (∼1.08╯×╯10−6–1.78╯×╯10−6╯N╯m/mm3 and ∼0.2╯×╯10−6– 0.9╯×╯10−6╯N╯m/mm3, respectively [see Table 15.2]).
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15.7 CONCLUDING REMARKS The following points can be summarized from the discussion in this chapter: a. The addition of ceramic fillers to HDPE does not have any significant influence on degradation in frictional properties under both ambient and SBF conditions. Broadly, a decrease in COF is measured during wear in SBF, irrespective of composite composition. Also, three-body abrasion resulted in lower COF (∼0.15) even under dry ambient conditions. b. A better wear resistance with increase in hardness has been recorded, in both ambient and SBF environments, and the wear rate varies on the order of 10−6–10−7â•›mm3/Nâ•›m. c. The worn surface observations reveal that severe abrasive wear, plastic deformation leading to deep grooves, formation of a transfer layer on the ball counterbody followed by spalling, and flaky wear debris are the major wear mechanisms observed in the case of unfilled HDPE. However, under similar operating conditions, the presence of Al2O3/HAp considerably reduces the severity of abrasion and deformation of the HDPE matrix. The generation of much less wear debris is observed for Al2O3/HAp-reinforced HDPE composites. d. On the basis of the possibility of achieving better mechanical properties as well as optimal combination of COF and wear resistance, the results presented in this chapter demonstrate that both HDPE–(40â•›vol%)Al2O3 and HDPE– (20â•›vol%)Al2O3–(20â•›vol%)HAp biocomposites can be considered as potential materials for orthopedic applications. The biocompatibility study revealing noncytotoxic response against fibroblast cells has already been reported else-
Enhanced E-modulus and Hertzian contact damage
Three-body abrasion (polymer transfer to counterbody) and low COF (0.1-0.15)
Polymer-ceramic biocomposites
- Less abrasion and cracking - Improved wear resistance
Figure 15.10â•… Schematic illustration of various key points of the aspects of the tribological properties of polymer–ceramic biocomposites, based on the results of HDPE–HAp–Al2O3 discussed in this chapter.
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where,26 and preliminary osseointegration tests in a rabbit model confirm good in vivo biocompatibility properties of HDPE–(20â•›vol%)Al2O3–(20â•›vol%)HAp biocomposites. Various important aspects of the tribological properties of polymer–ceramic composites, as expected on the basis of the results obtained for HDPE–Al2O3–HAp are schematically shown in Figure 15.10. While improved contact damage resistance can be expected due to increase in E-modulus with ceramic reinforcement, the severity of the abrasion as well as better wear resistance can be achieved by carefully selecting the type and amount of ceramic particulates. Another important aspect is that the beneficial frictional property (i.e., low COF) is not compromised to any significant extent, while enhancing physical and wear-resistance properties with ceramic reinforcement in polymer matrix. Similar studies in other polymer systems in future will confirm the general predictions stated here.
REFERENCES ╇ 1â•… A. A. Edidin and S. M. Kurtz. Influence of mechanical behavior on the wear of 4 clinically relevant polymeric biomaterials in a hip simulator. J. Arthroplasty 15 (2000), 321–331. ╇ 2â•… M. Long and H. J. Rack. Titanium alloys in total joint replacement—A materials science perspective. Biomaterials 19 (1998), 1621–1639. ╇ 3â•… C. A. Holding, D. M. Findlay, R. Stamenkov, S. D. Neale, H. Lucas, A. S. S. K. Dharmapatni, S. A. Callary, K. R. Shrestha, G. J. Atkins, D. W. Howie, and D. R. Haynes. The correlation of RANK, RANKL and TNFa expression with bone loss volume and polyethylene wear debris around hip implants. Biomaterials 27 (2006), 5212–5219. ╇ 4â•… A. Sargeant and T. Goswami. Hip implants: Paper V. Physiological effects. Mater. Design 27 (2006), 287–307. ╇ 5â•… J. H. Dumbleton and M. T. Manley. Metal-on-metal total hip replacement: What does the literature say? J. Arthroplasty 20 (2005), 174–188. ╇ 6â•… S. M. Lee. Orthopedic Composites: International Encyclopedia of Composites, Vol. 4. VCH Publishers, New York, 1991, 74–86. ╇ 7â•… S. Ramakrishna, J. Mayer, E. Wintermantel, and K. W. Leong. Biomedical applications of polymercomposite materials: A review. Comp. Sci. Technol. 61 (2001), 1189–1224. ╇ 8â•… S. Nath and B. Basu. Designing biomaterials for hard tissue replacement. J. Kor. Cer. Soc. 45(1) (2008), 1–29. ╇ 9â•… S. Li and A. H. Burstein. Current concepts review: Ultra-high molecular weight polyethylene. J. Bone Joint Surg. 76A (1994), 1080–1090. 10â•… H. C. Amstutz, P. Campbell, N. Kossovsky, and I. C. Clarke. Mechanism and clinical significance of wear debris-induced osteolysis. Clin. Orthop. Relat. Res. 276 (1992), 18. 11â•… T. P. Schmalzried, M. Jasty, and W. H. Harris. Prosthetic bone loss in total hip joint arthroplasty. J. Bone Joint Surg. 74A (1992), 849–863. 12â•… H. J. Fruh and G. Willimanm. Tribological investigations of the wear couple alumina-CFRP for total hip replacement. Biomaterials 19 (1998), 1145–1150. 13â•… Q. Guo and W. Luo. Mechanisms of fretting wear resistance in terms of material structures for unfilled engineering polymers. Wear 249 (2002), 924–931. 14â•… B. A. Kehler, N. P. Baker, D. H. Lee, C. J. Maggiore, M. Nastasi, J. R. Tesmer, K. C. Walter, Y. Nakamura, and B. M. Ullrich. Tribological behavior of high-density polyethylene in dry sliding contact with ion-implanted Co-Cr-Mo. Surf. Coatings Technol. 114 (1999), 19–28. 15â•… J. C. Anderson. High density and ultra-high molecular weight polyethenes: Their wear properties and bearing applications. Tribology Int. 15 (1982), 43–47.
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16â•… Y. Xue, W. Wu, O. Jacobs, and B. S. Del. Tribological behaviour of UHMWPE/HDPE blends reinforced with multi-wall carbon nanotubes. Poly Testing 25 (2006), 221–229. 17â•… S. K. Roy Chowdhury, A. Mishra, B. Pradhan, and D. Saha. Wear characteristic and biocompatibility of some polymer composite acetabular cups. Wear 256 (2004), 1026–1036. 18â•… W. Bonfield. Hydroxyapatite-reinforced polyethylene as an analogous material for bone replacement. In: Ducheyne P, Lemons JE, editors. Bioceramics: Materials characteristics versus in vivo behavior. Ann. N Y Acad Sci. 253 (1988), 173–177. 19â•… W. Bonfield, M. D. Grynpas, A. E. Tully, J. Bowman, and J. Abram. Hydroxyapatite reinforced polyethylene—A mechanically compatible implant. Biomaterials 2 (1981), 185–189. 20â•… J. Huang, L. D. Silvio, M. Wang, K. E. Tanner, and W. Bonfield. In vitro mechanical and biological assessment of hydroxyapatite reinforced polyethylene composite. J. Mater. Sci. Mater. Med. 8 (1997), 775–779. 21â•… M. Wang, D. Porter, and W. Bonfield. Processing, characterisation, and evaluation of hydroxyapatite reinforced polyethylene composites. Br. Ceram. Trans. 93 (1994), 91–95. 22â•… F. J. Guild and W. Bonfield. Predictive modelling of hydroxyapatite–polyethylene composites. Biomaterials 14 (1993), 985–993. 23â•… M. Wang, S. Deb, and W. Bonfield. Chemically coupled hydroxyapatite-polyethylene composite: Processing and characterization. Mater. Lett. 44 (2000), 119–124. 24â•… M. Wang and W. Bonfield. Chemically coupled hydroxyapatite-polyethylene composite: Structure and properties. Biomaterials 22 (2001), 1311–1320. 25â•… M. Wang, M. Chandrasekaran, and W. Bonfield. Friction and wear of hydroxyapatite reinforced high density polyethylene against the stainless steel counterface. J. Mater. Sci. Mater. Med. 13 (2002), 607–611. 26â•… S. Nath, S. Bodhak, and B. Basu. HDPE-Al2O3-HAp composites for biomedical applications: Processing and characterization. J. Biomed. Mater. Res. B Appl. Biomater. 88B (2009), 1–11. 27â•… S. Bodhak, S. Nath, and B. Basu. Friction and wear properties of novel HDPE-HAp-Al2O3 biocomposites against alumina counterface. J. Biomater. Appl. 23 (2009), 407–433. 28â•… S. Bodhak, S. Nath, and B. Basu. Fretting wear properties of hydroxyapatite, alumina containing high density polyethylene biocomposites against zirconia. J. Biomed. Mater. Res. A 85 (2008), 83–98. 29â•… L. L. Hench and J. Wilson. An Introduction to Bioceramics. World Scientific, London, UK, 1993. 30â•… M. Kumagai, Y. H. Kim, N. Inoue, E. Genda, K. Hua, B. T. L. Liong, T. Koo, and Y. Chao. 3-D dynamic hip contact pressure distribution in daily activities. 2003 Summer Bioengineering Conference, June 25-29, Sonesta Beach Resort in Key Biscayne, Florida, 2003, 53–54. 31â•… J. Hanks. Hanks’ balanced salt solution and pH control. Method Cell Sci. 1 (1975), 3–4. 32â•… B. Bhushan. Principles and Applications of Tribology. John Wiley & Sons, New York, 1999, 201–202. 33â•… M. Peters, H. Hemptenmacher, J. Kumpfert, and C. Leyens. Titanium and Titanium Alloys, C. Leyens and M. Peters (Eds.). Wiley-VCH, 2003, 1–57.
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16
CASE STUDY: NATURAL TOOTH AND DENTAL RESTORATIVE MATERIALS For the development of new dental restorative materials, it is essential to evaluate and understand the structure-property relationships of the human tooth. This chapter first discusses microstructure–property correlation of three major structural parts of the human tooth: enamel; dentin; and the dentin–enamel junction (DEJ). To evaluate the tribological properties of the human tooth, fretting wear test results will be discussed to illustrate how the wear of teeth is dominated by fretting fatigue and adhesive wear. Tooth wear involves the formation of the oxidized calcium phosphate based compounds and their subsequent transfer from tooth to alumina surface. In the context of dental restorative materials, various glass-ceramic (GC) materials are being developed. Considering the potential of the mica-based GCs as dental implants, an understanding of their wear behavior in the oral environment is important. The later part of this chapter discusses studies on K2O–B2O3–Al2O3– SiO2–MgO–F glass-ceramics containing about 70% crystals, heat-treated at 1040°C for 12 hours and subjected to fretting against a steel ball in an artificial saliva (AS) environment. It will be shown that the tribochemical layer formed in AS medium undergoes brittle fracture, whereas mica crystal pullout was the dominant mechanism in dry conditions.
16.1 INTRODUCTION Human teeth act as a mechanical device during masticatory processes such as cutting, tearing, and chewing of food particles.1 It is the only mineralized organ whose location is part internal and part external to the human body. The tooth is a functionally graded composite material with mineralized matrix and organic reinforcements. It is composed of three basic structural parts, namely enamel, dentin, and the dentin–enamel junction (DEJ). Different regions along with the principal structural parts of the human tooth are shown in Figure 16.1. The anatomical crowns of teeth are covered by the dental enamel, which is the hardest tissue in the body. Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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Buccal Region
Axial Direction
Lingual Region
Enamel Dentin
DEJ
Pulp Cavity Cervical Margin
Figure 16.1â•… Schematic of a tooth section cut perpendicularly to the occlusal surface along the long axis showing the important parts.15
The enamel is composed of 92–96% inorganic matter or mineral phase, 1–2% of organic material, and 3–4% of water by weight.2 The mineral phase consists primarily of calcium phosphate salts in the form of large hexagonal carbonated and defective hydroxyapatite (HAp) crystals.3 Tooth enamel has a unique microstructure consisting of aligned prisms or rods, running approximately perpendicular from the dentine–enamel junction toward the tooth surface. The ideal structure of these rods is keyhole-like, having an average width of about 5â•›mm.4 The DEJ can be considered as a biological interface between the enamel part and the dentin. It is a unique junction region between highly mineralized tissues of different embryogenic origins, matrix composition, and physical properties. Enamel covers the underlying dentin, which is a hydrated biological composite, consisting of 70% inorganic material, 18% organic matrix, and 12% water (wt%).2 Dentin contains dentinal tubules, extending through its entire thickness. The tubules are surrounded by highly mineralized cylinders of peritubular dentin in the crown, and the latter are separated by intertubular dentin.5 It has been reported that the number of dental tubules can vary from 4900 to 57,000 per square millimeter cross-sectional area. The number of these tubules progressively decreases from the crown toward the apical direction.6 To prepare clinical tooth replacements and develop new artificial dental restorative materials, it is important to understand the mechanical properties of the human tooth.7 Experimental results show that both enamel and dentin are fractured in a brittle manner.7 The fracture mode of enamel is highly anisotropic, while that of dentin is less anisotropic.7 Advanced characterization techniques, such as the Moiré fringes technique5 and microtensile testing,2 have been used to record the mechanical response of the different anatomical parts (enamel, dentins, and DEJ) of the human tooth. However, the relationship between basic mechanical properties, such as hardness, and the microstructural composition of various structural parts of the human tooth is lacking. One part of this chapter attempts to fill this lacuna. Another important property to be critically considered for dental restorative materials is their tribological behavior. The wear of human teeth is a natural and
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unavoidable phenomenon, resulting from the physiological and pathological functions in the oral cavity. Excessive wear may result in a lack of perfect contact between opposite teeth, loss in the efficiency of the masticatory system, and obliteration of chewing surfaces. Therefore, understanding the friction and wear behavior of human teeth would be helpful for the clinical treatment of teeth and also for the development of new dental restorative materials. The inherent anisotropy of the human tooth, in terms of mineral concentration gradient, and the resulting mechanical property variation in the enamel and dentin also influence their tribological behavior. Review of the existing literature reveals that limited work has been carried out to evaluate and understand the friction and wear properties of the human tooth. Burak et al.8 demonstrated that, at lower loads, enamel, with its high mineral enamel content, exhibits relatively low wear rates as compared with dentine, which contains a higher amount of organic phase. However, the brittle nature of enamel leads to relatively high wear rates at higher loads, whereas the connective tissue matrix of dentin makes it less susceptible to fracture under tribological conditions. Al-Hiyasat et al.9 carried out an in vitro study to compare the wear behavior of enamel against aluminous porcelain, bonded porcelain, low-fusing hydrothermal ceramic, feldspathic machinable ceramic, and cast gold. It was observed that gold caused significantly less enamel wear than the tested ceramic mating body, and the amount of enamel wear in the aluminous and bonded porcelain groups was significantly higher than that with the hydrothermal and machinable ceramics. However, there was no significant difference between the amount of enamel wear produced by the aluminous and bonded porcelains or between that produced by the hydrothermal and machinable ceramics. Seghi et al.10 performed a similar study on wear and found that Dicor GCs and Dicor coated with shading porcelain causes the lowest wear of enamel. Since 1990, different biocompatible ceramic and GC materials have been researched for biomedical and dental applications.11 The dental ceramics and GCs are especially considered for application such as restorative materials and supporting structures. Considering the increasing clinical problems associated with the longterm wear of dentin, some investigators attempted to study the wear mechanisms of human teeth and GC materials.12–15 While substantial work has been carried out in the microstructural evolution of various GC materials,16–22 systematic research in understanding the dominant wear mechanisms is rather limited for this important class of material. In an investigation of dry unlubricated sliding behavior, Park et al.23 demonstrated the dominating influence of the amount and orientation of wollastonite phase on the hardness and wear of apatite–wollastanite (A-W) GCs. The wear resistance of the material decreased with the decreasing amount of wollastonite. In a different work, Xiao et al. observed that the frictional properties of CaO–MgO–Al2O3–SiO2 (CMAS) GC were related to the contact pressure and sliding speed, while plastic deformation and recrystallization have been identified as dominant wear mechanisms during dry sliding tests.24 The tribological study with Dicor GC against alumina showed higher coefficient of friction (COF) of 0.8 in dry conditions.25 Based on Hertzian contact studies, it was concluded that the wear of mica-containing Macor
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GC is controlled by its short-crack toughness as well as by the size and volume fraction of mica plates.25 In a tribological study of feldspathic porcelain ceramics against silicon nitride in simulated oral conditions, Yu et al.26 illustrated the dominant effect of load compared with sliding speed and the frequency of oscillation. The reduction in friction and wear of cordierite GC resulted from the formation of thin gel-like reaction layers as a result of ion exchange in hydrochloric acid, while the network dissolution process led to increased wear and friction in caustic soda solutions.27 Nagarajan and co-workers28 reported a clear transition of wear from localized fracture at low load to contact or spallation mode at high loads, during sliding of mica-containing GC against alumina under distilled water lubrication. It has been further argued that microcrack-induced fracture occurs either along the weak mica– glass interface or mica cleavage planes.28 In another study29 of the sliding of glass composites against alumina in distilled water, it was found that wear predominantly occurs by the formation, delamination, and dissolution of the tribochemical layer and products. In more recent work,30 the development of some unusual spherulitic–dendritic crystals as well as their in vitro dissolution properties in K2O–B2O3–Al2O3–SiO2– MgO–F glass-ceramics have been reported. In subsequent work,31 it was shown how the variation in fluorine content over 1–4% influences the crystallization as well as microstructure development and mechanical properties (hardness, strength) in K2O– B2O3–Al2O3–SiO2–MgO–F glass-ceramics. Particularly, these GC materials showed higher hardness and comparable elastic modulus values than those observed in any other dental restorative materials.21,32 To explore the potential of this compositionally varied GC system to be used as dental restorative material, a clear understanding of its tribological behavior is essential. In this perspective, a part of this chapter reports the tribological behavior of the K2O–B2O3–Al2O3–SiO2–MgO–F glass-ceramic material with 70% mica crystal content, when subjected to fretting against steel. A major part of the work is focused on showing the influence of environment on the mechanisms of material removal, when tested under dry conditions and in AS medium.
16.2 MATERIALS AND METHODS 16.2.1â•… Preparation of Human Tooth Material For the study under discussion, six freshly extracted human mandibular third molar teeth were taken from six persons of varying age and food habits. All the teeth were completely free from caries (i.e., devoid of any observable cavities) and did not have any visible crack at the surface. One of these teeth was cut at right angles to the long axis, as shown in Figure 16.1. Most of the region in the root and pulp cavity was cleared and it was mounted in a die using a chemically reacting resin complex. It was polished using diamond slurry, of particle sizes 9, 6, 3, 1, and 0.25â•›µm.
16.4 Production and Characterization of Glass-Ceramics
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16.3 TRIBOLOGICAL TESTS ON TOOTH MATERIAL For the study on tribology, a specially developed tooth ball was rubbed against a sintered alumina base plate (99% theoretical density and hardness ∼ 18â•›GPa). For the preparation of the tooth balls, one groove each of 4â•›mm diameter and 4â•›mm depth was made in 10-mm-diameter polyethylene balls reinforced with steel particles. Small conical sections were cut from the remaining five molar teeth using a diamond blade cutter in such a way that the upper part contained only enamel and the lower part could be inserted and fixed into the groove of the polymer balls with the help of organic glue. The tooth sections were made carefully to have almost the same roundness as the polymer balls, to ensure point contact during fretting. The alumina base plate was about 8â•›mm╯×╯4â•›mm╯×╯1.5â•›mm and smoothly polished using the diamond slurry (particle sizes of 9, 6, 3, 1, and 0.25â•›µm). Alumina was selected as a mating counterbody, because alumina is one of the potential materials for dental restoration. Also, alumina, being much harder than natural human teeth, would not undergo much wear and instead, the tooth would wear extensively so that severe tooth wear could be studied in the planned experiments. The fretting tests were conducted under unlubricated condition at a load of 1 N, frequency of 8â•›Hz, and linear stroke length of 30â•›µm for five different test durations (2000, 4000, 6000, 8000, and 10,000 cycles). Relative humidity of 45╯±â•¯5% and temperature of 30╯±â•¯3°C were maintained during all the experiments. Such close control over temperature and humidity in the testing environment was possible by placing the test assembly in a humidity-controlled chamber. The testing under dry and unlubricated conditions was intentionally performed to study the severe wear of natural tooth material, which would not occur if tested in AS. The fretting parameters are selected so as to simulate the actual tribological conditions that prevail during mastication phenomenon.
16.4 PRODUCTION AND CHARACTERIZATION OF GLASS-CERAMICS The precursor materials used to produce the base glass in the K2O–B2O3–Al2O3– SiO2–MgO–F system. include high purity optical grade Quartz Flour (Sipur A1 Bremthaler Quartzitwerk, Germany), aluminum nitrate nona hydrate (Riedel-deHahn AG, Germany), magnesuim hydroxide carbonate (Merck, Germany), potassium nitrate (Merck KGaA, Germany), boric acid (Merck KGaA), and MgF2 (Merck KGaA). The glass batch of appropriate composition was mixed by Eirich Mixer (Germany) for 5 minutes, followed by melting of the glass composition, carried out in a platinum crucible at 1550°C for 2 hours using an electrical furnace. The glass melts were cast into a mild steel mold to obtain plates and then subjected to annealing for 2 hours in the temperature range of 600–650°C. The analysis of the composition of the base glass was done by inductively coupled plasma atomic emission spectroscopy (ICP-AES) (spectroflame modula FTM 08, Germany). Two-stage heat treatment was done to crystallize the glass and thus convert it into GC. Nucleation
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TABLE 16.1â•… Mechanical and Physical Properties of the K2O–B2O3–Al2O3–SiO2–MgO–F Glass-Ceramic (GC)31
Density (g/cm3) 3-point flexural strength (MPa) Elastic modulus (GPa) Vickers hardness (GPa)
2.83 80.6╯±â•¯7.7 69.7╯±â•¯2.9 6.4╯±â•¯1.2
was carried out in the temperature range of 750–850°C for 6 hours, followed by crystallization at 1040°C for a holding time of 12 hours. A number of studies investigated the wear of human teeth and glass-ceramic materials in view of their potential use as dental restorative materials.23–27,29,33,34 Our earlier study34 illustrated that the combination of this processing scheme results in a GC material with an optimal combination of crystal volume fraction and mechanical properties. The heating and cooling rates were maintained at 180°C/h throughout the heat-treatment cycle. The sample surfaces were polished using standard metallographic papers and finally with diamond paste (from 9â•›µm to 0.25â•›µm). The final surface roughness of the samples before wear testing was measured using a laser surface profilometer (PGK-120, Mahr GmbH, Gottingen, Germany) and was found to be on the order of 0.05â•›µm. The spatial resolution and vertical resolution of the profilometer are 0.1â•›µm and 0.5â•›nm, respectively (claimed by the manufacturer). The polished surface was etched with 12% hydrofluoric acid (HF) with time varying between 1 and 3 minutes at different zones of the sample for revealing the crystal morphology in the bulk GC sample. The conventional point counting method using scanning electron microscopy (SEM) images of the etched surfaces showed a large amount (70 vol%) of mica crystals in the investigated GC material. Various mechanical properties of the GC material are listed in Table 16.1. More details about the properties of the GC can be found in our paper.34
16.5 WEAR EXPERIMENTS ON GLASS-CERAMICS Polished GC bars of 25╯×╯5╯×╯4â•›mm3 were used as flat (moving) materials and 10-mm-diameter balls of commercial SAE 52100 grade steel (hardness ∼7â•›GPa) were chosen as counterbody (stationary) materials. The ideal mating material would have been natural human tooth. However, steel counterbody was chosen based on two reasons: (1) similarity in hardness with that of the developed GC or dental prosthetic materials (Ni–Cr alloy, ceramic, etc.) and (2) the difficulties in the preparation of ball samples with human tooth material. Similar studies on the wear of dental materials against steel were also previously reported.33 Prior to wear testing, both the GC and steel samples were ultrasonically cleaned using acetone. The fretting tests were carried out in ambient conditions in air without the use of any lubricant (dry conditions), and using AS medium with selected parameters of 1-N load, 8-Hz oscillation frequency, and 100-µm (gross slip regime) linear stroke length. The composition of AS is given in Table 16.2. The experiments were carried out for different test durations of 5000, 10,000, 50,000, and 100,000 cycles.
â•… 257
16.6 Microstructure and Hardness of Human Tooth Material
TABLE 16.2â•… Typical Composition of Artificial Saliva and All the Precursor Materials Are of 99% Purity or Above (AR grade)34
Material
Amount (g) per 1â•›L of distilled water
NaCl KCl CaCl2 2H2O NaH2PO4 2H2O Na2S 9H2O Urea
0.4 0.4 0.795 0.78 0.005 1
100 µm
100 µm
(a)
(b) E
D
DEJ
Cracks 10 µm (c)
Figure 16.2â•… SEM micrographs showing (a) enamel prisms running trough the thickness of the tooth, (b) dentinal tubules embedded in mineralized matrix, and (c) dentin–enamel junction (DEJ) in between the enamel (left side) and dentin (right side) appearing to be a zone of nearly 40â•›µm.15
16.6 MICROSTRUCTURE AND HARDNESS OF HUMAN TOOTH MATERIAL As stated earlier, a human tooth sample was cut along the longitudinal axis, and SEM observations and hardness measurements were performed on the polished cross-sectioned sample. Figure 16.2a–c present the representative SEM images,
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showing the microstructure of enamel, dentin, and DEJ, respectively. It can be seen from Figure 16.2a that, although most of the enamel prisms are almost parallel to each other and run through the thickness of the entire enamel from the DEJ to the occlusal surface of the tooth, some of them are almost interwoven with each other. In Figure 16.2b, the dentinal tubules appear as darker cavities in a matrix, which appears to consist of similar HAp prisms. In Figure 16.2c, the DEJ, separating the enamel (left) from the dentin (right), appears as a distinct zone rather than a single line. A closer look at Figure 16.2c shows that some cracks in the enamel part are effectively stopped at the DEJ and are not allowed to enter into the dentin part. This observation agrees well with the earlier findings of Xu et al.,7 who found that cracks in the enamel propagating toward the dentin were stopped at the DEJ and, therefore, could not penetrate into the dentin. This observation, combined with the resilient nature of the underlying dentin, supports the mechanical integrity of the enamel by preventing its fracture during masticatory action. The variation of hardness at different locations of the tooth and the corresponding composition gradient along the line parallel to indentations is shown in Figure 16.3a,b, respectively. The microstructural regions that have been studied are shown in Figure 16.3a. It can be seen that the hardness is highest at the outermost surface of the enamel (around 3.5â•›GPa) and it decreases with increasing depth inside the tooth. It is also seen that, at a distance of 100–600â•›µm from the DEJ, the hardness of the enamel is nearly constant with marginal fluctuation (ranging from 2 to 2.5â•›GPa). It may be mentioned here that earlier researchers also used different characterization methods to study the anisotropic nature of teeth. Meredith et al.12 have shown that Knoop hardness of enamel decreases with an increase in distance from the enamel surface at a linear rate of approximately 0.023â•›HK/µm and the hardness of dentin increased with distance from the amelodentinal junction (ADJ). Zheng et al.13 have shown that hardness value differs not only between enamel and dentin, but also between different occlusal layers of the enamel for a tooth. They reported that the hardness decrease is 17% from the outer layer of enamel to the DEJ and a further 77% from the DEJ to the dentin. Habelitz et al.4 used a nanoindentation technique to demonstrate that higher elastic modulus (E-modulus) and hardness can be measured in enamel when indentations are taken parallel to the enamel rod orientation. They recorded an average E-modulus of 87.5 (±2.2) and 72.7 (±4.5) GPa and an average hardness of 3.9 (±0.3) and 3.8 (±0.4) GPa in directions parallel and perpendicular to the enamel rods, respectively. It has been seen that hardness drops sharply to a very low value (less than 1â•›GPa) at the DEJ and again increases in the dentin region (nearly 1.5â•›GPa). At a distance of about 200â•›µm from DEJ, dentin hardness decreases below 1â•›GPa and remains nearly constant afterward. Earlier, Wang and Weiner5 found that the dentin adjacent to the DEJ had a low microhardness that increases rapidly to a peak value and then decreases slowly toward the pulp cavity. Meredith et al.12 showed also that both hardness and elastic modulus of dentin increase with distance from the ADJ. To determine the chemical composition of different regions, the SEM energydispersive x-ray spectrometry (SEM-EDS) line scan data in terms of Ca, P, and O content have been plotted in Figure 16.3c. It is quite obvious that enamel contains higher amounts of Ca, P, and O than dentin and that the concentration of these
Enamel
Indentation Direction
Line Scan Direction
DEJ
Dentin
200 µm (a) 3.5 Enamel
3.0
DEJ
Dentin
Hardness (GPa)
2.5 2.0 1.5 1.0 0.5 1000
0 500 500 Distance from DEJ (µm) (b) DEJ
Enamel
1000
1500
Detin
Intensity (arbitary unit)
O P Ca
2000
1000
0 1000 2000 Distance from DEJ (µm) (c)
3000
Figure 16.3â•… (a) SEM image showing Vickers hardness indentations taken along the line from enamel to dentin, (b) hardness values measured against the distance from DEJ, and (c) composition variation measured along the line parallel to the indentation line.15
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elements suddenly decreases at the DEJ. The significant fall in mechanical properties therefore appears to have a strong correlation with reduction in the amount of calcium phosphate derivatives, that is, with mineral concentration from the enamel surface to the DEJ. In earlier research, Cuy et al.3 observed that both hardness and elastic modulus decreased from 4.6 to 3.4â•›GPa (26% decrease) and from 91.1 to 66.2â•›GPa (27% decrease), respectively, on going from the enamel surface to the enamel–dentin junction. Wong et al.14 observed an increasing mineral concentration from the ADJ to the external surface of the enamel, ranging from 0.08 to 0.60â•›mg/ cm3â•›mm with an average of 0.366â•›mg/cm3â•›mm. However, the aforementioned variation in composition, as recorded in our work, does not appear to explain the hardness variation completely. It is possible that the microstructural features of different regions, such as prism orientation in the enamel part as well as the density and orientation of tubules in the dentin part, also play an active role in determining the hardness at those locations. The observed anisotropy in the mechanical properties of enamel is possibly due to the anisotropy and alignment of fiberlike apatite crystals within the rods and to the composite architecture of enamel, as has been found by earlier researchers.4,7 It has been observed by Habelitz and co-workers4 that elastic modulus and hardness are lower in the tail part of each enamel rod than in the head area. This is caused by the changes in crystal orientation in the tail area. Xu et al.7 observed that, during indentation, cracks in the axial section of the tooth were longer in the direction approximately parallel to the enamel rods. Also, the cracks, initiated in the direction perpendicular to the rods, are bent to align along the axes of the rods. Xu and co-workers7 also measured a fracture toughness value of 1.30╯±â•¯0.18â•›MPaâ•›m½, associated with the cracks approximately perpendicular to the enamel rods. A lower fracture toughness value of 0.52╯±â•¯0.06â•›MPaâ•›m½ has been measured for cracks nearly parallel to the enamel rods. The argument presented to explain the hardness variation in enamel is also applicable to the observed anisotropy in the hardness of dentin. Similar to the present case, a mineral concentration gradient has been observed1 inside the dentin, with mineral content increasing from the area near the pulp to the outside surface, which is closely related to the variation in hardness and E-modulus. Also, the anisotropy in the mechanical properties of dentin may be the result of the distribution of microstructural features. Mannocci et al.6 observed that high values of tensile strength of dentin are associated with low number density of dentinal tubules. The apical areas of root dentin are found to be more resistant to tension than the coronal ones.6 Thus, in the work presented here, hardness variation, and compositional gradient separately establish the anisotropic nature of human tooth material.
16.7 TRIBOLOGICAL PROPERTIES OF HUMAN TOOTH MATERIAL 16.7.1â•…Friction Behavior Figure 16.4 plots the experimentally measured COF against the number of fretting cycles, during rubbing of tooth balls against sintered alumina. It is observed that,
â•… 261
16.7 Tribological Properties of Human Tooth Material
TEETH BALL VS. ALUMINA BASE PLATE; 100 g LOAD; 8 Hz FREQUENCY; 30 MICRON STROKE; 2000 CYCLE 10,000 CYCLES 4000 CYCLES 6000 CYCLES 8000 CYCLES
0.6 0.5
COF
0.4 0.3 0.2 0.1 0.0 0
2000
4000 6000 8000 10,000 NO. OF CYCLES
Figure 16.4â•… Variation of COF with number of cycles, when human tooth balls are fretted against sintered alumina plate at 1-N load, 8-Hz frequency, under unlubricated conditions.15
for all the different experiments, the COF initially increased to around 1000 cycles and then remained almost constant with certain fluctuations of the steady-state COF value. It may be noted here that a range of COF values, with a maximum measured steady-state COF of 0.55 and the minimum measured at 0.12, have been recorded. For certain cases, such as the tooth sample tested for 8000 cycles, larger variation in the otherwise steady-state frictional stage does appear. The general observation of a sharp rise in COF within the initial fretting cycles is possibly due to the asperities getting knocked off by abrasion during the running-in period. It may be mentioned here that Zheng et al.13 studied the wear behavior of human tooth against titanium on a conventional pin-on-disk machine under both dry and AS conditions and found that the COF is slightly lower (0.2) in AS compared with that under the dry condition (COFâ•›≈â•›0.3). It was also seen in this work that the dental tissue is less burnt and carbonized under the AS condition. This is because AS plays the role of a lubricant as well as a coolant. An important observation from Figure 16.4 is that there is no definite relation between the COF and the number of testing cycles. The initial increase in the COF may be due to the high hardness of the enamel that is being worn with time and adhered to the alumina base plate. That means that, after the initial 1000 cycles, the tribocontact is no longer between the tooth and the alumina rather it is between the enamel and the enamel layer on the alumina. This is possibly why there is a slight decrease and subsequent attainment of the steady-state COF value. The experimental measurements in Figure 16.4 can be further comprehended when we consider the fact that the mechanical properties of teeth are strongly affected by individual differences (e.g., age, sex, physical condition) and also on the location of the contact zone inside a single tooth.15 Zheng et al.13 also observed that tribological behavior of teeth is strongly sensitive to the microstructural orientations and it varies depending on location for the same tooth. Therefore, a difference in hardness and tribological behavior in the present work under discussion is expected in the tested teeth,
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which have been taken from six different persons. Furthermore, in spite of our great effort to produce the tooth balls with identical geometry and location of the tooth sections, it is quite possible that point contacts between the balls and the alumina base plate vary in different testing conditions. All these facts may explain the discrepancy in the relation between the measured COF and the number of cycles during the fretting test.
16.7.2â•… Wear Mechanisms To determine the dominant wear mechanisms of the tooth/alumina fretting couple, SEM-EDS analysis of the worn surface of both tooth ball and alumina base plate was carried out. To show the damage caused at lower numbers of fretting cycles, the evidence of cracked tribolayer formation has been presented in Figure 16.5. From Figure 16.5a, it is clear that the tribolayer formation took place during early stages of fretting wear. Also, the tribolayer appears to be rather thin. The overview of the damage experienced by the human tooth material is presented in the SEM images in Figure 16.6a. A large circular damage zone with diameter of around 15–20â•›mm can be clearly observed. This indicates that the human tooth experiences extensive damage induced by fretting. Such observations need to be correlated with the much lower hardness of the human tooth compared with the Al2O3 counterbody. Some important details of the worn surface are shown in Figure 16.6b. The presence of many long cracks and the dense tribolayer are the common features of the worn surface. EDS analysis shows much stronger P and O peaks from the tribolayer, indicating the dominant formation of Ca–P compounds. Based on the analysis of the surface topography of worn surfaces, it can be summarized that during the initial stages of fretting cycles the tooth surface, being the softer of the two counterbodies, is worn away faster involving oxidation of the tooth surface. Also, a comparison among Figures 16.5 and 16.6 indicates much thicker tribolayer formation on the tooth surface. From the preceding observations, it can be inferred that fretting fatigue appears to be a dominant wear mechanism, involving cracking and material transfer to a softer counterbody. Once the transferred material fragments adhere to the opposing mating solid, tooth–enamel contact occurs. Such a phenomenon is expected to influence the frictional behavior in the case under discussion.
16.8 WEAR PROPERTIES OF GLASS-CERAMICS In a 2009 paper,34 the results of tribological experiments on K2O–B2O3–Al2O3–SiO2– MgO–F glass-ceramics containing about 70% crystals in AS environment are reported. A summary of the important results is provided in this section. The friction and wear properties of the investigated GC against steel were evaluated both in air (dry) and in AS medium for different test durations. Continuous measurement of the COF was done during the entire test period. All the tests were repeated twice or thrice and the reproducibility of the frictional behavior is fairly confirmed. In Figure 16.7, the average value of experimentally measured COF is plotted against the test
16.8 Wear Properties of Glass-Ceramics
â•… 263
Ca
4K Cycles
O
Al Al
Au
5 µm 1.00
2.00
3.00
4.00
3.00
4.00
(a) Ca
6K Cycles
O
Al Al
Au
10 µm 1.00
2.00
(b)
Figure 16.5â•… SEM images revealing the details of the tribolayer as well as EDS compositional analysis of the tribolayer on the worn alumina, after it was fretted against human tooth ball at 1-N load for various time durations: (a) 4000 cycles and (b) 6000 cycles. The cracking can be seen with more clarity in the inset of (a).15
duration for GC/steel tribocouples in air as well as in AS. In general, COF sharply increases within the first few thousand cycles and subsequently reaches the steadystate condition. A higher COF value of 0.88 was recorded for these couples when tests were conducted in dry conditions, while a COF of 0.67 was recorded in the case of AS-medium tests. Also, noticeable variations still appear throughout the steady-state region for both cases. The high COF value obtained under dry conditions can be attributed either to the interaction of the asperities of the two mating surfaces or to the interaction of GC asperities with hard debris particles. As is mentioned later, the wear involving a steel counterbody always produces iron oxide debris particles, which are harder
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Circular damage zone in tooth ball
rom k f hery c a Cr perip the 10 mm (a) Ca P
O
Ca O
P
200 µm (b)
Figure 16.6â•… (a) SEM image (BSE mode) illustrating the overall damage experienced by human tooth after initial 2000 fretting cycles. A crack can be seen to propagate from the periphery of the damaged zone. (b) Higher magnification SEM image (BSE mode), revealing the formation of tribolayer along with the EDS analysis obtained from different parts of the worn surface.15
16.8 Wear Properties of Glass-Ceramics
â•… 265
1.2 1.0
COF
0.8 0.6 0.4 0.2
Dry AS
0.0 0
20,000
40,000
60,000
80,000
100,000
Number of cycles
Figure 16.7â•… Evolution of COF during fretting of investigated glass-ceramics (GC) in dry and AS environment. Counterbody, 10-mm-diameter steel ball; load, 1â•›N; stroke length, 100â•›µm; oscillation frequency, 8â•›Hz.34
than the steel mating body. Therefore, the abrasion caused by hard oxide debris in the case of dry conditions will definitely lead to increased COF. The particular effect of AS medium on the friction and wear behavior will be further discussed in a later section. The wear rate data for the investigated GC materials is given in Figure 16.8. Irrespective of the environment, the wear rate systematically decreases with the fretting test duration; the wear rate for the investigated GCs in the selected tribological environment appears to vary in the range of 10−4–10−5â•›mm3/Nâ•›m. Considering the application of extremely low load (1â•›N) and also low sliding velocity (much less than 0.1â•›m/s), the wear rate of the investigated GC against steel is higher by one order of magnitude than that of various structural ceramic materials (typical wear rate ∼ 10−6â•›mm3/Nâ•›m) investigated earlier by our group.35 Such a difference was obviously to be expected because of the much lower hardness of the GC materials being investigated. However, compared with other GC materials, the wear rate of the investigated GCs is at least one order of magnitude lower.15 Further, it is evident from Figure 16.8 that the wear rate is less in AS medium than in dry conditions for any given test duration. The wear rate decreases systematically from 3.5╯×╯10−4â•›mm3/Nâ•›m to 1.2╯×╯10−5â•›mm3/Nâ•›m during fretting of GC in dry or AS environment. After fretting for 100,000 cycles, a maximum wear rate of 9╯×╯10−5â•›mm3/Nâ•›m was measured for GC under dry conditions, while a minimum wear rate of 1╯×╯10−5â•›mm3/Nâ•›m was measured for GC in AS medium. From this observation, longer durability or better wear resistance during long-term application of these investigated GCs could be expected.
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40.0
Dry AS
Wear rate (×10−5 mm3/Nm)
35.0 30.0 25.0 20.0 15.0 10.0 5.0 0.0 5K
10K 50K Number of cycles
100K
Figure 16.8â•… Wear rate of glass-ceramics after fretting under dry and AS medium conditions for different time durations. Around 10–15% deviation around the reported data were measured in our experimental results.34
16.9 DISCUSSION OF WEAR MECHANISMS OF GLASS-CERAMICS To understand the nature of dominant wear mechanisms, detailed SEM-EDS analysis of worn surfaces of investigated GC samples was conducted after fretting against steel in different environments. Some representative SEM images of the worn surfaces for different durations of fretting (cycles) are presented in Figures 16.9–16.12. During the initial stages, the overall fretting damage of the GC under dry conditions reveals dispersion of small-size debris particles (Fig. 16.9a,b). EDS analysis (inset Fig. 16.9b) reveals the presence of iron along with the glass elements (aluminum, magnesium, silicon, fluorine, potassium, oxygen, etc.), indicative of material transfer from steel ball. However, the tribolayer formation is significantly absent under dry conditions. On the other hand, the worn surface is chiefly characterized by the formation of a tribolayer after fretting in AS medium for 5000 cycles (Fig. 16.9c,d). Though the tribolayer does not cover entire area of the wear scar, deformation or smearing of the smooth layer is certainly evident. Figure 16.10 represents the worn surfaces of GC after 10,000 cycles. The striking difference with the earlier stage is the occurrence of severe wear in terms of considerable pullouts and increased amount of debris when fretted under dry conditions (Fig. 16.10b). However, the absence of the tribolayer indicates that its formation is not feasible after 10,000 cycles. In the case of the AS medium, a thick tribolayer appears to form, and its removal is evident through peel-off or delamination. Further, it may be noted
16.9 Discussion of Wear Mechanisms of Glass-Ceramics
(a)
(b)
(c)
(d)
â•… 267
Figure 16.9â•… SEM image illustrating the overall fretting damage experienced by glass-ceramic plate, after testing against steel ball at 1-N load for 5000 cycles, 8-Hz frequency, 100-µm stroke length: (a,b) under dry conditions and (c,d) in artificial saliva medium. The double-pointed arrows indicate fretting directions.34
that the wear scar is elliptical, when compared with the circular wear scar under dry conditions (Fig. 16.10c). The roughness and pullouts as well as the number of agglomerated brighter debris particles increased when the number of cycles increased to 50,000 under dry conditions. When the EDS spectrum of the unworn GC surface (inset in Fig. 16.11a) is compared with that of the debris (inset in Fig. 16.11b), it can be observed that fretting results in the pullout and oxidation of the constituents of the GC and the counterbody steel material. Further, the iron found in the debris on the worn surface of the GC is indicative of material transfer from the steel ball. On the other hand, very long duration of fretting in AS medium resulted in the small and elliptical wear scar that is completely covered with the dense layer (Fig. 16.11c). The dispersion of debris around the periphery of the wear scar is an interesting feature, which is otherwise absent in case of the dry conditions. Such dispersion of debris might be attributed to the continuous flushing in and out of the AS caused by the reciprocating action of fretting. High-magnification SEM image of the GC worn surface after
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(a)
(b)
(c)
(d)
Figure 16.10â•… SEM image illustrating the fretting damage experienced by glass-ceramic plate, after testing against steel ball under dry conditions (a,b) and in AS (c,d) at 1-N load for 10,000 cycles, 8-Hz frequency, 100-µm stroke length. The double-pointed arrows indicate fretting directions.34
50,000 fretting cycles (Fig. 16.11d) reveals the presence of several cracks, indicating the characteristic brittle fracture and spalling of the tribolayer. Similar features were explained as a result of contact fracture and spalling, when Dicor GC was rubbed against alumina in a distilled-water environment.28 The opening-up of cracks is due to the hygroscopic nature of the layers formed during fretting in AS medium. The severity of the pullouts is further increased with the increase in the number of cycles (100,000) under dry conditions (Fig. 16.12). Further, Figure 16.12b reveals the formation of thin-layer fragments and their smearing on the worn surface after 100,000 cycles. The characteristic features of the worn surface after 100,000 cycles in AS medium are, in general, similar to those after 50,000 cycles, namely, the elliptical and smaller wear scar, formation of a potential tribolayer protecting the base surface, the dispersion of fine debris around the periphery of the wear scar, and the brittle nature of the tribolayer with many cracks. In addition, the tribochemical layer has fine grooves in the fretting direction, indicating mild abrasion due to the sliding of hard debris particles (see Fig. 16.12d).
16.9 Discussion of Wear Mechanisms of Glass-Ceramics Si
Si
O F
Al Mg
â•… 269
Al Au O Mg
Au K
B F
Fe
K
20 µm
500 µm (a)
(b) Au Si
Si
Al Mg O K F
Au Al Mg F K O
Fe
Fe
10 µm
500 µm (c)
(d)
Figure 16.11â•… SEM image illustrating the overall fretting damage experienced by glass-ceramic plate, after testing against steel ball at 1-N load for 50,000 cycles, 8-Hz frequency, 100-µm stroke length: (a,b) under dry condtions and (c,d) in AS medium. The double-pointed arrows indicate fretting directions.34
Based on the aforementioned observations of friction, wear, and topographical features of the worn surfaces, the effect of the AS medium on the fretting wear behavior of the investigated GC during fretting against steel can be discussed. As the hardness difference between the mating materials is much less, the local welding and detachment of asperity junctions result in significant adhesion during the initial stages of fretting under dry conditions. The steep rise of the COF plot during the running-in period could be attributed to the previously stated phenomenon. As fretting progresses, asperities are subjected to deformation or fracture, resulting in the formation of debris. The interaction of either surface with continuously formed debris provides the steady state in the frictional behavior. The debris are transferred to the countersurface and/or subjected to oxidation in ambient testing conditions. Further, the hard debris particles are responsible for causing abrasion on the countersurface. Thus, the debris consists of oxides of both the mating surfaces (see EDS analysis). The hard oxide debris formed during continuous fretting leads to three-body wear under dry conditions; the prolonged fretting results in the formation of increased
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(a)
(b)
(c)
(d)
Figure 16.12â•… SEM image illustrating the overall fretting damage experienced by glass-ceramic plate, after testing against steel ball at 1-N load for 100,000 cycles, 8-Hz frequency, 100-µm stroke length (a,b) under dry conditions and (c,d) in AS medium. The double-pointed arrows indicate fretting directions.34
amounts of agglomerated debris. Further, the higher value of average COF (∼0.9) observed throughout the test duration indicates that the wear occurs mainly through the removal of crystalline mica phase. For tests conducted in AS medium, it is highly possible that the pulled-out mica debris particles are subjected to hydroxylation. This could be considered as a significant effect of corrosion occurring at the wet tribocontacts, as also observed in the case of Ti-Ca-P biocomposites.36 Eventually the hydroxides are compacted to form a tribochemical layer during extended fretting such that it protects the base GC surface from further wear. The reduced friction and wear rate against those under dry conditions is in agreement with this fact. The presence of macrocracks on the tribochemical layer indicates drying of the viscous layer, resulting in brittle fracture. To analyze thermomechanical effects at the dry tribocontact, both Hertzian contact stress and the maximum contact temperature rise have been analytically computed. Using the Hertzian contact mechanics approach,37 the initial contact pressure is calculated as 196â•›MPa and the initial contact diameter is approximately
16.10 Comparison with Existing Glass-Ceramic Materials
â•… 271
40â•›µm. Observation of Figures 16.9–16.12 makes it clear that the fretting at the tribocontact considerably expand the scar diameter to 500–600â•›µm, depending on the number of testing cycles. Such a considerable increase in scar diameter decreases the contact pressure to 0.51–0.35â•›MPa. Additionally, in the present case, Archard’s model38 estimated very low rise in contact temperature (17°C) at the dry tribocontact. Such negligible contact temperature rise may be attributed to the combination of low load (1â•›N) and sliding speed (0.0016â•›m/s). From this, it is therefore clear that the extent of damage caused by wear must be related to the severity of contact stress, rather than to the thermal effect when tests were conducted under dry conditions. Further, surfaces worn under dry conditions reveal rough surfaces compared with those in AS medium. The contrast in the back scattered electron (BSE) images suggests that the crystalline glass phase is pulled out during the wear. Considering the large amount (70%) of mica crystalline phase in the unworn GC, its easy removal from the soft glass phase and subsequent fracture is highly possible during the fretting process. On the other hand, the formation of the hydroxide-rich tribolayer is responsible for the reduction in wear rate in AS medium. The failure of the layer and the fracture of crystalline phases resulted in the higher value of COF (∼0.67), which is still less than that observed under dry conditions. Therefore, it is understood from the preceding observation that the dominant wear mechanism changes from stress-induced mechanical wear under dry conditions to the corrosion-induced tribochemical wear at wet contacts, during fretting of GC against steel. The implication of this present investigation is significant due to the influence of saliva on the wear performance of the developed dental restorative GC material. Under dry conditions, the COF and wear rate are high due to the removal of mica crystals and subsequent abrasion. On the other hand, the presence of AS medium helps in reducing the friction and wear by forming a dense viscous tribochemical layer. Based on the experimental results presented here, obtained under the selected fretting conditions, and also against the backdrop of the good combination with their in vitro reactivity,31 the investigated mica-based GC materials appear to be a good choice for dental applications.
16.10 COMPARISON WITH EXISTING GLASSCERAMIC MATERIALS It is necessary to compare the tribological properties of the investigated GC material discussed here with those of previously developed materials. For this purpose, a summary of the friction and wear rate data along with the operating parameters for some of the commercial GCs (such as Dicor) or GCs developed in various laboratories is provided in Table 16.3. Additionally, some results of the test with human teeth are also mentioned in Table 16.3. It is known that the Dicor material, developed commercially for dental restorative applications, shows a modest combination of hardness (350â•›MPa), E-modulus (66.2–91.1â•›GPa), and strength (127â•›MPa). A comparison with Dicor in terms of tribological properties demonÂ� strates a better wear-resistance property and comparable frictional property of
272â•…
CHAPTER 16â•… Case Study: Natural Tooth and Dental Restorative Materials
TABLE 16.3â•… Summary of the Tribology Test Results Obtained with Some of Previously Developed Glass-Ceramics as well as Human Teeth and Comparison with the Presently Investigated GC34
Tribocouple Human teeth vs. steel Human teeth vs. Al2O3 Dicor vs. Al2O3 Dicor vs. Al2O3 CaO–MgO– Al2O3–SiO2 (self-mated) MgO–CaO–SiO2 P2O5–F vs. ZrO2 K2O–B2O3– Al2O3–SiO2– MgO–F vs. steel
Operating conditions 20â•›N, dry/AS 0.002â•›m/s 1â•›N, AS; 0.0005â•›m/s, 8000 cycles 4.9â•›N, 0.0014â•›m/s, dry 1â•›N, 0.0025â•›m/s, distilled water 0.01–0.5â•›m/s; dry; contact pressure- 0.1– 1.4â•›MPa 10â•›N, 0.025â•›m/s, dry 1â•›N, 0.0016â•›m/s, dry/AS; 100,000 cycles
COF
Wear rate (mm3/N/m)
Wear mechanisms
0.8–1.2 (dry) 1.0 (AS) 0.12–0.55
—
Oxidative wear and microfracture
—
Fretting fatigue; adhesive wear
0.7–0.077
2.6╯×╯10−3
Microfracture
0.4–0.6
10−3 to 10−4
0.05–0.65
10−3 to 10−4
Localized fracture Microcracking, abrasion
0.75
0.7╯×╯10−4
Abrasive and adhesive wear
0.88 (dry); 0.67(AS)
12╯×╯10−5 (dry); 2╯×╯10−5 (AS)
Tribomechanical wear (dry); tribochemical wear (AS)
The variation in friction/wear rate depends on the variation in operating conditions.
(K2O3-B2O3-Al2O3-SiO2-MgO-F GC) in AS under similar operating conditions (see Table 16.3). It is therefore expected that the newer GC material will have better durability than Dicor. As far as the frictional properties of human teeth are concerned, the investigated GC-material/steel couple discussed here has a lower value of COF than the human-teeth/steel mating couple. Also, the material exhibits tribological properties comparable to those of the MgO–CaO–SiO2–P2O5–F/ZrO2 tribocouple. Consideration of the mechanisms of material removal, Table 16.3 indicates that the microfracture–tribomechanical wear is the major wear mechanisms for GCs. While this is true under dry conditions for the case under discussion, this chapter also indicates that the tribochemical layer formation governs wear in AS medium. In closing, a comprehensive comparison, in terms of mechanical and wear properties, of the investigated GC discussed here with previously developed GCs indicate the former is a good potential dental restorative material (also see Ref. 23). However, further study is needed to assess the optical translucence, machinability, and long-term durability of the newer material. Recent work has revealed the good biological compatibility, in terms of good cell viability (MTT tests) of human osteoblast cells as well as better antimicrobial properties (Escherichia coli and
16.11 Concluding Remarks
â•… 273
Staphylococcus epidermidis cell lines) of the GC system under discussion here.39 However, in vivo osseointegration (clinical trials) needs to be carried out before applications can be realized.
16.11 CONCLUDING REMARKS a) The mechanical property measurements on human teeth reveal the variation in hardness, with enamel being the hardest (around 3.5â•›GPa) and dentin being softer than enamel (less than 1â•›GPa). Hardness value decreases with progressively increasing depth from the surface. Such a variation partially depends upon the variation in mineral concentration in enamel and dentin, with a possible dependence of local microstructural features, such as enamel rod orientation and dentinal tubule density. b) While alumina does not undergo any measurable wear loss, the human tooth material exhibits severe fretting-induced damage. The transfer layer (from the tooth ball) on alumina is found to be (Ca,P,O)-rich and the layer is found to contain numerous microcracks. No indication of abrasion is observed on any of the mating counterbodies. c) Fretting fatigue and adhesive wear seem to be the dominating wear mechanisms. A thick tribolayer rich in Ca, P, and O is observed to adhere strongly to the worn tooth surface. During the initial fretting cycles, the COF decreases as the tribological contact is established between the tooth enamel and the adhered layer of enamel on the alumina surface. To observe the wear behavior of dental restorative materials, a GC with 70% mica crystals in a K2O–B2O3–Al2O3–SiO2–MgO–F system was subjected to fretting wear against steel in dry and AS environments. The initial Hertzian contact stress was found to be 196â•›MPa for the selected operating parameters. The influence of the environment and duration of fretting on the friction and wear behavior of GC material can be summarized as follows: a) The COF increased during initial running-in period and subsequently attained a steady-state value, irrespective of the fretting environment. However, a higher value of steady-state COF (∼0.88) was measured in dry contact, while in AS medium a much lower COF (0.67) was recorded. A higher COF under ambient conditions correlates well with the severity in abrasion. b) Wear rate varied on the order of 10−4–10−5â•›mm3/Nâ•›m and a systematic decrease in the wear rate with test duration was recorded in both testing media. A comparison of the newly developed GC discussed here with the previously developed GCs as well as commercial Dicor material shows a better combination of friction and wear-resistance properties for the newer GC. c) The topographical observations of the worn surfaces using SEM-EDS analysis indicate that the material was removed mainly by tribomechanical wear assisted by the removal of crystals under dry conditions, whereas the formation and brittle fracture of the tribochemical layer was dominant in AS medium.
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CHAPTER 16â•… Case Study: Natural Tooth and Dental Restorative Materials
Also, material transfer from the counterbody was also observed in the ambient environment as well as in the AS medium. d) Irrespective of the test environment, a large amount of debris particles were formed and found to be dispersed mainly around wear scars. In the AS medium, the stability of the thick tribochemical layer seems to control material removal. The occasional spalling of the tribolayer results in wear, but to a lesser extent than that under dry and ambient conditions.
REFERENCES ╇ 1â•… A. Kishen, U. Ramamurty, and A. Asundi. Experimental studies on the nature of property gradients in the human dentine. J. Biomed. Mater. Res. 51 (2000), 650–659. ╇ 2â•… M. Giannini, C. J. Soares, and R. M. D. Carvalho. Ultimate tensile strength of tooth structures. Dent. Mater. 20 (2004), 322–332. ╇ 3â•… J. L. Cuy, A. B. Manna, K. J. Livi, M. F. Teaford, and T. P. Weihs. Nanoindentation mapping of the mechanical properties of human molar tooth enamel. Arch. Oral Biol. 47 (2002), 281–291. ╇ 4â•… S. Habelitz, S. J. Marshall, G. W. Jr. Marshall, and M. Balooch. Mechanical properties of human dental enamel on the nanometre scale. Arch. Oral Biol. 46 (2001), 173–183. ╇ 5â•… R. Z. Wang and S. Weiner. Strain structure relations in human teeth using Moire fringes. J. Biomech. 31 (1998), 135–141. ╇ 6â•… F. Mannocci, P. Pilecki, E. Bertelli, and T. F. Watson. Density of dentinal tubules affects the tensile strength of root dentin. Dent. Mater. 20 (2004), 293–296. ╇ 7â•… H H K Xu, D. T. Smith, S. Jahanmir, E. Romberg, J. R. Kelly, V. P. Thompson, and E. D. Rekow. Indentation damage and mechanical properties of human enamel and dentin. J. Dent. Res. 77 (1998), 472–480. ╇ 8â•… N. Burak, J. A. Kaidonis, L. C. Richards, and G. C. Townsend. Experimental studies of human dentine wear. Arch. Oral Biol. 44 (1999), 885–887. ╇ 9â•… A. S. Al-Hiyasat, W. P. Saundersb, S. W. Sharkey, R. Smith G Mc, and W. H. Gilmour. Investigation of human enamel wear against four dental ceramics and gold. J. Dent. 26 (1998), 487–495. 10â•… R. R. Seghi, S. F. Rosenstiel, and P. Bauer. Abrasion of human enamel by different dental ceramics in vitro. J. Dent. Res. 70 (1991), 221–225. 11â•… M. Z. A. M. Sulong and R. A. Aziz. Wear of materials used in dentistry: A review of the literature. J. Prosthet. Dent. 63 (1990), 342–349. 12â•… N. Meredith, M. Sherriff, D. J. Setchell, and S. A. V. Swanson. Measurement of the microhardness and Young’s modulus of human enamel and dentine using an indentation technique. Archives of Oral Biology 41(6) (1996), 539–545. 13â•… J. Zheng, Z. R. Zhou, J. Zhang, H. Li, and H. Y. Yu. On the friction and wear behaviour of human tooth enamel and dentin. Wear 255(7–12) (2003), 967–974. 14â•… F. S. L. Wong, P. Anderson, H. Fan, and G. R. Davis. X-ray microtomographic study of mineral concentration distribution in deciduous enamel. Archives of Oral Biology 49(11) (2004), 937–944. 15â•… S. Roy and B. Basu. Mechanical and tribological characterization of human tooth. Mater. Character. 59 (2008), 747–756. 16â•… A. Gebhardt, T. Höche, G. Carl, and I. I. Khodos. TEM study on the origin of cabbage-shaped mica crystal aggregates in machinable glass-ceramics. Acta Mater. 47 (1999), 4427–4434. 17â•… T. Höche, S. Habelitz, and I. I. Khodos. Origin of unusual fluorophlogopite morphology in mica glass-ceramics of the system SiO2–Al2O3– MgO–K2O–Na2O–F2. J. Cryst. Growth 192 (1998), 185–195. 18â•… Lj. Radonjic´ and Lj. Nikolic´ . The effect of fluorine source and concentration on the crystallization of machinable glass-ceramics. J. Eur. Ceram. Soc. 7 (1991), 11–16. 19â•… K. Cheng, J. Wan, and K. Liang. Crystallization of R2O–MgO–Al2O3–B2O3–SiO2–F (R=K+, Na+) glasses with different fluorine source. Mater. Lett. 47 (2001), 1–6.
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20â•… S. C. V. Clausbruch, M. Schweiger, W. Höland, and V. Rheinberger. The effect of P2O5 on the crystallization and microstructure of glass ceramics in the SiO2–Li2O–K2O–ZnO–P2O5 system. J. NonCrystalline Solids 263 & 264 (2000), 388–394. 21â•… S. Habelitz, T. Höche, R. Hergt, G. Carl, and C. Rüssel. Microstructural design through epitaxial growth in extruded mica glass-ceramics. Acta Mater. 47 (1999), 2831–2840. 22â•… T. Höche, S. Habelitz, and I. Avramov. Crystal morphology engineering in SiO2–Al2O3–MgO–K2O– Na2O–F−mica glass-ceramics. Acta Mater. 47 (1999), 735–744. 23â•… J. Park and A. Ozturk. Tribological properties of MgO–CaO–SiO2–P2O5–F− based glass-ceramic for dental applications. Mater. Lett. 61 (2007), 1916–1921. 24â•… H. Xiao, Y. Cheng, Q. Yang, and T. Senda. Mechanical and tribological properties of calcia– magnesia–alumina–silica-based glass–ceramics prepared by in situ crystallization. Mater. Sci. Eng. A 423 (2006), 170–174. 25â•… S. Jahanmir and X. Dong. Wear mechanism of a dental glass-ceramic. Wear 181–183 (1995), 821–825. 26â•… H. Y. Yu, Z. B. Cai, P. D. Ren, M. H. Zhu, and Z. R. Zhou. Friction and wear behavior of dental feldspathic porcelain. Wear 261 (2006), 611–621. 27â•… G. K. H. Zhum and P. Neumann. Oscillating sliding wear of cordierite glass and ceramic in liquid media. Wear 203–204 (1997), 107–118. 28â•… V. S. Nagarajan and S. Jahanmir. The relationship between microstructure and wear of micacontaining glass-ceramics. Wear 200 (1996), 176–185. 29â•… V. S. Nagarajan, S. Jahanmir, and V. P. Thompson. In vitro contact wear of dental composites. Dent. Mater. 20 (2004), 63–71. 30â•… S. Roy and B. Basu. On the development of two characteristically different crystal morphology in SiO2-MgO-Al2O3-K2O-B2O3-F glass-ceramic system. J. Mater. Sci. Mater. Med. 20(1) (2008), 51–66. 31â•… A. R. Molla and B. Basu. Microstructure, mechanical, and in vitro properties of mica glass-ceramics with varying fluorine content. J. Mater. Sci. Mater. Med. 20(4) (2009), 869–882. 32â•… S. Roy and B. Basu. In vitro dissolution behaviour of SiO2-MgO-Al2O3-K2O-B2O3-F glass-ceramic system. J. Mater. Sci. Mater. Med. 19 (2008), 3123–3133. 33â•… H. Li and Z. R. Zhou. Wear behaviour of human teeth in dry and artificial saliva conditions. Wear 249 (2002), 980–984. 34â•… A. R. Molla, B. V. Manoj Kumar, and B. Basu. Friction and wear mechanisms of K2O-B2O3-Al2O3SiO2-MgO-F glass-ceramics. J. Eur. Ceram. Soc. 29 (2009), 2481–2489. 35â•… D. Sarkar, B. Basu, M. C. Chu, and S. J. Cho. Is glass infiltration beneficial to improve fretting wear properties for alumina? J. Am. Ceram. Soc. 90 (2007), 523–532. 36â•… M. Karanjai, B. V. M. Kumar, and B. Basu. Fretting wear study on Ti-Ca-P biocomposites in dry and simulated body fluid. Mater. Sci. Eng. A 475 (2008), 299–307. 37â•… K. L. Johnson. Contact Mechanics. Cambridge University Press, UK, 2001, 93. 38â•… J. F. Archard. Contact and rubbing of flat surfaces. J. Appl. Phys. 24 (1953), 981–988. 39â•… S. Kalmodia, A. R. Molla, and B. Basu. In vitro cellular adhesion and antimicrobial property of SiO2-MgO-Al2O3-K2O-B2O3-F glass ceramic. Journal of Materials Science: Materials in Medicine 21 (2010), 1297–1309.
CH A P T E R
17
CASE STUDY: GLASSINFILTRATED ALUMINA In this chapter, the tribological properties of glass-infiltrated alumina are discussed with a major aim of scientifically understanding the influence of glass infiltration on the underlying mechanisms of material removal (when fretted against steel). The formation of tribochemical reaction products is explained using the thermodynamically feasible reactions. The major wear mechanisms are identified as fatigue wear, three-body abrasion, and formation of a tribochemical layer.
17.1 INTRODUCTION Among the structural ceramics, sintered alumina has been widely researched for various tribological applications. However, because of low fracture toughness and mechanical strength, Al2O3 components are prone to catastrophic failures with major damage to the workpiece.1 Hence, alumina ceramics fail to reach their full potential, because of strength-limiting surface flaws.2 The machining-induced damage can penetrate approximately one grain diameter from the surface.3 It is therefore expected that an appropriate surface coating or strengthening technique of sintered Al2O3 surface can have a wide range of advantages such as high strength and wear resistance.4 Recent research results on the possibility of improving of surface properties of wear-resistant materials suggest that deposition of coatings by several processes could be a potential solution.5–8 A well-recognized method of increasing the strength of brittle materials is the introduction of surface compressive stress by depositing a coating material of lower coefficient of thermal expansion (CTE).9 An important factor is that high residual stresses are introduced during postdeposition cooling, which in turn affect the behavior of the coating. The residual stresses are potentially beneficial in improving the strength and tribological behavior of the component. For example, compressive stresses develop in physical vapor deposited (PVD) coating, which increase the fracture strength.10 However, large compressive stresses can cause failure of the coating by debonding, especially at coating edges.11
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
276
17.2 Materials and Experiments
â•… 277
Among several alternative approaches, a common glass-infiltration technique can increase the strength of the alumina surface through generation of compressive residual stress and blunting of the fine surface cracks generated during machining.12 For example, the use of a relatively thick mullite layer (>300â•›µm) on alumina substrate has been reported.13 Interestingly, Kim et al.14 reported at least 60% increase in flexural strength of Al2O3 through the “SiO2-smoke” deposition process, where the deposited coating thickness was only 1â•›µm. Although such coatings show promise in strengthening ceramics, their potential to improve tribological properties has not been investigated. In one among the few such studies, Kalin et al.15 studied the tribological behavior of glass-infiltrated alumina (AG) against alumina balls. Their experiments revealed that the wear rate of both the alumina balls and the glassinfiltrated alumina disks increased with disk roughness. In a subsequent study,16 the same group of researchers reported that the wear of AG against Al2O3 increases linearly with load and sliding distance in water. Their research results also suggested the possibility of the application of AG in a wider tribological domain, besides the well-established use in dental restoration. Against this backdrop, the experimental results summarized elsewhere17 are presented in this chapter to describe the wear characteristics of glass-infiltrated alumina.
17.2 MATERIALS AND EXPERIMENTS Dense, 99.5% pure Al2O3 ceramic plates with an average grain size of 5â•›µm were machined into flexural specimens of 3â•›mm╯×╯4â•›mm╯×╯40â•›mm. The chemical and physical properties of both the substrate and coating substance are provided in Table 17.1. Small pieces (1–2â•›mm) of the glass cullet were placed on a 4-mm╯×╯40-mm surface, and the specimens were heat-treated at 1500°C for 5 minutes in air. Flexural strength was measured on a four-point fixture configuration, having inner and outer spans of 10 and 30â•›mm, respectively. The measured hardness, indentation toughness, and strength values of as-received alumina (AS) and glass-infiltrated alumina (AG) are summarized in Table 17.2. As can be seen, the as-sintered alumina has a lower strength than the glass-infiltrated alumina. The average microhardness of as-sintered Al2O3 is 17.6â•›GPa, while that of surface-treated Al2O3 is 19.2â•›GPa. A noticeable increase in indented toughness from 3.9 to 4.6â•›MPaâ•›m1/2 with glass infiltration was also recorded. As can be noticed from Table 17.2, the as-sintered alumina has a lower TABLE 17.1â•… Chemical and Physical Properties of Supplied Substrate and Coating Substances17
Material Alumina (substrate) Pyrex glass (coating substance)
Composition
CTE (×10−6)/°C
Density (gm/cc)
Elastic modulus (GPa)
Al2O3╯=╯99.5% and rest oxide impurities SiO2╯=╯81% and rest oxide impurities
9.1
3.89
378
3.2
2.23
62
278â•…
CHAPTER 17â•… Case Study: Glass-Infiltrated Alumina
TABLE 17.2â•… Mechanical Properties of As-Received Alumina (AS) and Glass-Infiltrated Alumina (AG)17
Material
Sintered alumina Glass-infiltrated alumina
Ceramic designation
Flexural strength (MPa)
Vickers hardness, HV5 (GPa)
Indentation fracture toughness, KC (MPaâ•›m1/2)
AS AG
330╯±â•¯50 530╯±â•¯35
17.6╯±â•¯1.6 19.2╯±â•¯1.2
3.9╯±â•¯0.6 4.6╯±â•¯0.8
strength than the glass-infiltrated alumina. Also, the increase in hardness and toughness of AG (compared with AS) is attributed to the residual compressive stresses. The fretting wear experiments on AS and AG surfaces against bearing steel can predict the tribological behavior for hybrid (ceramic-metal) bearings and, accordingly, the results summarized in this chapter were obtained after experiments with varying loads (2, 5, and 10â•›N) and varying cycles (10,000, 50,000, and 100,000), at constant frequency (8â•›Hz) and constant displacement stroke (100â•›µm) in air at room temperature (28╯±â•¯2°C) with relative humidity (RH) 50╯±â•¯5%.
17.3 FRICTIONAL PROPERTIES It can be noted from Figure 17.1 that the evolution of frictional behavior is strongly dependent on normal load as well as fretting cycles. A distinct transition in the frictional behavior is recorded at 5â•›N for the AG/steel tribocouple. As shown in Figure 17.1b, the initial friction coefficient of AG/steel rose suddenly to 0.7 at intermediate load (5â•›N), then dropped to 0.54 at 40,000 cycles due to formation of a tribolayer and subsequently maintained a steady-state behavior. However, a lower steady-state COF of less than 0.5 was measured at 10-N load.
17.4 WEAR RESISTANCE AND WEAR MECHANISMS The wear volume, measured using a laser surface profilometer, was normalized with respect to normal load and total sliding distance in order to obtain the specific wear rate for both specimens. The measured wear rate data are mentioned in Table 17.3, whereas Figure 17.2 illustrates the variation of wear rate (after 100,000 cycles) with respect to load. A careful observation of data in Table 17.3 reveals that the engineered surface exhibits better wear resistance at lower load, whereas at higher load the wear rate of both materials are almost identical, on the order of 10−6â•›mm3/N·m. The wear rate of hard ceramics such as Al2O3, SiC, and Si3N4 were measured to be on the order of 10−6â•›mm3/N·m (lowest up to 10−9â•›mm3/N·m).18,19 It can be noted here that the wear rate for water-lubricated self-mated Al2O3 can vary by one order of magnitude (10−6 to 10−5â•›mm3/N·m), depending on load (10–40â•›N).20 In case of glassinfiltrated alumina, the wear rate of the alumina counterface increases as the initial
17.4 Wear Resistance and Wear Mechanisms
â•… 279
0.9 0.8 0.7
COF
0.6 0.5 0.4 0.3 0.2
2N 5N 10 N
0.1 0.0 0
20,000
40,000 60,000 No. of Cycles
80,000
100,000
(a) 0.9 0.8 0.7
COF
0.6 0.5 0.4 0.3 0.2
2N 5N 10 N
0.1 0.0 0
20,000
40,000
60,000
80,000
100,000
No. of Cycles (b)
Figure 17.1â•… Plot depicting the evolution of coefficient of friction (COF) when (a) as-sintered alumina (AS) and (b) glass-infiltrated alumina (AG) surfaces were fretted against steel at different loads for 100,000 cycles with a frequency of 8â•›Hz and 100-µm displacement.17
roughness of the glass-infiltrated surface is increased, when sliding against highpurity alumina.15 To show the difference in wear mechanisms, some representative SEM images revealing the worn surface topographical features are provided in Figure 17.3. From the features of the worn surface, it is clear that abrasion results in wear of AS ceramic at lower (2â•›N) load (Fig. 17.3a,b), whereas the combination of abrasion, adhesion,
280â•…
CHAPTER 17â•… Case Study: Glass-Infiltrated Alumina
TABLE 17.3â•… Coefficient of Friction and Wear Rate of Investigated Materials, after Being Fretted at Various Load and Cycles17
Sample code
Load
No. of cycles
AG
2
5
10
AS
2
5
10
Wear Rate (mm3/N·m)
1.5 × 10−5
Wear rate
COF
mm3/Nm
(N) 10,000 50,000 100,000 10,000 50,000 100,000 10,000 50,000 100,000 10,000 50,000 100,000 10,000 50,000 100,000 10,000 50,000 100,000
1.5╯×╯10−4 6╯×╯10−5 7.5╯×╯10−6 7╯×╯10−5 2.8╯×╯10−5 3╯×╯10−6 5╯×╯10−6 2╯×╯10−6 1.5╯×╯10−6 2╯×╯10−4 5╯×╯10−5 1.1╯×╯10−5 9╯×╯10−5 3.2╯×╯10−5 5╯×╯10−6 1╯×╯10−5 4╯×╯10−6 1.6╯×╯10−6
0.63 0.59 0.54 0.57 0.54 0.52 0.53 0.49 0.44 0.62 0.58 0.59 0.62 0.56 0.57 0.56 0.52 0.49
Alumina AG
1.0 × 10−5
5.0 × 10−6
1.0 × 10−6
2
4
6 Load (N)
8
10
Figure 17.2â•… Wear rates of alumina and AG surfaces at different loads for 100,000 cycles. Typically, 5% deviation around the average wear data is observed in our experiments.17
17.4 Wear Resistance and Wear Mechanisms
10 µm
2N
(d)
1 µm
1 µm
10 µm
5N
(e)
10 µm (c)
10 µm
5N
(b)
10 N
10 µm
2N
(a)
â•… 281
10 N
10 µm (f )
Figure 17.3â•… Representative SEM images (back-scattered electron [BSE] mode), illustrating the topographical features of the fretted surfaces of as-sintered alumina (a–c) alumina and glass-infiltrated alumina (d–f), after fretting tests for 100,000 cycles at varying loads, as mentioned in individual figure. Double-pointed arrow indicates the fretting direction. The details of the deformed tribolayer and debris particle are also shown in the inset of (b) and (e).17
and tribochemical wear contributes to the material removal at load ≥2â•›N for both the AS and AG ceramics. In particular, at higher loads (≥2â•›N), the tribochemical layer formation as well as spalling of the tribolayer can be noticed (Fig. 17.3c,d). In addition, a different contrast on SEM images in Figure 17.3e,f reveals the presence of a transfer layer and wear debris particles, adhering on the abrasive scratches. The characteristics of the tribochemical layer are discussed in the following section.
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CHAPTER 17â•… Case Study: Glass-Infiltrated Alumina
17.5 WEAR DEBRIS ANALYSIS AND TRIBOCHEMICAL REACTIONS After the wear testing, the debris particles are found to be mostly of submicron size and spherical in nature. An average size (a representative line across the image) of the debris particles is approximately 1â•›µm. Occasionally, agglomerated wear debris of size up to 2–5â•›µm and a smaller fraction of submicron-sized debris particles are observed. Two dominant x-ray peaks, corresponding to the presence of Fe2SiO4 and Fe(OH)3 are recorded from wear debris (see Fig. 17.4). Also, an overall wavy baseline is observed, because of the presence of a large amount of amorphous SiO2 on the engineered surface. The x-ray diffraction (XRD) analysis of wear debris particles, therefore, indicates that a probable chemical reaction occurred between Fe (from the steel ball) and SiO2 (from the engineered alumina surface). From the Ellingham diagram, it can be commented that the formation of Fe3O4 is thermodynamically more favorable at lower temperature and Fe3O4 can be further oxidized to form Fe2O3 at higher temperature, 3Fe + 2O2 = Fe 3 O 4, 4 Fe3 O4 + O2 = 6 Fe 2 O3.
(17.1a) (17.1b)
Once Fe oxides are preferentially formed during early tribological interaction, they can be hydrolyzed favorably in the presence of water molecules to form Fe(OH)3, by the following probable reaction: Fe 2 O3 + 3H 2 O = 2Fe(OH)3.
(17.2)
Fe(OH)3
Intensity (A.U.)
Fe2SiO4
20
30
40
50
2*θ (in degrees)
60
70
Figure 17.4â•… XRD spectra obtained from the debris particle of glass-infiltrated alumina surface, after fretting against bearing steel at 2-N load for 100,000 cycles. The baseline was observed for the presence of amorphous glassy coating.17
17.6 Influence of Glass Infiltration on Wear Properties
â•… 283
Simultaneously, in the case of glass-infiltrated alumina, more possible reactions can take place and those can explain the formation of Fe2SiO4 (crystalline fayalite):
2 Fe + SiO2 + O2 = Fe 2 SiO 4 , 2 Fe3 O4 + 3SiO2 = 3Fe 2 SiO4 + O2.
(17.3) (17.4)
Based on the available thermodynamic data (HSC Chemistry® Version 5.1), the calculation reveals the ΔG for the Reaction 17.2 is negative (−1000 to −600â•›kJ) over the temperature variation of room temperature and 1000°C. Similarly, the ΔG for the Reaction 17.3 varies in the range of −1379â•›kJ to −1025â•›kJ within the temperature range of 25–1127°C.21 For Reaction 17.4, the thermodynamic property calculation reveals that the ΔG is, however, positive and decreases with temperature to reach a value lower than 100â•›kJ at a temperature of 1000°C. The preceding discussion signifies that only Reactions 17.2 and 17.3 are thermodynamically feasible and that such reactions, in combination, can be used to describe the tribochemical reactions at the investigated fretting contacts.
17.6 INFLUENCE OF GLASS INFILTRATION ON WEAR PROPERTIES This section focuses on two major aspects: (1) mechanical property improvement due to the glass-infiltration process and (2) underlying mechanisms to explain observed differences in friction and wear properties of AS and AG specimens. In the work discussed here, the conventional approach of firing and subsequent grinding–polishing technique is adopted to infiltrate the amorphous glass on alumina surface. On cooling from the firing temperature, the Al2O3 substrate tries to shrink more than the glass layer but, due to the displacement compatibility, the coated surface experiences compressive stress. The maximum value of the residual stress, which arises because of a difference in expansion coefficient, is given by:22
σ max = E ∆α T/[(1 − ν)],
(17.5)
where E and ν are Young’s modulus and Poisson’s ratio (0.26) of the surface layer, Δα is the difference in the thermal expansion coefficient of the bulk and the surface, and ΔT is the difference between the final temperature and the stress-free temperature. From Equation 17.5, it is clear that the formation of a glass layer on the surface of Al2O3 results in compressive residual stress (∼715â•›MPa), because of the difference in the CTE of Pyrex glass (3.2╯×╯10−6/°C) and that of alumina (9╯×╯10−6/°C).23 It is reported that compressive residual stresses on the ceramic surface can increase the strength of the material.9,23,24 In the case of the as-sintered alumina surface, major wear mechanisms were recorded as, (a) abrasion and (b) tribochemical layer formation. However, additional wear mechanisms for the engineered alumina surface include (a) static fatigue and (b) adhesive wear. Based on the indentation fracture mechanics concepts, Roberts23 proposed a theoretical model to compute the depth of cracks on brittle materials due to abrasion, and this model enables us to predict the minimum load required for fracture due to abrasion:
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P* =
5447β K IC K IC , πη2 θ 4 H
(17.6)
where P* is the minimum load required to produce fracture from a point contact (in newtons), η is a constant, β is the constant relating hardness to diagonal (2.16 for Vickers indentation), θ is the geometrical constant ( ≈0.2), KIC is the fracture toughness (MPaâ•›m1/2), and H is the hardness (GPa) of the material indented (GPa). Incorporating the material properties (see Table 17.2) into Equation 17.6, it has been found that a minimum of ∼3-N load is necessary for abrasion-induced fracture of the investigated engineered alumina surface (AG sample). Hence, the possibility of abrasion-induced surface fracture and subsequent debris formation at 2-N load is likely to take place to a lower extent. In contrast, high COF and high wear can be explained by the static fatigue wear mechanism. Static fatigue results from a stress-dependent chemical reaction between water vapor and the surface of a ceramic.25 During fretting, the formation of numerous finer cracks can proceed at a faster rate at the root of cracks. The small cracks gradually lengthen, and failure occurs when the cracks are long enough to satisfy the Griffith failure criteria for brittle fracture.26 Thus, two stages of crack growth can be visualized: (1) slow crack growth because of chemical attack at the crack tip and (2) a catastrophic crack growth, when the crack is long enough to satisfy the Griffith criteria. Such moisture-assisted crack propagation and fracture is called static fatigue, also known as stress corrosion cracking. Importantly, the stresses (introduced during static or dynamic contact) at the crack tip control the rate of crack growth.27 The basic crack tip reaction for glass can be as follows:
(H-O-H) + (-Si-O-Si-) → (-Si-OH HO-Si-)
(17.7)
For a given material and the environment, wear due to static fatigue can be reduced with lower residual stresses.28 Hence, the AS sample is expected to be less prone to static fatigue wear. However, the high wear rate (2╯×╯10−4â•›mm3/N·m) of the original alumina surface at 2-N load is presumably due to brittle fracture of the AS surface and simultaneous abrasion. From Equation 17.6 and mechanical properties of the AS specimen, a low load of ∼1â•›N is sufficient to fracture the surface during fretting contact, which produced a significant amount of debris particles after fretting at 2-N load. As a concluding note, it can be stated that the investigation discussed here illustrates that while glass infiltration improves the mechanical properties of Al2O3, such improvement in wear resistance at fretting contact is only limited to low load (2â•›N). At higher load (5, 10â•›N) during fretting, comparable wear resistance is recorded with both GI-Al2O3 and as-sintered Al2O3.
17.7 CONCLUDING REMARKS The research results discussed in this chapter demonstrate better mechanical and tribological performance of the glass-infiltrated alumina over conventional sintered
REFERENCES
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alumina ceramic only at low load (2 N). At lower load, static fatigue wear initiated the fracture, and tribochemical wear directed material removal from glass-infiltrated alumina, whereas abrasion fracture is dominant for as-sintered alumina. The comparable wear resistance of glass-infiltrated alumina with that of as-sintered alumina is attributed to the presence of third-body abrasion and noticeable tribochemical wear, which accommodates the interfacial shear stresses.
REFERENCES ╇ 1â•… R. W. Davidge. Mechanical Behaviour of Ceramics. Cambridge University Press, Cambridge, UK, 1979. ╇ 2â•… B. Lawn. Fracture of Brittle Solids. Cambridge University Press, Cambridge, UK, 1993. ╇ 3â•… M. W. Barsoum. Fundamentals of Ceramics. McGraw-Hill International Editions, 1997. ╇ 4â•… B. R. Marple and D. J. Green. Mullite/alumina particulate composites by infiltration processing: III, mechanical properties. J. Am. Ceram. Soc. 74(10) (1991), 2453–2459. ╇ 5â•… M. Kalin, B. Hockey, and S. Jahanmir. Wear of hydroxyapatite sliding against glass-infiltrated alumina. J. Mater. Res. 18(1) (2003), 27–36. ╇ 6â•… Z. C. Wang, J. Shemilt, and P. Xiao. Fabrication of composite coatings using a combination of electrochemical methods and reaction bonding process. J. Eur. Ceram. Soc. 20 (2000), 1469–1473. ╇ 7â•… B. Giovanni, C. Valeria, L. Luca, M. Tiziano, S. Cristina, B. Cecilia, L. Alessio, and V. Teodoro. Plasma-sprayed glass-ceramic coatings on ceramic tiles: Microstructure, chemical resistance and mechanical properties. J. Eur. Ceram. Soc. 25 (2005), 1835–1853. ╇ 8â•… D. T. Gawne, Z. Qiu, Y. Bao, T. Zhang, and K. Zhang. Abrasive wear resistance of plasma-sprayed glass-composite coatings. J. Therm. Spray Technol. 10(4) (2001), 599–603. ╇ 9â•… D. J. Green, R. Tandon, and V. M. Sglavo. Crack arrest and multiple cracking in glass through the use of designed residual stress profiles. Science 283(5406) (1999), 1295–1297. 10â•… M. Berger and M. Larsson. Mechanical properties of multilayer PVD Ti/TiB2 coatings. Surf. Eng. 16(2) (2000), 122–126. 11â•… J. Gunnars and A. Alahelisten. Thermal stresses in diamond coatings and their influence on coating wear and failure. Surf. Coatings Technol. 80 (1996), 303–312. 12â•… R. Tandon and D. J. Green. Crack stability and T-curves due to macroscopic residual compressive stress profiles. J. Am. Ceram. Soc. 74(8) (1991), 1981–1986. 13â•… B. R. Marple and D. J. Green. Mullite/alumina particulate composites by infiltration processing: IV, residual stress profiles. J. Am. Ceram. Soc. 75(1) (1992), 44–51. 14â•… H. E. Kim, A. J. Moorhead, and S. H. Kim. Strengthening of alumina by formation of a mullite/ glass layer on the surface. J. Am. Ceram. Soc. 80(7) (1997), 1877–1880. 15â•… M. Kalin and S. Jahanmir. Influence of roughness on wear transition in glass-infiltrated alumina. Wear 255 (2003), 669–676. 16â•… M. Kalin, S. Jahanmir, and G. Dražicˇ. Wear mechanisms of glass-infiltrated alumina sliding against alumina in water. J. Am. Ceram. Soc. 88(2) (2005), 346–352. 17â•… D. Sarkar, B. Basu, M. C. Chu, and S. J. Cho. Is glass infiltration beneficial to improve fretting wear properties for alumina? J. Am. Cer. Soc. 90(2) (2007), 523–532. 18â•… M. Chen, K. Kato, and K. Adachi. Friction and wear of self mated SiC and Si3N4 sliding in water. Wear 250 (2001), 246–255. 19â•… M. C. Jeng and L. Y. Yan. Environmental effects on wear behavior of Al2O3. Wear 161 (1993), 11–16. 20â•… M. Kalin, S. Novak, and J. Vižintin. Wear and friction behavior of alumina ceramics in aqueous solutions with different pH. Wear 254 (2003), 1141–1146. 21â•… R. A. Robie, C. B. Finch, and B. S. Hemingway. Heat capacity and entropy of fayalite (Fe2SiO4) between 5.1 and 383 K: Comparison of calorimetric and equilibrium values for the QFM buffer reaction. Am. Mineral. 67 (1982), 463–469.
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22â•… X. Wang and P. Xiao. Characterization of ceramic coatings sintering using residual stress measurements. J. Eur. Ceram. Soc. 24 (2004), 283–288. 23â•… S. G. Roberts. Depths of cracks produced by abrasion of brittle materials. Scr. Mater. 40(1) (1999), 101–108. 24â•… W. D. Kingery, H. K. Bowen, and D. R. Uhlmann. Introduction of Ceramics, 2nd ed. John Wiley & Sons, New York, 1991, 841. 25â•… S. M. Wiederhon. Influence of water vapor on crack propagation in soda-lime glass. J. Am. Ceram. Soc. 50 (1967), 407–414. 26â•… B. Bhushan. Principles and Applications of Tribology. John Wiley & Sons, New York, 1999. 27â•… A. R. C. Westwood. Environment-sensitive fracture of ionic and ceramic solids. Proc. Int. Conf. on Mechanisms of Environmnet Sensitive Cracking of Materials (P. Roland, S. Swann, F. P. Ford, and A. R. C. Westwood, eds), Metals Soc. London, 1974, 283–297. 28â•… B. Bhushan and D. V. Khatavkar. Role of water vapor on the wear of Mn-Zn ferrite heads sliding against magnetic tapes. Wear 202 (1996), 30–34.
CHAPTER
18
TRIBOLOGICAL PROPERTIES OF CERAMIC BIOCOMPOSITES In the Chapter 15, it was clearly mentioned that the tribological properties of hard tissue implant material need to be evaluated before their biomedical application can be realized. Ideally, an implant should survive for the major part of a patient’s lifespan and revision surgery due to intermediate failure should be avoided. Additionally, wear particles, once released from an implant’s surface, can cause cytotoxicity and genotoxicity of biological cells. In the perspective of such broader implications, this chapter discusses tribological properties of hydroxyapatite (HAp)-based bioceramic composites.
18.1 BACKGROUND Hydroxyapatite and various calcium phosphate (CaP)-based biomaterials have been widely researched but the application of monolithic CaP is rather limited because of poor fracture toughness.1–6 Moreover, earlier wear studies have demonstrated that pure HAp does not show acceptable wear resistance, even under water lubrication, for many envisaged applications. A wear factor in the range of 7╯×╯10−6 to 1╯×╯10−5 was reported7 with high dependence on roughness properties.8 Hence, a superior CaP-based biocomposite is being developed with mullite (3Al2O3 2SiO2) reinforcement, which can exhibit single edge V-notched beam (SEVNB) fracture toughness of up to 1.5â•›MPaâ•›m0.5, while showing compressive strength of more than 300â•›MPa. Also, in vitro and in vivo studies on mullite (up to 30 wt%)–reinforced CaP illustrated similar biocompatibility of the composite as that of pure HAp.9,10 Consequently, it is imperative to realize the tribological response of CaP–mullite composite as a potential biomaterial. The tribological studies of a material can be more relevant, if the test conditions are similar to those the material will experience in vivo. Hence, a possible approach to achieve this is to perform fretting wear tests in simulated body fluid (SBF, with or without proteins).
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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Fretting wear is a typical wear mechanism in articulating surfaces, for example, the neck region, taper junction, and knee and hip joints. When such wear occurs, the wear particles can lead to allergy, septic and aseptic loosening of the implant, toxicity, and so on, and this may ultimately require a revision surgery. Hence, fretting damage evaluation of biomaterials is required before total hip arthroplasty is considered, and enhanced tribological properties are necessary for good design of body implants.11 Accordingly, tailoring the implant properties can substantially increase the fretting wear resistance of the implant material, and thus enhance the implant’s life.
18.2 TRIBOLOGICAL PROPERTIES OF MULLITEREINFORCED HYDROXYAPATITE Several studies have reported various tribological mechanisms of diverse implant materials under both dry and lubricated conditions (SBF).12–21 Comparison of these results is difficult due to the variety of parameters used; thus consistent analyses of tribological behavior under the same conditions with a controlled variation in material composition would allow better overview of the major influences and material responses. To achieve this and to attain superior properties, a comprehensive and well-defined set of composites needs to be prepared. Fu et al.22 reported that thermally sprayed HAp coating showed a coefficient of friction (COF) of 0.2–0.3 in the presence of bovine albumin lubrication; unlubricated fretting wear showed a COF value between 0.7 and 0.8. Mullite-containing (10–30â•›wt%) CaP composites were pressureless sintered at 1350°C for 2 hours and, for comparison, monolithic HAp and mullite were sintered at 1200°C for 2 hours and at 1700°C for 4 hours, respectively. The density as well as the phase assemblage in various sintered ceramics, based on x-ray diffraction (XRD) analysis, is summarized in Table 18.1. The influence of tricalcium phosphate (α-TCP or β-TCP) relative to HAp on friction and wear behavior will be analyzed later. Some representative bright field transmission electron microscopy (TEM) images showing the phase assemblage are provided in Figure 18.1. Some characteristic observations include the presence of mullite needles with aspect ratio of 7–8 as well as the presence of gehlenite at the triple pockets of HAp/TCP/mullite. A more detailed description of the phase assemblage can be found elsewhere.10
18.3 FRICTION AND WEAR RATE To understand the wear under severe conditions, a harder counterbody needs to be used; to illustrate this, ZrO2 balls (Hvâ•›≈â•›12â•›GPa) were used in the fretting tests against sintered CaP–mullite biocomposites (Hvâ•›≈â•›4â•›GPa) in both ambient and SBF environments. The results presented here are based on the fretting tests carried out at 5-Hz frequency swapping a linear distance of 80â•›µm at the normal load of 10â•›N. A total of 100,000 cycles were given and the frictional response was monitored in situ to evaluate the fretting wear resistance of the HAp–mullite composites.
18.3 Friction and Wear Rate
â•… 289
TABLE 18.1â•… Summary of the Sinter Density and Phase Assemblage (Determined on the Basis of the X-Ray Peak Intensities) in Various Pressureless Sintered HAp-Mullite Biocomposites17
Samples
Maximum achievable Densification (w.r.t. initial powder mixture) %
Sintering Conditions
Phases present
Pure HAp HAp–10M
99.17 94
1200°C for 2 hours 1350°C for 2 hours
HAp–20M
98.1
1350°C for 2 hours
HAp–30M
95.6
1350°C for 2 hours
Pure Mullite
98.7
1700°C for 4 hours
HAp-ss α-TCP-ss, HAp-ss, β-TCP-w, mullite-ww, CaO-ww α-TCP-s, β-TCP-ss, HAp-ww, mullite-w, gehlenite-ww, CaO-ww, alumina-ww β-TCP-ss, HAp-ww, mullite-s Gehlenite-ww, CaO-ww, mlumina-s Mullite-ss
ss, very strong; s, strong; ww, very weak; w, weak.
After fretting, a laser surface profilometer can be used to scan the worn surface to estimate wear volume. Coefficient of friction (COF) is plotted in Figure 18.2. The recorded wear resistance of the composites with varying mullite content is compared in Figure 18.3. A comparison of friction plots in Figure 18.2a reveals that pure HAp exhibited COF of 0.35, whereas a much higher COF of 0.62 was recorded with pure mullite. All the HAp–mullite composites experienced COF within this rage (∼0.5) during the steady-state condition. However, noticeable fluctuations appear to be caused by three-body wear, or entrapment of debris particles during fretting. Similar to the COF values observed under dry conditions (Fig. 18.2a), the lowest COF (0.3) was exhibited by pure HAp samples under wet conditions (Fig. 18.2b). Overall, a lower COF (0.5) can be obtained with CaP–mullite composites and stable frictional behavior was observed.
18.3.1â•… Effect of Simulated-Body-Fluid Medium on Wear of Mullite-Reinforced Hydroxyapatite Some interesting observations can be made when wear rate is plotted along with hardness, as shown in Figure 18.3. While the wear rate does not correlate with the corresponding hardness under dry conditions, there is an inverse proportionality of wear rate with hardness under lubricated conditions. The maximum wear rate (8.54╯×╯10−6â•›mm3/N m) was measured for the HAp–10â•›M sample, whereas pure mullite experienced the lowest wear rate (1.22╯×╯10−6â•›mm3/Nâ•›m). Under both dry and SBF conditions, pure mullite exhibited the lowest wear rate (0.94╯×╯10−6â•›mm3/Nâ•›m) and this could be due to higher hardness. The decrease
290â•…
CHAPTER 18â•… Tribological Properties of Ceramic Biocomposites
M
ul
lit
CaO
e
TCP HAp
Gehlenite 200 nm
TCP
e nit
Mullite
Ge
hle
Al
um
in
a
(a)
TCP
Alumina
200 nm (b)
Figure 18.1â•… (a) Bright field TEM images revealing presence of CaO, gehlenite, and mullite needles in pressureless sintered biocomposites: (a) HAp–(10 wt%)mullite and (b) HAp–(30 wt%)mullite.10
18.3 Friction and Wear Rate
â•… 291
0.7 0.6
COF
0.5 0.4 0.3 Pure HAp HAp-10M HAp-20M HAp-30M Pure Mullite
0.2 0.1 0.0 0
20,000
40,000 60,000 Cycles (a)
0
20,000
40,000 60,000 Cycles (b)
80,000
100,000
0.7 0.6
COF
0.5 0.4 0.3 0.2 0.1 0.0 80,000
100,000
Figure 18.2â•… Plots of coefficient of friction (COF) versus number of fretting cycles for HAp–mullite composites (a) under dry ambient condition (b) in simulated body fluid with albumin protein environment.17
in the wear rate under the lubricating SBF condition was noticeable for pure HAp and HAp–10â•›M. Overall, the wear rate varies in the order of 10−6â•›mm3/Nâ•›m and this reflects reasonably good wear resistance of the CaP–mullite composites. Also, the mullite reinforcement importantly increases the wear rate with respect to monolithic HAp and this aspect is discussed later. The 2D surface depth profiles after fretting under dry and SBF conditions are shown in Figure 18.4a,b, respectively.
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20
Fretted in air Fretted in SBF+Albumin
18
hardness
8
16 14
Hardness (GPa)
Wear Rate (¥10-6mm3/Nm)
10
12
6
10 4
8 6 4
2
2 0
0 Pure HAp HAp10M
HAp20M
HAp30M Pure Mullite
Figure 18.3â•… Wear rate of different samples (pure HAp, HAp–mullite composites, and pure mullite) are plotted when fretted under dry conditions as well as in SBF╯+╯albumin medium. Around 10% deviation of the mean wear rate was measured in our experiments.17 20.00 µm
Air
20.00
Pure HAp
SBF
µm
0.00
Pure HAp
0.00
−20.00
−20.00 1.67 mm
0.00 mm 16.00
0.00 mm 10.00
HAp-10M
µm
HAp-10M
2.00 mm
µm
0.00
0.00
−16.00
−10.00 2.30 mm
0.00 mm 16.00
10.00
HAp-20M
µm
0.00 mm
HAp-20M
2.00 mm
µm 0.00
0.00
−16.00
−10.00 0.00 mm
20.00
6.00
HAp-30M
µm
0.00 mm
2.35 mm
2.00 mm
HAp-30M
µm 0.00
0.00
−6.00
−20.00 3.00 mm
0.00 mm 4.00
Pure Mullite
µm
4.00
2.50 mm
0.00 mm
Pure Mullite
µm 0.00
0.00
−4.00
−4.00 0.00 mm
(a)
1.86 mm
0.00 mm
(b)
2.00 mm
Figure 18.4â•… Two-dimensional wear depth profiles for pure HAp, HAp–30â•›M, and pure mullite after fretting wear tests under dry and SBF conditions.17
18.3 Friction and Wear Rate
â•… 293
Figure 18.4a,b reveal that the pure HAp surface is very rough, independent of tribological environment, whereas the addition of mullite to HAp matrix elicits a smooth surface appearance after testing in SBF. On one hand, where pure HAp elicits maximum wear depth of 20â•›µm, the wear depth decreases to ∼16â•›µm with mullite addition, and to a much lower value (∼2.5â•›µm) in the case of pure mullite during fretting under dry conditions. A common observation was that higher wear depths were measured under dry conditions compared with those under SBF medium.
18.3.2â•…Surface Topography of Mullite-Reinforced Hydroxyapatite after Fretting Wear Scanning electron microscopy (SEM) images revealing surface topography of pure HAp, HAp–20â•›M, and HAp–30â•›M samples are presented in Figure 18.5a–e, respectively, showing damage caused by fretting under dry conditions. The delamination of the HAp surface layer (Fig. 18.5a) indicates severe damage due to plowing and microcrack formation (Fig. 18.5b). Fatigue cracks were observed in the HAp–20â•›M sample due to abrasion during fretting (Fig. 18.5c,d), but a reduced severity of cracking is observed in the HAp–30â•›M sample (Fig. 18.5f). As for the HAp–10â•›M sample, severe wear and cracking were observed (not shown here), whereas only abrasive wear scratches were observed for the pure mullite sample. The dominance of abrasive wear was observed after fretting under SBF conditions, for 10% and 20% mullite composites (Fig. 18.6a,b), abrasive scratches with mild fatigue cracking were observed in the damaged zone, whereas plowing and grain pullout were dominant in the HAp–30â•›M sample (Fig. 18.6c). The large delaminated flakes with the so called “mud” cracks are a result of the postwear drying process of the hydrated layers, typically during sample preparation and SEM analyses, as reported by Kalin et al.7 The wear depth values presented in Figure 18.4 do not exactly follow the wear rate trends, since wear rate depends on both depth and diameter of wear scar. Lower depth but larger wear scar may produce higher wear rate. For HAp–20â•›M and HAp–30â•›M samples, the wear rate values were higher due to larger wear scar diameter (Table 18.2) despite showing lower fretting wear depth under SBF conditions than under dry conditions.
18.3.3â•…Frictional Behavior In this subsection, an attempt has been made to explain the difference in friction behavior, as shown in Figure 18.2. Pure HAp and pure mullite possess the lowest and highest COF values, respectively, and all three composites show COF values intermediate to these. In dry conditions, the formation of wear debris and their escape or displacement from the contact zone during fretting can cause fluctuations in COF, which is evident in Figure 18.2a. On the other hand, friction is quite smooth and steady under lubricated conditions, with only minimal friction fluctuations (except for pure HAp) as evident in Figure 18.2b. It has to be mentioned that fretting wear has a characteristic potential of forming considerable amounts of very fine oxide debris and these wear debris particles often govern the wear of the
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500 µm
20 µm
(a)
(b) P
Ca
Full Frame O
Al Si
Ca
400 µm
20 µm
(c)
(d) Full Frame O
P
Ca
Al Si
Ca
500 µm (e)
20 µm (f )
Figure 18.5â•… SEM images of worn surfaces of (a,b) pure HAp sintered at 1200°C, (c,d) HAp–20M, and (e,f) HAp–30â•›M sintered at 1350°C, after testing against zirconia in air. The double-pointed arrow indicates fretting direction.17
materials.23,24 Young’s moduli of HAp–xM composites are inferior compared with those of pure HAp (∼120â•›GPa) and pure mullite (240â•›GPa) samples. Due to these lower elastic modulus values (∼45–70â•›GPa), the initial contact diameter was larger (Table 18.2) than contact diameters of other materials. The strong asperity contact between mullite and a zirconia counterbody was responsible for the highest COF
18.3 Friction and Wear Rate P Full Frame
â•… 295 Ca
P
Ca
Full Frame Al
O
O
Al
Si
Si
Ca
Ca
100 µm
100 µm (a)
(b) Full Frame
P
Ca
Full Frame
Al
Al O
Si
100 µm (c)
Si
O
Ca
100 µm (d)
Figure 18.6â•… SEM images of worn surfaces of (a) HAp–10â•›M sintered at 1350°C, (b) HAp–20â•›M, sintered at 1350°C, (c) HAp–30â•›M sintered at 1350°C, and (d) pure mullite sintered at 1700°C, after testing against zirconia in SBF medium containing albumin. The double-pointed arrow indicates fretting direction.17
for pure mullite samples. The data in Table 18.2 also indicate that the tests were conducted with variation in contact stress of ∼ 809–1038â•›MPa. This kind of contact pressure is much higher than the upper end of the commonly experienced contact pressure at various orthopedic joints (hip joint, maximum contact pressure ∼25â•›MPa).25 The COF values seem to be quite unexpected in the SBF environment. The highest COF values recorded for a HAp–10â•›M sample can be attributed to the dominant presence of α-TCP phase (Table 18.1). In the presence of SBF medium and constant fretting action, α-TCP can be dissolved locally at the contact and, therefore, causes localized spalling, leading to an increase in COF. This type of phenomenon was not observed in the case of HAp–20â•›M and HAp–30â•›M samples as the major phase is β-TCP. It is already well known that β-TCP is more stable than α-TCP in body fluid. Except for HAp–10â•›M samples, the lowering of COF in SBF condition compared with dry contact can be attributed to lubrication effects in SBF solution. Considering various parameters, the liquid film at the contact would be in the
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TABLE 18.2â•… Maximum and Mean Hertzian Contact Pressures as well as the Initial Hertzian Contact Diameter for the Investigated Fretting Couple (Containing Pure HAp, Pure Mullite, and HAp–mullite Composites) Are Presented along with Measured Wear Scar Diameter17
Material
Medium
Pure HAp
Dry SBF Dry SBF Dry SBF Dry SBF Dry SBF
HAp–10â•›M HAp–20â•›M HAp–30â•›M Pure mullite
Maximum Hertzian contact pressure (initial) (MPa)
Mean Hertzian contact pressure (initial) (MPa)
Initial Hertzian contact diameter (µm)
Measured wear scar diameter (µm)
E-Modulus (GPa)
809 809 521 521 621 621 597 597 1038 1038
540 540 347 347 414 414 398 398 692 692
153.7 153.7 191.6 191.6 175.4 175.4 178.9 178.9 135.7 135.7
767 377 801 671 564 636 563 614 530 344
120 45 72 68 240
boundary regime and hence real asperity contact cannot be avoided in spite of the presence of a liquid film.
18.3.4â•… Wear Micromechanisms of Hydroxyapatite-Based Materials in Simulated Body Fluid In this section, the wear mechanisms are analyzed on the basis of wear scar topographical observations and Hertzian contact stress conditions. An effort is made to discuss the effect of fretting environment (ambient and SBF lubrication) on the wear mechanisms. Following the Hertzian contact mechanics theory,26 the maximum and mean contact pressures along with initial Hertzian contact diameter are presented in Table 18.2. Hertzian contact pressure and contact diameter are mainly functions of elastic modulus of the mating materials. Having highest E-modulus value, mullite possesses very high contact pressure (1038â•›MPa) and the lowest contact diameter (136â•›µm). The contact pressure is the lowest (521â•›MPa) and the diameter is the highest (192â•›µm) for HAp–10â•›M composite. The low elastic modulus tends to reduce contact pressure by increasing the contact diameter. After the fretting experiments, the wear scar diameters were increased due to the oscillatory motion and consequently, the contact pressures were reduced. In dry contact, the fretted contact diameters (Table 18.2) are proportional to initial contact diameter. However, in SBF medium discrepancy was found for HAp–20â•›M and HAp–30â•›M samples. In these two cases, the fretted wear scar diameters were higher than those of dry contact. Therefore, the higher wear rate for these two composites in SBF medium than under dry conditions can
18.3 Friction and Wear Rate
â•… 297
be understood. These two composites mainly contained β-TCP, and the wear resistance of β-TCP is inferior in SBF medium. In dry contact, the wear of pure HAp mainly occurred by delamination and microcrack formation. The high contact pressure (809â•›MPa) was sufficient to initiate the microcracks at the contact region. Also, being very brittle material (toughness ∼ 0.6â•›MPa m0.5), HAp could easily form microcracks at the fretting contact. The microcracks were also generated in the subsurface region. Due to prolonged cyclic loading, the microcracks grew and combined with each other, culminating in delamination of the surface layer. For HAp–10â•›M and HAp–20â•›M composites, the fatigue cracks were formed due to the cyclic loading over prolonged time (Fig. 18.6). For HAp– 30â•›M composites, the extent of fatigue cracks was much lower. Here, the grains were pulled out mainly due to plowing by counterbody asperities. For pure mullite, the counterbody asperities plastically deformed the surface, resulting in abrasive scratches due to very high contact pressure. Under SBF conditions, (Fig. 18.6) the contact damage was considerably less than that of dry contact. Due to the presence of liquid film, only abrasion (except for HAp–30â•›M) and mild plowing occurred. However, the wear debris formed under SBF conditions was constantly flushed off due to the rapid vibrating motion at contacts. For the samples containing TCP phase, such an effect was more prominent as the TCP phase has higher solubility than pure HAp or mullite. Therefore, the wear rate was increased to some extent. Even for HAp–20â•›M and HAp–30â•›M samples, the wear rates exceeded the value obtained in dry contact. The effect of albumin present in SBF needs to be analyzed. As a direct observation from SEM, it is revealed that albumin did not adhere or got absorbed on the fretted zone. It is possible that due to continuous reciprocatory motion at the fretting contact, it was difficult for protein molecules to be adhered on the surface. Also, if there was some adherence that might be very loosely bonded, which was washed away during cleaning of samples after the experiments. However, albumin might play an important role in controlling the wear rate and COF during fretting experiments. Marques et al.27 studied the effect of albumin on the mineralization process in Hank’s balanced salt solution (HBSS). From their results, it was revealed that albumin helps in TCP dissolution due to its Ca affinity in aqueous HBSS solution. For a similar reason, serum albumin inhibited apatite crystal growth on the sample surface in SBF solution.28–30 Their results showed that, at concentration of 40â•›g/L, bovine serum albumin has a significant retardation effect on apatite precipitation from SBF onto pure or fluoridated HAp coatings. It was concluded that, depending on the surface composition, the apatite formation kinetics even became faster. From this discussion, it is clear that albumin in general retards apatite formation on the sample surface in SBF. Also, an apatite layer could act as a protective layer against further material removal. Therefore, the wear rate of CaP-containing materials in SBF was found to be higher than expected, due to such inhibitory effect. However, in some places on the fretted surface the apatite layer was found to be adhered discretely and later, upon drying, the layer was cracked. SEM images probably show the presence of apatite with the signature of cracking due to drying. For pure mullite, this kind of effect was not observed as pure mullite cannot induce any apatite layer formation when immersed in SBF solution.
298â•…
CHAPTER 18â•… Tribological Properties of Ceramic Biocomposites
On the other hand, tribochemical reactions of pure HAp in a distilled-water environment were reported to play a major role in wear behavior.7,8 It was shown using SEM and TEM observations that a “crystalline-to-amorphous transition” occurs during wear of HAp in water and forms a third-body interface boundary layer, which importantly affects the wear and friction. This amorphous surface layer could form as a direct result of extreme, localized plastic deformation or as the result of an irreversible transition of hexagonal HAp to a high-pressure phase that is unstable, except at the peak contact pressures. As an alternative, the amorphous wear products (tribolayer) could be the result of rapid solution–reprecipitation (see Fig. 18.7). In this case, accelerated dissolution of HAp into water occurs within the contact area. Since the two surfaces are in constant contact, localized saturation and subsequent precipitation is possible within the contact area. In either case, the presence of largescale pores within the amorphous particles suggests that wear products are in a hydrated form (possibly gel-like) during the active wear process, and the pores form during drying as the water evaporates. This hydrated form appears sufficiently attached or bonded at the interface to form a quasi-stable “matrix” that also contains the crystalline wear debris particulates. The continued presence of such a debris layer during active wear is, of course, necessary for three-body wear.7,8 A comparison of the tribological results obtained with pure HAp and HApbased composites, as reported in the literature, is presented in Table 18.3. Broadly, the wear rate varies between 10−5 and 10−9â•›mm3/Nâ•›m depending on operating parameters. Differences between COF and wear rate, as observed from Table 18.3, can be explained on the basis of difference in tribological environment as well as counterbody. Although the coefficient of friction as recorded with HAp–mullite composites was on the higher side compared with other competing materials, the wear rate, varying on the order of 10−6â•›mm3/Nâ•›m, showed good wear-resistance properties. In summary, mild abrasion, fracture, and fatigue wear were reported as dominant wear mechanisms for different HAp-based composites. It must also be mentioned that rise in contact temperature can also play an important role in adversely affecting the tribology of HAp composites. In some studies on various sets of biomaterials under dry and SBF conditions,19,21 it has already been shown that the contact temperature rise was rather low in fretting contact due to its lower sliding velocity, although in some single contact spots the temperature might increase (depending on the conditions). This statement is certainly correct, when full apparent contact area is considered.31–34 Hence, the temperature would not be a major factor in the case of fretting contacts for CaP– mullite composites.
18.4 CONCLUDING REMARKS Based on the discussion in this chapter, the following key observations emerge: The steady-state COF values for pure HAp were always lower (0.30–0.35), irrespective of the fretting environment. In the SBF environment, HAp–(10â•›wt%) mullite composite shows the highest COF value (0.55). In both dry and SBF environments, the wear rate varied within the same order of magnitude (10−6â•›mm3/Nâ•›m).
18.4 Concluding Remarks
â•… 299
Electron Diffraction
nano-crystalline “ring” pattern
100 nm
(a) Electron Diffraction
diffuse “ring” pattern smorphous material
200 nm
(b)
Figure 18.7â•… (a) Polycrystalline wear debris particle; insert is characteristic electron diffraction pattern. Note nanoscale crystalline structure. (b) Amorphous debris; insert is diffuse ring electron diffraction pattern characteristic of amorphous materials. Note extensive porosity within the amorphous particle.7
No systematic trend between wear rate (dry contact) and hardness was observed. In the SBF environment, the wear rate was noticeably reduced in the case of pure HAp, mullite, and HAp–10â•›M composite. The wear depth varied between 2 and 20â•›µm, with higher depth being consistently measured in the case of dry contact. In dry contact, the wear mechanism was mainly guided by microcracking, delamination, plowing, and fatigue cracking. However, in SBF medium, the major wear mechanism for the mullite-containing HAp composites was found to be mild abrasion and/or plowing, eventually leading to mild fracture and removal of
300
Counterbody
UHMWPE
Stainless steel
ZrO2
Unknown
Stainless steel
Material
HAp–PSZ
HAp–collagen with 10% Hyaluronic acid
HAp–CNT
HAp-NFSS
HDPE–HAp– Al2O3
Hank’s balanced salt solution (HBSS)
Air
Load: 10â•›N
0.07–0.11
0.35–0.75 (45â•›N)
Load: 15, 45â•›N Ball-on-flat (Fretting)
0.5–0.85 (15â•›N)
Pin-on-disk
Load: 8.8â•›N
—
Pin-on-disk
c-SBF
0.01–0.03
Load: 10–70â•›N
Bovine serum
0.01–0.04
0.04–0.06
COF
Pin-on-disk
Pin-on-disk Load: 4.3â•›N
Experimental conditions
Carboxymethyle cellulose solution
Human plasma
Medium of testing
Abrasive wear at low load; Abrasive, Adhesive and third-body mechanism at higher load Mild abrasive wear and limited plastic deformation
—
2.3╯±â•¯0.7╯×╯10−6â•›mm3/Nâ•›m
38.9â•›g m−2/104 revolutions
Abrasion, fracture, plastic deformation fragmentation, chipping and plowing
Deformation and fracture of HA layer, Transfer layer buildup
4.1╯×╯10−5â•›mm3/Nâ•›m
—
Fatigue wear
Wear mechanisms
5.9╯±â•¯0.7╯×╯10−9â•›mm3/Nâ•›m
Wear rate
TABLE 18.3â•… Summary of Literature Results Reporting the Tribological Properties of Several HAp-Based Materials17
301
Al2O3
ZrO2
ZrO2
Hardened SS
Al2O3 (with infiltrated glass)
HDPE–HAp– Al2O3
HDPE–HAp– Al2O3
HAp–(30 Wt%)mullite
HAp
HAp
water
Dry
SBF
SBF
SBF
Medium of testing
Load: 10â•›N, 20â•›N, 30â•›N
Ball-on-disk
Load: 10â•›N
Ball-on-flat
Load: 10â•›N
Ball-on-flat (fretting)
Load: 10â•›N
Ball-on-flat (fretting)
Load: 10â•›N
Ball-on-flat (fretting)
Experimental conditions
0.65
0.7–0.8
0.4
0.05
0.11
COF
7╯×╯10−6â•›mm3/Nâ•›m −1.1╯×╯10−5â•›mm3/Nâ•›m
1.36–11.5╯×╯10−6â•›mm3/Nâ•›m
5╯×╯10−6â•›mm3/Nâ•›m
1.1╯±â•¯0.7╯×╯10−6â•›mm3/Nâ•›m
5.9╯±â•¯0.9╯×╯10−7â•›mm3/N m
Wear rate
UHMWPE, ultra-high-molecular-weight polyethylene; PSS, partially stabilized ZrO2, CNT, carbon nanotube; NFSS, nickel-free stainless steel.
Counterbody
Material
Fracture and deformation, with tribochemical film and phase transformation
Delamination and abrasive wear
Mild plowing
Mild Abrasive wear and limited plastic deformation
Mild abrasive wear and limited plastic deformation
Wear mechanisms
302â•…
CHAPTER 18â•… Tribological Properties of Ceramic Biocomposites
the grains. Based on comparison among HAp–mullite composite materials, HAp– (30â•›wt%)mullite ceramics appears to have the best combination of COF and wear resistance under dry and SBF conditions.
REFERENCES ╇ 1â•… S. Gautier, E. Champion, and D. B. Assollant. Processing, microstructure and toughness of Al2O3 platelet-reinforced hydroxyapatite. J. Eur. Ceram. Soc. 17 (1997), 1361–1369. ╇ 2â•… J. Li, B. Fartash, and L. Hermansson. Hydroxyapatite–alumina composites and bone-bonding. Biomaterials 16 (1995), 417–422. ╇ 3â•… R. R. Rao and T. S. Kannan. Synthesis and sintering of hydroxyapatite–zirconia composites. Mater. Sci. Eng. C 20 (2002), 187–193. ╇ 4â•… V. V. Silva, F. S. Lameiras, and R. Z. Domínguez. Microstructural and mechanical study of zirconiahydroxyapatite (ZH) composite ceramics for biomedical applications. Comp. Sci. Tech. 61 (2001), 301–310. ╇ 5â•… G. Gollera, H. Demirkıran, F. N. Oktar, and E. Demirkesen. Processing and characterization of bioglass reinforced hydroxyapatite composites. Ceram. Inter. 29 (2003), 721–724. ╇ 6â•… W. Suchanek, M. Yashima, M. Kakihana, and M. Yoshimura. Hydroxyapatite/hydroxyapatite-whisker composites without sintering additives: Mechanical properties and microstructural evolution. J. Am. Ceram. Soc. 80 (1997), 2805–2813. ╇ 7â•… M. Kalin, B. Hockey, and S. Jahanmir. Wear of hydroxyapatite sliding against glass-infiltrated alumina. J. Mater. Res. 18(1) (2003), 27–36. ╇ 8â•… M. Kalin, S. Jahanmir, and L. K. Ives. Effect of counterface roughness on abrasive wear of hydroxyapatite. Wear 252 (2002), 679. ╇ 9â•… (a) S. Nath. Development of novel calcium phosphate-mullite composites for orthopedic applications, PhD thesis, IIT Kanpur, India, 2008. (b) S. Nath, B. Basu, M. Mohanty, P. V. Mohanan. In vivo response of novel Hydroxyapatite-mullite composites: Results up to 12 weeks of implantation. J. Biomed. Mater. Res. B Appl. Biomater. 90B (2009), 547–57. 10â•… S. Nath, K. Biswas, and B. Basu. Phase stability and microstructure development in hydroxyapatite– mullite system. Scr. Mater. 58 (2008), 1054–1057. 11â•… P. Schaaff. The role of fretting damage in total hip arthroplasty with modular design hip joints— Evaluation of retrieval studies and experimental simulation methods. J. Appl. Biomater. Biomech. 2 (2004), 121–135. 12â•… A. Choubey, B. Basu, and R. Balasubramaniam. Tribological behavior of Ti-based alloys in simulated body fluid solution at fretting contacts. Mater. Sci. Eng. A 379 (2004), 234–239. 13â•… D. Sheeja, B. K. Tay, S. P. Lau, and L. N. Nung. Tribological characterization of diamond-like carbon coatings on Co–Cr–Mo alloy for orthopedic applications. Surf. Coatings Technol. 146–147 (2001), 410–416. 14â•… D. Sheeja, B. K. Tay, and L. N. Nung. Tribological characterization of surface modified UHMWPE against DLC-coated Co–Cr–Mo. Surf. Coatings Technol. 190 (2005), 231–237. 15â•… R. Hauert and U. Müller. An overview on tailored tribological and biological behavior of diamondlike carbon. Diam. Relat. Mater. 12 (2003), 171–177. 16â•… S. Nath, S. Bodhak, and B. Basu. Tribological investigation of novel HDPE-HAp-Al2O3 hybrid biocomposites against steel under dry and simulated body fluid condition. J. Biomed. Mater. Res A 83A (2007), 191–208. 17â•… S. Nath, R. Ummethala, and B. Basu. Fretting wear behavior of calcium phosphate-mullite composites in dry and albumin-containing simulated body fluid conditions. J. Mater Sci. Mater. Med. 21 (2010), 1151–1161. 18â•… S. Bodhak, S. Nath, and B. Basu. Fretting wear properties of hydroxyapatite, alumina containing high density polyethylene biocomposites against zirconia. J. Biomed. Mater. Res A 85A (2008), 83–98.
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19â•… S. Nath, S. Bajaj, and B. Basu. Microwave-sintered MgO-doped zirconia with improved mechanical and tribological properties. Int. J. Appl. Ceram. Tech. 5 (2008), 49–62. 20â•… S. Nath, S. Bodhak, and B. Basu. HDPE-Al2O3-HAp composites for biomedical applications: Processing and characterization. J. Biomed Mater. Res. B Appl. Biomater. 88B (2009), 1–11. 21â•… S. Bodhak, S. Nath, and B. Basu. Friction and wear properties of novel HDPE-HAp-Al2O3 composites against alumina counterface. J. Biomater. Appl. 23 (2009), 407–433. 22â•… Y. Fu, A. W. Batchelor, and K. A. Khor. Fretting wear behavior of thermal sprayed hydroxyapatite coating lubricated with bovine albumin. Wear 230 (1999), 98–102. 23â•… M. Varenberg, G. Halperin, and I. Etsion. Different aspects of the role of wear debris in fretting wear. Wear 252 (2002), 902–910. 24â•… I. L. Singer. How third-body processes affect friction and wear. MRS Bull. 23 (1998), 37–40. 25â•… M. Kumagai, Y. H. Kim, N. Inoue, E. Genda, K. Hua, B. T. L. Liong, T. Koo, and Y. Chao. 3-D dynamic hip contact pressure distribution in daily activities. Summer Bioengineering Conference, June 25–29, Sonesta Beach Resort in Key Biscayne, Florida 2003, 53–54. 26â•… B. Bhushan. Principles and Applications of Tribology. A wiley-Interscience Publication, John Wiley & Sons, INC., New York, 1999, 201–202. 27â•… P. A. A. P. Marques, A. P. Serro, B. J. Saramago, A. C. Fernandes, M. C. F. Magalhães, and R. N. Correia. Mineralisation of two phosphate ceramics in HBSS: Role of albumin. Biomaterials 24(3) (2003), 451–460. 28â•… H. Gilman and D. W. L. Hukins. Seeded growth of hydroxyapatite in the presence of dissolved albumin at constant composition. J. Inorg. Biochem. 55 (1994), 31–39. 29â•… J. Garnett and P. Dieppe. The effects of serum and human albumin on calcium hydroxyapatite crystal growth. Biochem. J. 266 (1990), 863–868. 30â•… A. P. Serro, A. C. Fernandes, B. Saramago, J. Lima, and M. A. Barbosa. Apatite deposition on titanium surfaces—the role of albumin adsorption. Biomaterials 16 (1997), 963–966. 31â•… M. Kalin and J. Vizintin. Comparison of different theoretical models for flash temperature calculation under fretting conditions. Tribol. Int. 34 (2001), 831–839. 32â•… M. Kalin and J. Vizintin. High temperature phase transformations under fretting conditions. Wear 249 (2001), 172–181. 33â•… M. Kalin and J. Vizintin. A tentative explanation for the tribochemical effects in fretting wear. Wear 250 (2001), 681–689. 34â•… M. Kalin. Influence of flash temperatures on the tribological behaviour in low-speed sliding: A review. Mater. Sci. Eng. A 374 (2004), 390–397.
SECTION
IV
FRICTION AND WEAR OF NANOCERAMICS
CHAPTER
19
OVERVIEW: NANOCERAMIC COMPOSITES Nanostructured ceramics and composites have been attracting wider attention, and an extensive review of this type of materials was published in 2007.1 This chapter presents some salient parts of this review, which is intended to act as an introduction to the unique property combination that one can achieve with this unique class of materials. Also, the processing of these materials requires the use of advanced sintering techniques and this is also discussed in this overview chapter. The microstructural characteristics of some of the developed nanoceramics and nanocomposites are briefly presented. An overview of tribological properties of a few nanostructured ceramics is also provided.
19.1 INTRODUCTION Conventionally, nanostructured materials are defined as materials with structural units having a size scale of less than a hundred nanometers in any dimension, and such materials have a unique combination of physical or functional properties, which cannot be obtained in materials with structural units having length scale in microns or larger. This length scale is typically particle diameter, grain size, layer thickness, or the width of a conducting line on an electronic chip. Based on length scale, nanostructured materials can be classified as zero-dimensional (nanosized powders), one-dimensional (nanocrystalline multilayer), two-dimensional (filamentary rods of nanoscaled thickness), and three-dimensional (bulk materials with at least one nanocrystalline phase).2,3 A wide variety of applications have been proposed for bulk nanoceramics and nanoceramic composites, such as durable ceramic parts for automotive engines, cutting tools, heat engine components, wear-resistance parts, aerospace-related industrial applications, ultrafine filters, flexible superconducting wire, and fiber-optic connector components. The major challenge in nanoceramic development is the restriction of grain growth during processing, which is often found to be difficult using conventional
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
307
308â•…
CHAPTER 19â•… Overview: Nanoceramic Composites
(a) Intergranular nanocomposite
GB pinning: improved creep resistance (c) Inter/intragranular nanocomposite
Good creep resistance and high σb & KIC
(b) Intragranular nanocomposite
Transgranular fracture: high σb & KIC (d) Nano/nano composite
GB sliding: superplasticity
(e) Nano/micro composite
Figure 19.1â•… Schematic representation of the microstructural features of the various nanocomposites as well as nano/nano composites and nano/micron composites.1
sintering techniques. The advanced sintering techniques, in particular spark plasma sintering (SPS), sinter–hot isostatic pressing (sinter-HIPing), are some of the wellknown and successful laboratory-scale processes for synthesizing bulk nanomaterials. Among them, SPS is currently one of the most widely used processing routes for developing bulk nanostructured ceramics. As pioneered by Niihara,4 significant improvement in mechanical properties, in particular strength and hardness, can be achieved in ceramics by nanocomposite design. Ceramic nanocomposites can also be classified, based on their microstructural design, depending on whether the matrix or the reinforcement or both are nanocrystalline as well as on the distribution of the nanocrystalline reinforcement (Fig. 19.1). Accordingly, specific improvements in mechanical properties may be obtained via microstructural designing, and the development of ceramic materials by nanocomposite design results in improvement of not only the room-temperature mechanical properties such as hardness and strength, but also high-temperature mechanical properties such as hot hardness, high-temperature strength creep, and fatigue fracture resistance.1
â•… 309
19.3 Overview of Developed Nanoceramics and Ceramic Nanocomposites
19.2 PROCESSING OF BULK NANOCRYSTALLINE CERAMICS The major processing-related challenges in ceramic nanocomposite development are schematically shown in Figure 19.2a. Although various powder synthesis routes are now available to produce nanocrystalline ceramic powders, the success in consolidation of nanopowders to yield bulk nanocomposites lies with the precise control of the coarsening–densification competition during the sintering process. A combination of much lower sintering temperature and shorter sintering time, in short, “activated sintering,” is adopted in order to realize processing of nanocomposites. Fundamentally, activated sintering results in the augmentation of densification with reduced sintering time and temperature. The enhanced duration of elevated temperature exposure in the case of pressureless sintering results in significant grain growth, especially in the last stage of sintering.5 The enhancement of the densification rate can be achieved to some extent, by application of external pressure during sintering (hot pressing, sinter forging, sinter-HIPing, and SPS). Hot pressing and sinter forging involve the application of uniaxial pressure to a powder compact in the presence or absence of a constraining die, respectively. A variant of activated sintering, known as field-assisted sintering technique (FAST), involves the imposition of an electric field simultaneously with applied pressure during densification. FAST is known under different names, such as SPS and plasma activated sintering (PAS). From the phenomenological point of view, SPS is similar to conventional hot pressing in the sense that sintering occurs upon application of uniaxial mechanical pressure, simultaneously during heating. However, the main difference is that, in SPS, the heating is achieved by passing current through the electrically conducting pressure dies (typically graphite dies) and, in the case of conducting powders, through the powders as well. This is considered to have a major influence on the ability of SPS to consolidate even some difficult-to-sinter ceramic powders.6 A schematic view of the current flow during SPS is shown in Figure 19.2b. Based on the electrical conductivity of the ceramic powders, a fraction of the total current will always pass through the porous powder compact (path 1) and the remaining fraction will flow via the graphite die walls (path 2). In practice, the external pressure is applied throughout the consolidation process, and it may be kept constant throughout or varied from an initial light pressure (10–15â•›MPa) to a higher pressure (up to ∼100â•›MPa) during heating. From a mechanistic point of view, the consolidation process can be divided into two overlapping stages, which occur simultaneously or in sequence: (1) initial activation of the powders by the application of a pulsed voltage, and (2) resistance sintering under pressure by application of a steady DC current.
19.3 OVERVIEW OF DEVELOPED NANOCERAMICS AND CERAMIC NANOCOMPOSITES With a view to exploring the postulated superiority of ceramic nanomaterials in terms of better material properties, considerable efforts are being invested in
310â•…
CHAPTER 19â•… Overview: Nanoceramic Composites
Compaction of nanoparticles to obtain green body without cracks/density gradients
Cost-effective synthesis route to obtain non-agglomerated nanocrystalline ceramic powders
Processing of Ceramic Nanocomposites
Retention of nanosized grains/particulates and dispersion of nano-reinforcements in matrix
Inhibition of grain growth while achieving full densification
(a) Pressure
DC Pulse Generator
Pulse Current
Powder
Thermocouples 2
3
1
3
2
Displacement ∆Z
3 3 Graphite Punch 2 1 2 Graphite Die Vacuum Chamber Pressure (b)
Figure 19.2â•… (a) Summary of major challenges associated with the processing of ceramic nanocomposites and (b) schematic representation of the current flow in spark plasma sintering (SPS), one of the most widely used processing routes in the development of bulk ceramic nanocomposites.1
â•… 311
19.3 Overview of Developed Nanoceramics and Ceramic Nanocomposites
processing and characterization of these materials. In the following, those materials having grain sizes, of at least one of the phases, less than about 200â•›nm will be referred to as nanocrystalline materials.
19.3.1â•…Monolithic Nanoceramics Among various materials, WC-based cemented carbides are used in machining as well as in other damage-tolerant applications. However, the metallic binders in conventional cemented carbides limit their high-temperature applications. Therefore, the successful fabrication of inherently difficult-to-sinter WC demands attention. To this end, Omori7 and Cha et al.8 used the SPS route to consolidate binderless WC at 1900 and 1700°C, respectively. However, such high sintering temperatures induced appreciable grain growth, resulting in extremely poor mechanical properties. Eskandarany et al.9 synthesized pure WC powders (∼ 7â•›nm size) via mechanical solid-state reduction of WO3 and Mg, followed by solid-state reaction of W and C using high-energy ball milling (HEBM). These powders were consolidated at 1690°C via the PAS route. It was reported that the as-consolidated WC, having grain size of around 25â•›nm (Fig. 19.3a), could exhibit extremely high hardness of around 23â•›GPa. Field-assisted sintering of TiN to 97% densification at 1200°C resulted in the retention of 90-nm grain size in the sintered product, starting with powders of ∼70â•›nm.10 Also, dense TiN with grain size of 0.5–1â•›µm could be obtained by conventional sintering of the same powders at 1500°C. Lee et al.11 reported grain size of around 200â•›nm in 99% densified TiO2 by SPS at 700°C. Compared to this, the microwave sintering resulted in grain size of around 300â•›nm, as against 1–2â•›µm grain size in pressureless-sintered TiO2. Among oxide ceramics, 3â•›mol% yttria-doped ZrO2 can be sintered to nearly theoretical density by SPS at 1200°C.12 The nanoceramics (average grain size ∼90â•›nm) exhibited higher hardness (∼14.5â•›GPa), compared with conventional monolithic ZrO2 (∼11â•›GPa). One of the chapters of this section discusses the wear properties of ZrO2 nanoceramics. Importantly, the hardness possessed by nanocrystalline zirconia is even superior to those obtained by conventional composites of zirconia with 30 vol% harder reinforcements, such as ZrB2,13,14 TiB2,15 and WC.16 In a different set of SPS experiments at 1100°C, Nygren et al.17 consolidated nanocrystalline powders (20â•›nm) of yttria-stabilized tetragonal zirconia (Y-TZP) to near-theoretical densification, while maintaining final grain size to below 100â•›nm. In another work, pressureless sintering of Y-TZP powders (colloidal route) at a lower temperature of 1150°C, resulted in maintaining nanocrystalline grain size (∼110â•›nm).18 Nanocrystalline structure and tailoring of yttria content enable the attainment of superior mechanical properties with a maximum hardness of ∼13â•›GPa and fracture toughness of ∼14â•›MPaâ•›m1/2 for 2Y-TZP. For comparison, conventional submicron-grained (<1â•›µm) Y-TZP ceramics (hot pressed, 1450°C) exhibit a hardness of ∼11â•›GPa and a maximum fracture toughness of ∼10â•›MPa.m1/2.19 Nanocrystalline α-Al2O3 (∼150â•›nm), developed via high-pressure sintering (∼1â•›GPa, ∼1000°C), exhibited excellent hardness (25.3â•›GPa), which is higher than that of conventional monolithic alumina (∼20â•›GPa).20 A combination of HEBM and SPS was used to develop nanocrystalline β-Si3N4 ceramics (grain size 68â•›nm)21 with
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50 nm (a)
100 nm (b)
Figure 19.3â•… Some selected bright field TEM images representative of various dense monolithic nanoceramics: (a) pure WC with an average grain size of about 25â•›nm consolidated via PAS at 1690°C48 and (b) nanoscaled grain size of SiC (∼60â•›nm) after HIPing.22
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a better hardness (18â•›GPa), compared with conventional β-Si3N4 ceramics (∼16â•›GPa). Besides SPS, sinter forging also enabled the development of dense nanocrystalline SiC.22 A bright field transmission electron microscopy (TEM) image, revealing the finer scale microstructure of the developed nanoceramic, is shown in Figure 19.3b. Critical consideration of the published results, as discussed previously, implicates that development of bulk nanocrystalline ceramics, without second phase reinforcement, can be successfully achieved via advanced sintering techniques such as PAS, sinter-HIPing, and SPS. Also, judicious combinations of novel powder synthesis techniques with the sintering processes can be helpful in maintaining nanocrystalline grains in the densified nanoceramics. Also, superior mechanical properties impart better wear resistance to the nanoceramics, compared with those of conventional ceramics. However, to date no significant improvement of fracture toughness could be achieved in nanocrystalline monoliths. To couple the improvements from advanced processing techniques with fracture toughness improvement, nanocomposite design is recommended.23
19.3.2â•… Alumina-Based Nanocomposites The following is a brief survey of some of the literature reports24–35 pertaining to the sintering conditions, microstructure, mechanical properties, and toughening mechanisms of Al2O3-based nanocomposites. Zhan et al.24 developed ZrO2-toughened Al2O3 nanocomposite via SPS at 1100°C and the microstructure is characterized by alumina matrix grain size of 96â•›nm and the zirconia reinforcements of 265â•›nm. A good combination of mechanical properties, in particular, hardness of approximately 15.2â•›GPa and fracture toughness of around 8.9â•›MPaâ•›m1/2, was recorded with Al2O3–(20 vol%)ZrO2 nanocomposite. With respect to Al2O3–SiC nanocomposites,25 it was reported that SPS of Al2O3–(5 vol%)SiC powders at 1450°C resulted in superior strength of around 980â•›MPa, which is much higher than the 350â•›MPa of monolithic Al2O3.25 The property improvement corroborates well with earlier findings of Niihara, for hot pressed Al2O3–(5–10 vol%)SiC nanocomposites.25 In another work, Gao et al.26 developed Al2O3–SiC–ZrO2(3Y) nanocomposites via SPS under similar conditions and measured an even superior strength of ∼1.2â•›GPa. Despite ZrO2, incorporation, the nanocomposite exhibited a modest fracture toughness of around 4â•›MPaâ•›m1/2 (slightly better than that of monolithic alumina ∼3.5â•›MPaâ•›m1/2), while maintaining high hardness of around 19â•›GPa, similar to that of pure alumina. A novel processing route for Al2O3–SiC nanocomposite fabrication involves infiltration of porous alumina matrix by a polymer (polycarbosilane) precursor, followed by sinter-HIPing.27 Sintering at 1800°C, followed by HIPing at 1700°C, enabled achievement of neartheoretical densification of Al2O3–(x vol%)SiC nanocomposites (x ∼ 3–8). The nanocomposites had a modest toughness of ∼5â•›MPaâ•›m1/2, irrespective of the reinforcement content. An alumina-based nanocomposite, containing nanocrystalline Ti(C0.7N0.3) (30 wt%) and nano-SiC (5 wt%) second phases, was processed by hot pressing at 1650°C (30â•›MPa).28 The novel nanocomposite possesses a combination of high fracture toughness (∼8â•›MPaâ•›m1/2) and improved strength (∼800â•›MPa). In accordance with the available toughening models, the change in fracture mode and
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tortuosity in crack propagation by the nanoreinforcements have been cited as the possible reasons for the improved fracture properties. As an alternative processing approach, reactive hot pressing of a powder mixture, containing mullite, aluminum, and carbon, at 1800°C was used to develop Al2O3–(18 vol%)SiC nanocomposites. TEM analysis identified the presence of nanocrystalline SiC particles within the alumina matrix, while submicron SiC reinforcements were present at the intergranular regions (Fig. 19.4). The nanocomposite exhibited high strength (∼800â•›MPa), but moderate fracture toughness (∼3.1â•›MPaâ•›m1/2).31 The combination of HEBM and SPS (1480°C, 4 minutes) of powder mixture of Ti, graphite, and Al2O3 led to the development of Al2O3–(35 vol%)TiC nanocomposite (grain sizes were Al2O3, 400â•›nm, and TiC, 200â•›nm; see Fig. 19.4b). Importantly, Al2O3–TiC in situ nanocomposites exhibited superior fracture strength of ∼ 950â•›MPa, compared with that of hot pressed Al2O3–TiC microcomposites (∼800â•›MPa). However, the hardness (∼20â•›GPa) and fracture toughness (∼4â•›MPaâ•›m1/2) were found to be comparable, independent of processing route.32 In efforts to enhance fracture toughness of brittle alumina, the incorporation of nanoscaled piezoelectric ceramic was attempted. Zhan et al.29 sintered Al2O3– Nd2Ti2O7 nanocomposite via SPS in the temperature range of 1000–1150°C and among these nanocomposites, Al2O3–(9 vol%)Nd2Ti2O7 exhibited a superior fracture toughness of 5.7╯±â•¯0.4â•›MPaâ•›m1/2. The fracture toughness increment was attributed to the conversion of mechanical energy to electrical energy during localized deformation in the presence of piezoelectric second phase (Nd2Ti2O7). In their follow-up work, γ-Al2O3–(7.5 vol%)BaTiO3 nanocomposite was sintered using the SPS route under similar conditions.30 The grain size of the alumina matrix was maintained at 190â•›nm, with BaTiO3 particles in the nanoscale range. The fracture toughness was appreciably increased to 5.3â•›MPaâ•›m1/2 due to the domain switching in ferroelectric BaTiO3.33,34 Another Al2O3 nanocomposite, that is, Al2O3–(10 vol%)Nb, with high fracture toughness (∼8â•›MPaâ•›m1/2) was sintered via high-pressure (2â•›GPa) sintering at 900°C.35 The incorporation of a metallic phase did not degrade the excellent hardness of alumina ceramics, and the composite possessed a hardness of 20–23â•›GPa. In the same work, the preparation of Al2O3–(6 vol%)diamond nanocomposite with excellent hardness (∼32â•›GPa) was also reported. From this, it is evident that SPS has greatly assisted the development of alumina-based ultrafine grain nanocomposites. While strength and hardness are enhanced relative to conventional alumina-based ceramics, the most significant outcome has been a threefold increase in fracture toughness. Furthermore, the incorporation of a softer phase to enhance fracture toughness is possible without degrading hardness properties.
19.3.3â•… Tungsten Carbide-Based Nanocomposites Due to the attractive combination of properties such as high hardness, high elastic modulus, superior wear resistance, good thermal conductivity, and low thermal expansion coefficient, WC-based cermets are traditionally used as cutting tools, wear parts, tools, and dies. The addition of ductile metals (cobalt) has traditionally been
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19.3 Overview of Developed Nanoceramics and Ceramic Nanocomposites
SiC
SiC
100 nm (a)
TiC Al2O3
500 nm (b)
100 nm (c)
Figure 19.4â•… Microstructures of ceramic nanocomposites showing (a) TEM micrograph of alumina– silicon carbide nanocomposites consolidated via reactive hot pressing,25 (b) TEM image and selected area diffraction pattern (SADP) of Al2O3–TiC nanocomposite, developed in situ via combination of HEBM and SPS,32 and (c) SEM image of fracture surface of ZrO2–(30 vol%) ZrB2 nanocomposites developed via SPS.49
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used to enhance sinterability and fracture resistance. Conventional cemented carbides [WC–(6â•›wt%)Co], sintered via pressureless sintering, commonly exhibit a fracture toughness of around 14â•›MPaâ•›m1/2, which is significantly higher than the fracture toughness of monolithic WC (∼4â•›MPaâ•›m1/2). However, the hardness is lowered compared with monolithic WC (∼16â•›GPa vs. ∼25â•›GPa) due to the presence of softer metal binder. To obtain cemented carbides with improved combination of properties, the development of nanocrystalline cemented carbides was successfully attempted via a combination of spray conversion processing for nanoscaled powder synthesis36 and subsequent consolidation via advanced sintering techniques. The minimization of grain growth was achieved by addition of grain-growth inhibitors such as VC, TaC, while densifying via a conventional sintering technique.36 The reduction in grain sizes of the WC grains to ultra-fine/nanoscale range enhances the hardness and strength considerably. Additionally, the solid solution strengthening due to enhanced solubility of W and C in Co binder, with reduced WC grain size enhanced the mechanical properties. Sivaprahasam et al.37 developed WC–(12â•›wt%) Co cemented carbides via SPS with better hardness (∼15.2â•›GPa), as compared to conventional pressureless sintering (∼13.8â•›GPa). Similar effect of microstructural refinement has also been reported by Cha et al.38 who processed WC–(10â•›wt%)Co nanomaterial via SPS (1000°C, 10 minutes). The SPS-processed compacts possessed a better hardness of ∼18â•›GPa, while similar compositions, conventionally sintered (1100°C), had a lower hardness of ∼16â•›GPa. Michalski et al.39 reported a higher hardness of ∼22â•›GPa for pulse plasma sintered (1100°C, 5 minutes, 60â•›MPa) WC– (12â•›wt%)Co nanomaterials (WC grain size ∼50â•›nm). Similar hardness for bulk WC–Co nanomaterials has also been measured by Jia et al.40 Furthermore, Richter et al.41 observed that the hardness of nanocrystalline WC-based ceramics can reach a value close to, or even exceeding, that of WC single crystal. Despite excellent hardness, the fracture toughness of the nanocrystalline cemented carbides remains inferior. According to Richter et al.,41 the fracture toughness of ultrafine and nanocrystalline WC–Co can be as low as 5–6â•›MPaâ•›m1/2. The inferior fracture toughness of nanocrystalline cemented carbides was attributed to the reduction in mean free path of the binder phase and increase in constraint to plastic deformation ahead of the propagating crack tip.40 In fact, metallic Co phase actually behaves in a brittle manner below a critical mean free path.42 Another possible reason for lower fracture toughness is the reduction in the hcp/fcc ratio of Co phase due to increased concentration of solute (W and C).43 This minimizes the toughening, arising from the hcp-to-fcc transformation of Co binder phase. Compared with conventional cemented carbides, the fracture toughness of nanocrystalline WC–Co is not affected to a noticeable degree with increasing hardness (decreasing grain size). In fact, an interesting observation made by Jia et al.40 reveals that fracture toughness of nanocrystalline WC–Co can increase moderately with an increase in bulk hardness (Fig. 19.3). Ductile metal bridging was reported to be the dominant toughening mechanism in nanocermets.40 A study by Kim et al.44 reveals that fracture toughness increases significantly with decreasing hardness, even for SPS-processed nanocrystalline/ultra-fine WC–Co. In the same work,44 a fracture toughness of ∼ 12â•›MPaâ•›m1/2 was measured for WC–(10â•›wt%)Co (WC grain size ∼350â•›nm), which is superior to that of 6â•›MPaâ•›m1/2, as reported by Ritcher et al.41 for the same composi-
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tion with similar grain sizes (∼400â•›nm). Additionally, Michalski et al.39 processed WC–(12â•›wt%)Co nanomaterials (∼50â•›nm) with superior fracture toughness (∼15â•›MPaâ•›m1/2). It was also stated that, by increasing the binder content in nanocrystalline WC–Co, fracture toughness can be restored to nearly that of conventional cemented carbides, although at the same time not compensating with hardness owing to the grain size refinement.43 Though cemented carbides attracted wider attention as materials for cutting tools, dies, and wear-resistant parts, yet the presence of metal binder renders them unsuitable under corrosive industrial environments. In attempts to replace metallic binder with ceramic additives, Basu et al.45 reported the development of WC–ZrO2 nanocomposites. SPS of starting powders of WC (∼200â•›nm) and ZrO2 (3Y) (∼27â•›nm) at 1300°C (600â•›K/min) for 5 minutes enabled achievement of near-theoretical density of the WC–(6â•›wt%)ZrO2 nanocomposite. TEM investigation revealed the presence of 300- to 400-nm WC matrix grains, reinforced with nanocrystalline (<90â•›nm) ZrO2 particles.46 Also, the developed nanocomposites [WC–(6 wt%)ZrO2] exhibited superior hardness (∼24â•›GPa), compared with the WC–Co materials (∼16â•›GPa), and high wear resistance (wear rate ∼ 10−8â•›mm3/Nâ•›m).47 However, the fracture toughness (∼6â•›MPaâ•›m1/2) remains on the lower side, though comparable with most WC–Co nanocomposites.41 Eskandarany48 processed WC–(18â•›wt%)MgO nanocomposites, which were characterized by nanocrystalline WC grains and MgO (<50â•›nm). This nanocomposite exhibited a high fracture toughness of around 14â•›MPaâ•›m1/2, along with moderate hardness (15â•›GPa) and elastic modulus (413â•›GPa). Importantly, the presence of higher amounts of relatively softer MgO for fracture toughness enhancement does not degrade the hardness to values lower than those of conventional cermets. Thus, the developed nanocomposite, possessing a better combination of hardness and fracture toughness, can potentially replace conventional WC-based cermets in some of the existing as well as futuristic applications.
19.3.4â•… Zirconia-Based Nanocomposites Despite having high fracture toughness and excellent strength (700–1200â•›MPa), Y-TZP monoliths possess lower hardness (10–11â•›GPa) compared with other structural ceramics, such as Si3N4 or Al2O3. Therefore, various approaches, such as incorporation of hard reinforcements (such as ZrB2, TiB2) and, in particular, development of nanoceramics/nanocomposites based on TZP were attempted. The SPS route is utilized to develop ZrO2–(30â•›vol%)ZrB2 nanocomposites49 characterized by the presence of coarser tabular or elongated ZrB2 particles (∼2– 3â•›µm) and the equiaxed nano-ZrO2 matrix grains (∼100–200â•›nm). Owing to such characteristic microstructure, the composites are referred to as nano/microcrystalline composites (Fig. 19.1). The ZrO2–ZrB2 nanocomposites exhibited a superior hardness of ∼14â•›GPa, with fracture toughness varying between 4 and 10â•›MPaâ•›m1/2, depending on Y2O3 stabilization. Transformation toughening was identified as the main toughening mechanism, along with limited contribution from crack deflection by hard ZrB2 reinforcements. Yoshimura et al.50 sintered ZrO2–Al2O3 nanocomposites by SPS and the microstructure is characterized by nanosized (∼50â•›nm) matrix
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(ZrO2) grains. Jiang and co-workers51 developed ZrO2(2Y)–(5–40 vol%)WC nanocomposites via hot pressing at 1450°C and the nanocomposites with 40 vol% nanoWC exhibited an excellent combination of mechanical properties, such as strength of ∼ 2â•›GPa, hardness of ∼15â•›GPa, and fracture toughness of ∼9â•›MPaâ•›m1/2. The strength and hardness values are significantly higher than those of conventionally sintered ZrO2–WC composites, containing similar vol% of micron-sized WC.51 Hirvonen et al.52 sintered zirconia–zircon (ZrSiO4) nanocomposite, characterized by nanocrystalline (∼200â•›nm) ZrO2 grains and intergranular zircon glassy phase. In efforts to develop machinable ZrO2-based materials, Li et al.53 developed ZrO2 nanocomposites via hot pressing nanocrystalline BN-coated ZrO2 particles at 1400°C. Though the mechanical properties degraded to some extent with BN addition (up to 30 vol%), yet the properties were far superior to those of the microcomposites of corresponding compositions. For example, upon addition of 30 vol% BN, the nanocomposites exhibited a high strength of ∼800â•›MPa and fracture toughness of ∼8â•›MPaâ•›m1/2, while the microcomposites of similar composition had a moderate strength of ∼500â•›MPa and lower fracture toughness of ∼5â•›MPaâ•›m1/2.
19.4 OVERVIEW OF TRIBOLOGICAL PROPERTIES OF CERAMIC NANOCOMPOSITES It has been generally recognized that high strength, hardness, and chemical inertness, even at elevated temperatures, render ceramics potentially useful in demanding tribological applications. It has been experimentally confirmed by various research groups that reduction of microstructural scale leads to significant improvements in the wear resistance of ceramics.43,54–57 Two reasons for this grain size dependence of wear behavior were proposed. First, due to considerable improvement in hardness and yield strength, the rate of accumulation of plasticity-controlled damage during the initial deformation-controlled wear is significantly reduced. Second, the smaller flaw sizes result in a considerable increment in the plasticity-induced critical stress σD(t), a factor that controls the subsequent brittle-fracture-controlled wear. This can be explained on the basis of the following equation55:
K (t ) = (2 / π 0.5 )σ D (t ) + (q )βl 0.5
(19.1)
where l is grain size, β (≤1) is a factor relating flaw size to grain size, q is the effective tensile residual stress, and K(t) is the time-dependent stress intensity factor. In the event that K(t) becomes equal to grain boundary toughness, a fracture-controlled wear process is initiated.55 In accordance with this hypothesis, the delay in wear transition from mild to severe wear for nanostructured ceramics was experimentally observed by Chen et al.56 In the case of brittle-fracture-dominated material removal, the Evans model predicts that the lateral dimension of pullout is proportional to E3/5/ (Kc1/2H29/40) and the pullout depth is proportional to E2/5/H3/2 (Ref. 57). From this, it can be said that, owing to higher hardness of nanoceramics, the wear loss by grain pullout will be significantly lower than in the corresponding conventional ceramics. To illustrate the differences in wear properties of nanostructured ceramic composites with conventional composites, some selected results55 are presented in
â•… 319
19.4 Overview of Tribological Properties of Ceramic Nanocomposites
Figures 19.5–19.7. In particular, the results presented in Figure 19.5a,b, reveal significantly lower material damage by pullout in ultra-fine-grained (400â•›nm) αAl2O3/β-Al2TiO5 compared with the coarser grained (2.2â•›µm) counterpart.55 That nanocomposite design can lead to reduction in wear rate by reducing the dimension and rate of grain pullout, has also been reported by Merino and Todd.58 The change in fracture mode from intergranular to transgranular on reinforcing Al2O3 matrix by intergranular SiC nanoparticles as well as the suppression of twins and dislocation pile-ups by the transgranular SiC nanoparticles has been a beneficial factor behind the improved wear behavior of the Al2O3–SiC nanocomposites. Some typical worn surface features indicate lower severity of pullout and the corresponding fracture mode within the pullouts for the Al2O3–SiC nanocomposites; the evidences for such
2 µm
2 µm (a)
(b)
20 µm (c)
20 µm (d)
Figure 19.5â•… (a) SEM micrographs showing wear damage in ultra-fine-grained (UFG) (400â•›nm) α-Al2O3/β-Al2TiO5 ceramic and (b) coarse-grained (CG) α-Al2O3/β-Al2TiO5 ceramic (2.2â•›µm). One can observe the (a) severity and the large scale of the grain pullout in the CG ceramic, compared with (b) UFG ceramic.55 Selected SEM images revealing the worn surface features of (c) monolithic Al2O3 and (d) Al2O3–SiC nanocomposites. Such grain pullouts are largely minimized, indicating (d) better wear resistance of the Al2O3 nanocomposite reinforced with 10% SiC.58
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9 × 10–8
Specific wear rate (mm3/Nm)
8 × 10–8 7 × 10–8 6 × 10–8 5 × 10–8 4 × 10–8 3 × 10–8 2 × 10–8 1 × 10–8 Conventional Si3N4
Conventional Si3N4/TiN
Si3N4/TiN Nanocomposite
Figure 19.6â•… Comparison of wear rate in air for conventional Si3N4, conventional Si3N4–TiN composites and Si3N4–TiN nano/nanocomposites.61
observations is provided in Figure 19.5. An improvement in wear resistance by dispersing nanocrystalline SiC in Al2O3 has also been experimentally measured during sliding wear59 as well as erosive wear.60 A notable improvement in wear resistance of pulse electric current sintering (PECS) processed Si3N4–(30 vol%)TiN nanocomposite (grain size╯<╯20â•›nm) (Fig. 19.6) was reported by Yoshimura et al.61 In general, the cemented carbides are desired to exhibit excellent wear resistance for most of their applications, especially for use as cutting tools and die materials. However, there is ambiguity over the primary wear mechanism; it has been observed in most cases that the initiation of the wear process primarily results from the removal of the ductile binder phase, following plastic deformation and microabrasion.62–66 From this perspective, cemented carbides with finer grain size should ideally exhibit improved wear performance due to enhanced resistance to deformation of the binder phase and also the increased hardness of the finer grained composites. Also, it has been reported that the higher hardness of the nanostructured cemented carbides leads to an increase in wear resistance (see Fig. 19.7).34,36
19.5 CONCLUDING REMARKS As a concluding note, it is of interest to assess whether the approach of microstructural refinement leading to development of nanostructured ceramics or composites can lead to better wear-resistance properties. The enhancement of mechanical properties can evidently lead to an increase in transition load to severe wear; therefore, it is expected that the nanoceramic composites can be used in the mild wear regime over a broader spectrum of operating conditions than conventional micron-sized
19.5 Concluding Remarks
â•… 321
Sliding Wear Coefficient (K; ×10–7; mm3/N m)
10 9 8 7 6 5 4
Sliding Wear Coefficient (K; ×10–7; mm3/N m)
3 900
1000
1100 1200 1300 1400 Hardness (HV; kg/mm2) (a)
1500
1600
Nanostructured WC–Co cemented carbides Conventional WC–Co cemented carbides
7
6
5
4
3
2 800
1200
1600 2000 Hardness (HV; kg/mm2) (b)
2400
Figure 19.7â•… (a) Plot of wear coefficient (K) with hardness for WC–Co cemented carbides during sliding wear against steel (replotted with data from Reference 65); (b) variation of wear rate with hardness for conventional and nanostructured cemented carbides during sliding against silicon nitride (replotted with data from Reference 63).
materials. To this end, various factors, which are to be considered in analyzing the wear-resistance properties of nanostructured ceramics, are summarized in Figure 19.8. Against this backdrop, the next two chapters discuss the tribological properties of nanostructured ceramics and composites based on oxide and non-oxide ceramics.
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Phase transformation induced microcracking(e.g., ZrO2)
Enhancement of mechanical properties and less abrasive wear
Wear of nanoceramic composites
Critical load for abrasion induced brittle fracture
Tribochemical wear in case of non-oxide ceramic (WC)
Figure 19.8â•… Summary of various factors that determine the wear-resistance properties of nanoceramic composites; the influence of each factor is discussed in two case studies discussed in Chapters 20 and 21.
REFERENCES ╇ 1â•… A. Mukhopadhyay and B. Basu. Consolidation-microstructure-property relationships in bulk nanoceramics and ceramic nanocomposites: A review. Int. Mater. Rev. 52 (2007), 257–288. ╇ 2â•… B. Basu, G. B. Raju, and A. K. Suri. Processing and properties of TiB2-based materials: A review. Int. Mater. Rev. 51(6) (2006), 352–374. ╇ 3â•… V. Bounhoure, S. Lay, and M. Loubradou. Missiaen: Special WC/Co orientation relationships at basal facets of WC grains in WC-Co alloys JM. J. Mater. Sci. 43 (2008), 892–899. ╇ 4â•… K. Niihara and Y. Suzuki. Strong monolithic and composite MoSi2 materials by nanostructure design. Mater. Sci. Eng. A 261 (1999), 6–15. ╇ 5â•… M. J. Mayo, D. C. Hague, and D. J. Chen. Processing nanocrystalline ceramics for applications in superplasticity. Mater. Sci. Eng. A 166 (1993), 145–159. ╇ 6â•… J. R. Groza. ASM Materials Handbook, Vol. 7. ASM International, Materials Park, OH, 1998, 583. ╇ 7â•… M. Omori. Sintering consolidation, reaction and crystals growth by the spark plasma sintering. Mater. Sci. Eng. A 287 (2000), 183–188. ╇ 8â•… S. I. Cha and S. H. Hong. Microstructure of binderless tungsten carbides sintered by spark plasma sintering process. Mater. Sci. Eng. A 356 (2003), 381–389. ╇ 9â•… M. S. E. Eskandarany, A. A. Mahday, H. A. Ahmed, and A. H. Amer. Synthesis and characterizations of ball-milled nanocrystalline WC and nanocomposite WC-Co powders and subsequent consolidations. J. Alloy. Compd. 312 (2000), 315–325. 10â•… J. R. Groza, J. Curtis, and M. Kraemer. Field assisted sintering of nanocrystalline titanium nitride. J. Am. Ceram. Soc. 83 (2000), 1281. 11â•… Y. I. Lee, J. H. Lee, S. H. Hong, and D. Y. Kim. Preparation of nanostructured TiO2 ceramics by spark plasma sintering. Mater. Res. Bull. 38 (2003), 925–930. 12â•… B. Basu, J. H. Lee, and D. Y. Kim. Development of nanocrystalline wear resistant Y-TZP ceramics. J. Am. Ceram. Soc. 87(9) (2004), 1771–1774. 13â•… B. Basu, J. Vleugels, and O. Van Der Biest. Development of ZrO2-ZrB2 composites. J. Alloy. Compd. 334(1–2) (2002), 200–204. 14â•… A. Mukhopadhyay, B. Basu, S. Das Bakshi, and S. K. Mishra. Pressureless sintering of ZrO2-ZrB2 composites: Microstructure and properties. Int. J. Refract. Hard Mater. 25 (2007), 179–188. 15â•… B. Basu, J. Vleugels, and O. Van Der Biest. Development of ZrO2-TiB2 composites: Role of residual stress and starting powders. J. Alloy. Compd. 365(1–2) (2004), 266–270. 16â•… G. Anne, S. Put, K. Vanmeensel, D. Jiang, J. Vleugels, and O. Vander Biest. Hard tough and strong ZrO2–WC composite from nanocomposites powders. J. Eur. Ceram. Soc. 25 (2005), 55–63.
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17â•… M. Nygren and Z. Shen. On the preparation of bio-, nano- and structural ceramics and composites by spark plasma sintering. Solid State Sci. 5 (2003), 125–131. 18â•… O. Vasylkiv, Y. Sakka, and V. V. Skorokhod. Low-temperature processing and mechanical properties of zirconia and zirconia-alumina nanoceramics. J. Am. Ceram. Soc. 86(2) (2003), 299–304. 19â•… B. Basu, J. Vleugels, and O. Vander Biest. Microstructure-toughness-wear relationship of tetragonal zirconia ceramics. J. Eur. Ceram. Soc. 24 (2004), 2031–2040. 20â•… R. S. Mishra, C. E. Leshier, and A. K. Mukherjee. High pressure sintering of nanocrystalline-Al2O3. J. Am. Ceram. Soc. 79(11) (1996), 2989–2992. 21â•… X. Xu, T. Nishimura, N. Hirosaki, R. J. Xie, Y. Zhu, Y. Yamamoto, and H. Tanaka. New strategies for preparing nanosized silicon nitride ceramics. J. Am. Ceram. Soc. 88(4) (2005), 934–937. 22â•… R. Vaben and D. Stover. Processing and properties of nanophase ceramics. J. Mater. Process. Tech. 92–93 (1999), 77–84. 23â•… K. Niihara. New design concept for structural ceramic nanocomposites. J. Ceram. Soc. Jpn. 99(10) (1991), 974–982. The Centennial Memorial Issue. 24â•… G. D. Zhan, J. Kuntz, J. Wan, J. Garay, and A. K. Mukherjee. Spark-plasma sintering of silicon carbide whiskers (SiCw) reinforced nanocrystalline alumina. J. Am. Ceram. Soc. 86(1) (2003), 200–202. 25â•… L. Gao, H. Z. Wang, J. S. Hong, H. Miyamoto, K. Miyamoto, Y. Nishikawa, and S. D. D. L. Torre. Mechanical properties and microstructure of nano SiC–Al2O3 composites densified by spark plasma sintering. J. Eur. Ceram. Soc. 19 (1999), 609–613. 26â•… L. Gao, H. Z. Wang, J. S. Hong, H. Miyamoto, K. Miyamoto, Y. Nishikawa, and S. D. D. L. Torre. SiC-ZrO2(3Y)–Al2O3 nanocomposite superfast densified by spark plasma sintering. Nanostruct. Mater. 11(1) (1999), 43–49. 27â•… D. Galusek, J. Sedlacek, P. Svancarek, R. Riedel, R. Satet, and M. Hoffmann. The influence of post-sintering HIP on the microstructure, hardness, and indentation fracture toughness of polymerderived Al2O3–SiC nanocomposites. J. Eur. Ceram. Soc. 27 (2007), 1237–1245. 28â•… H. Liu, C. Huang, J. Wang, and X. Teng. Fabrication and mechanical properties of Al2O3/Ti(C0.7N0.3) nanocomposites. Mater. Res. Bull. 41 (2006), 1215–1224. 29â•… M. Sternitzke. Review. Structural ceramic nanocomposites. J. European Ceram. Soc. 17(9) (1997), 1061–1082. 30â•… G. D. Zhan, J. Kuntz, J. Wan, J. Garay, and A. K. Mukherjee. Alumina-based nanocomposites consolidated by spark plasma sintering. Scripta Mater. 47 (2002), 737–741. 31â•… G. J. Zhang, J. F. Yang, M. Ando, and T. Ohji. Reactive hot pressing of alumina-silicon carbide nanocomposites. J. Am. Ceram. Soc. 87 (2004), 299–301. 32â•… Y. Zhang, L. Wang, W. Jiang, L. Chen, and G. Bai. Microstructure and properties of Al2O3–TiC nanocomposites fabricated by spark plasma sintering from high-energy ball milled reactants. J. Eur. Ceram. Soc. 26 (2006), 3393–3397. 33â•… G. Winfield, F. Azough, and R. Freer. DiP224: Neodymium titanate (Nd2Ti2O7) ceramics. Ferroelectrics 133 (1992), 181. 34â•… G. H. Haertling. Ferroelectric ceramics: History and technology. J. Am. Ceram. Soc. 82 (1999), 797–718. 35â•… R. S. Mishra and A. K. Mukherjee. Processing of high hardness–high toughness alumina matrix nanocomposites. Mater. Sci. Eng. A 301 (2001), 97–101. 36â•… B. K. Kim, G. H. Ha, and D. W. Lee. Sintering and microstructure of nanophase WC/Co hardmetals. J. Mater. Process. Technol. 63 (1997), 317–321. 37â•… D. Sivaprahasam, S. B. Chandrasekar, and R. Sundaresan. Microstructure and mechanical properties of nanocrystalline WC–12Co consolidated by spark plasma sintering. Int. J. Refract. Metal. Hard Mater. 25(2) (2007), 144–152. 38â•… S. I. Cha, S. H. Hong, and B. K. Kim. Spark plasma sintering behavior of nanocrystalline WC–10Co cemented carbide powders. Mater. Sci. Eng. A 351 (2003), 31–38. 39â•… A. Michalski and D. Siemiaszko. Nanocrystalline cemented carbides sintered by the pulse plasma method. Int. J. Refract. Metal. Hard Mater. 25(2) (2007), 153–158. 40â•… K. Jia, T. E. Fischer, and B. Gallois. Hardness and toughness of nanostructured and conventional WC-Co composites. Nanostruct. Mater. 10(5) (1998), 875–891. 41â•… V. Richter and M. V. Ruthendorf. On hardness and toughness of ultrafine and nanocrystalline hard materials. Int. J. Refract. Metal. Hard Mater. 17 (1999), 141–152.
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42â•… G. Gille. Doctoral thesis, Akademie der Wissenschaften, 1977. 43â•… S. I. Cha, S. H. Hong, G. K. Ha, and B. K. Kim. Microstructure and mechanical properties of nanocrystalline WC-10Co cemented carbide. Scr. Mater. 44 (2001), 1535–1539. 44â•… H. C. Kim, I. J. Shon, J. K. Yoon, and J. M. Doh. Consolidation of ultra fine WC and WC–Co hard materials by pulsed current activated sintering and its mechanical properties. Int. J. Refract. Metal. Hard Mater. 25 (2007), 46–52. 45â•… B. Basu, J.-H. Lee, and D.-Y. Kim. Development of WC-ZrO2 nanocomposites by spark plasma sintering. J. Am. Ceram. Soc. 87(2) (2004), 317–319. 46â•… K. Biswas, A. Mukhopadhyay, B. Basu, and K. Chattopadhyay. Densification and microstructure development in spark plasma sintered WC-6 wt. % ZrO2 nanocomposites. J. Mat. Res. 22(6) (2007), 1491–1501. 47â•… T. Venkateswaran, D. Sarkar, and B. Basu. Tribological properties of WC-ZrO2 nanocomposites. J. Am. Ceram. Soc. 88(3) (2005), 691–697. 48â•… M. S. E. Eskandarany. Mechanical solid state mixing for synthesizing of SiCp/Al nanocomposites. J. Alloys Compounds 296 (2000), 175–182. 49â•… B. Basu, T. Venkateswaran, and D.-Y. Kim. Microstructure and properties of spark plasma sintered ZrO2-ZrB2 nanoceramic composites. J. Am. Ceram. Soc. 89(8) (2006), 2405–2412. 50â•… M. Yoshimura, M. Sando, Y. H. Choa, T. Sekino, and K. Niihara. Fabrication of dense ZrO2 based nano/nano type composites by new powder preparation method and controlled sintering processes. Key Eng. Mater. 423 (1999), 161–163. 51â•… D. Jiang, O. Van der Biest, and J. Vleugels. ZrO2-WC nanocomposites with superior properties. J. Eur. Ceram. Soc. 27 (2007), 1247–1251. 52â•… A. Hirvonen, R. Nowak, Y. Yamamoto, T. Sekino, and K. Niihara. Fabrication, structure, mechanical and thermal properties of zirconia-based ceramic nanocomposites. J. Eur. Ceram. Soc. 26(8) (2005), 1497–1505. 53â•… Y. Li, J. Zhang, G. Qiao, and Z. Jin. Fabrication and properties of machinable 3Y-ZrO2/BN nanocomposites. Mater. Sci. Eng. A 397 (2005), 35–40. 54â•… B. Basu, J.-H. Lee, and D.-Y. Kim. Development of nanocrystalline wear resistant Y-TZP ceramics. J. Am. Ceram. Soc. 87(9) (2004), 1771–1774. 55â•… X. Wang, N. P. Padture, H. Tanaka, and A. L. Ortiz. Wear-resistant ultra-fine-grained ceramics. Acta Mater. 53 (2005), 271–277. 56â•… H. J. Chen, W. N. Rainforth, and W. E. Lee. The wear behaviour of Al2O3-SiC ceramic nanocomposites. Scr. Mater. 42 (2000), 555–560. 57â•… A. G. Evans. The science of ceramic machining and surface finishing II, in National Bureau of Standards Sp. Pub. 562, B. J. Hockey and R. W. Rice (Eds.). US Govt. Printing Office, Washington, DC, 1979, 1–14. 58â•… J. L. O. Merino and R. I. Todd. Relationship between wear rate, surface pullout and microstructure during abrasive wear of alumina and alumina/SiC nanocomposites. Acta Mater. 53(12) (2005), 3345–3357. 59â•… J. Rodriguez, A. Martin, J. Y. Pastor, J. Llorca, J. F. Bartolome, and J. S. Moya. Sliding wear of nanocomposites. J. Am. Ceram. Soc. 82(8) (1999), 2252–2254. 60â•… R. W. Davidge, P. C. Twigg, and F. L. Riley. Effects of silicon carbide nanophase on the wet of alumina. J. Eur. Ceram. Soc. 16 (1996), 700–802. 61â•… M. Yoshimura, O. Komura, and A. Yamakawa. Microstructure and tribological properties of nanosized Si3N4. Scr. Mater. 44 (2001), 1517–1521. 62â•… J. Y. Sheikh-Ahmed and J. A. Bailey. The wear characteristics of some cemented tungsten carbides in machining particle board. Wear 225–229 (1999), 256–266. 63â•… K. Jia and T. E. Fischer. Sliding wear of conventional and nanostructured cemented carbide. Wear 203–204 (1997), 310–318. 64â•… Q. Yang, T. Sneda, and A. Ohmori. Effect of carbide grain size on microstructure and sliding wear behavior of HVOF-sprayed WC-12% Co coatings. Wear 254 (2003), 23–34. 65â•… J. Pirso, S. Letunovits, and M. Viljus. Friction and wear behavior of cemented carbides. Wear 257 (2004), 257–265. 66â•… H. Engqvist, G. A. Botton, S. Ederyd, M. Phaneuf, J. Fondelius, and N. Axen. Wear phenomena on WC-based face seal rings. Int. J. Refract. Met. Hard Mater. 18 (2000), 39–46.
CHAPTER
20
CASE STUDY: NANOCRYSTALLINE YTTRIASTABILIZED TETRAGONAL ZIRCONIA POLYCRYSTALLINE CERAMICS To illustrate the tribological properties of nanostructured ceramics, this chapter presents experimental results obtained with spark plasma sintered yttria-stabilizedtetragonal zirconia (Y-TZP). It will be shown that the Y-TZP nanoceramics exhibit better fretting wear resistance than the conventionally sintered Y-TZP. This will be followed by a discussion of how the yttria stabilization level or the toughness affects wear resistance of spark plasma sintered ZrO2. Fretting wear experiments against steel show a high coefficient of friction (COF) of 0.5 and a low wear rate, varying on the order of 10−7â•›mm3/Nâ•›m. Material removal from the contacting interfaces occurs mainly by grain pullout caused by intergranular fracture and delamination wear.
20.1 INTRODUCTION Since the seminal work by Niihara,1 an increase in research on the development of nanoceramics and nanoceramic composites has been triggered more recently.2–9 As regards the processing of nanoceramics, materials fabrication using an external electrical field is identified as an area of active scientific and technological interest due to specific processing advantages such as high heating rate (as high as 600– 650°C/min) and lower processing time.4 As discussed in Chapter 19, the fieldactivated sintering technique (FAST) consists of the imposition of an electrical field during sintering. Among the different FAST techniques, spark plasma sintering (SPS) is finding wide use for faster densification of ceramics and composite materials.5–7,9 In spite of having specific advantages, one of the major disadvantages
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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of the SPS process is that spark sintering can cause grain growth, if the proper processing window, in terms of heating rate, sintering temperature, and dwell time, is not precisely selected. Among the oxide ceramics, the densification of alumina via the SPS process has been investigated by many researchers.5–7 Hong et al. performed SPS experiments to develop zirconia ceramics.9 In more recent research, SPS processing has been successfully used to produce WC–(6 wt%)ZrO2 composites with significantly higher hardness (∼24â•›GPa).2 The development of nanoceramics has also triggered researchers to understand their friction and wear behavior. To assess the tribological properties of nanostructured ceramics, the sliding wear behavior of Al2O3–SiC nanocomposite containing nano-SiC (40–800â•›nm) reinforcement (5–20â•›vol%) against partially magnesiumstabilized zirconia (Mg-PSZ) has been performed.8 The wear resistance of this nanocomposite was found to be two orders of magnitude greater than that of unreinforced alumina at higher contact loads (125–150â•›N). Intergranular fracture followed by grain pullout is observed to be the governing wear mechanisms. The microstructural parameters, such as grain size, porosity, and grain pullout behavior, play a vital role in wear resistance. It can be reiterated here that, among oxide ceramics, the transformation-toughened zirconia ceramics with better toughness and relatively low elastic modulus are considered to be ideal wear-resistant materials for a variety of engineering applications.10–13 It is well known that the relative behavior of ceramic materials is influenced by normal load and that various wear processes operate under different loads. Low friction coefficient and negligible wear of zirconia ceramics were observed under mild combinations of low load and low speed.14–16 The ultra-fine-grained (180â•›nm) yttria-doped tetragonal zirconia (Y-TZP) was tested in dry N2 atmosphere by He and co-workers17 to determine the influence of porosity, varying from 1.5 to 7%, on the sliding wear properties. Stachowiak and Stachowiak18 reported that the wear mechanism for metal–ceramic sliding pairs is dominated by metallic film transfer onto the ceramic surface. For ceramic–ceramic sliding couples, plastic deformation and delamination are the dominating wear mechanisms. Evans and Marshall19 showed that the wear rate is inversely proportional to the square root of the toughness for the plasma sprayed nanostructured ZrO2 coatings, which have better wear resistance than traditional zirconia coatings. The wear rate of the nanostructured zirconia coatings is about two-fifths of that of traditional zirconia coatings under loads varying from 20 to 80â•›N. The improvement in wear resistance of nanostructured zirconia coatings is attributed to the optimization of microstructure and the enhancement of mechanical properties, which in turn increased the ability to undergo plastic deformation. Under high load (80â•›N), the wear of nanostructured zirconia coating is explained on the basis of plastic deformation and microcracking, while the interpretation for the wear of traditional zirconia coatings is in terms of deformation and brittle fracture.20 Recently, the tribological behavior of Al2O3/SiC8,21 and Si3N4/SiC22 nanocomposites has been reported in the literature. Kašiarová et al.22 investigated the tribological behavior of Si3N4/SiC nanoceramic composites against a Si3N4 ball using a pin-on-disk wear tester. The friction coefficient varied in the range of 0.4–0.6 under the experimental conditions of varying load (10, 15, and 20 N), sliding distance (600, 900, and 1200â•›m) and sliding speed (0.1, 0.2, and 0.3â•›m/s).22 The erosive wear behav-
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ior of Al2O3/SiC nanocomposites (against Al2O3) was studied by Davidge et al.21 The dispersion of secondary SiC nanoparticles in polycrystalline alumina reduced the wear rate to a significant extent and smooth transgranular fracture paths were observed in the worn nanocomposites. The tribological tests of the Al2O3/TiO2 nanocomposites against Si3N4 were carried out with the help of a ball-on-disk tribometer. The highest wear resistance was observed for 10â•›mol% of TiO2 reinforcement and the wear occurred predominantly by abrasive and plastic deformation.23 From the preceding perspective, this chapter discusses the wear properties of nanostructured tetragonal ZrO2 (t-ZrO2) ceramics. The experimental results on microstructure and properties of Y-TZP nanoceramics published elsewhere are summarized.24,27 A major emphasis is laid on understanding the material removal mechanisms. While the wear resitance properties of some nanostructured ceramics are reviewed in a recent work,25 the grain size dependent wear behavior of tetragonal zirconia ceramics are reported in a work by He et al.26
20.2 MATERIALS AND EXPERIMENTS The commercially available 3â•›mol% yttria co-precipitated ZrO2 (Tosoh grade TZ-3Y) powders were subjected to spark sintering at temperatures in the range of 1100– 1300°C under a pressure of 30â•›MPa. Bright field transmission electron microscopy (TEM) imaging (Fig. 20.1) of the ZrO2, SPS at 1200°C, confirmed the presence of finer grains of tetragonal zirconia (average size 70–80â•›nm). However, the occasional coarser t-ZrO2 grains with sizes around 110–130â•›nm were also observed. On critical assessment of the data presented in Reference 3, it is observed that the grain sizes of the pressureless sintered 3Y-TZP (starting powder TZ-3Y, Tosoh), densified at 1200°C for 2 hours with varying heating rate of 2–200°C/min, lie in the range of 240–280â•›nm. In the investigation under discussion here, the occurrence of much finer grain sizes (∼100â•›nm) in SPS-processed 3Y-TZP (1200°C, 5 minutes) can be attributed to extremely high heating rate (650°C/min) and lower holding time (5 minutes). Furthermore, the sintered grain size data clearly suggest that full densification without promoting significant grain growth is feasible in the SPS route. Although grain growth is visible in the SPS-1300 sample (not shown), the sizes of tetragonal ZrO2 grains are in the range of 200–300â•›nm. In a different set of experiments, 3Y-ZrO2 was mixed with TZ-O in various proportions to obtain powder mixtures with overall Y-content of 2â•›mol% (TM2 grade), 2.5â•›mol% (TM2.5 grade), and 2.75â•›mol% (TM 2.75 grade). The starting powders with overall yttria content in the range of 2 and 3â•›mol% are densified using SPS under identical sintering conditions of 1200°C for 5 minutes with a heating rate of ∼600°C/min. In the SPS experiments, the powders were inserted into a graphite die of 10-mm diameter. The die assembly was placed inside the SPS chamber between the two electrodes. During the SPS process, a high current of the order of 1–1.2â•›kA was passed through the graphite electrode, to carry out sintering. For comparison, the densification of 3Y-TZP powders was also accomplished by conventional hot pressing (HP) in vacuum at 1450°C for 1 hour. The heating rate in hot pressing experiments was 50°C/min. The hot pressed samples were designated as CS-1450. The average grain size of the CS-1450 sample is around 300â•›nm.
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Figure 20.1â•… Bright field TEM image of the 3Y-TZP ceramic spark sintered at 1200°C for 5 minutes.24
Figure 20.2 presents the mechanical property data, particularly hardness and indentation toughness, versus overall yttria content for Y-TZPs densified by SPS. The hardness varies between 11.75 and 14.25â•›GPa with the highest hardness recorded for SPS-processed 3Y-TZP. The lower hardness of TM2.5 and TM2 is mainly due to the presence of lesser amounts of monoclinic zirconia (m-ZrO2), possibly during cooling from SPS. The indentation toughness varies over a wide range of 4.5– 9.5â•›MPaâ•›m1/2 with the maximum value measured for TM2 ceramic. A general trend is that toughness increases with decrease in overall amount of yttria. The use of three different Y-TZP ceramic samples with varying yttria content, that is, with variation in toughness properties, makes it possible to analyze the role of toughness of nanostructured ceramics on wear damage behavior. X-ray diffraction (XRD) investigation showed the predominant presence of tetragonal zirconia in all the SPS samples. The density data, measured in water using Archimedes’ principle, show that more than 99.5% of theoretical density is attained in all SPS ceramics. For tribological experiments, dense Y-TZP nanoceramics processed via the SPS route are used as a flat material. The friction and wear properties were measured on a ball-on-flat type of fretting wear tester (mode I, linear reciprocatory motion) under gross slip conditions in ambient conditions of temperature and humidity.
Vickers Hardness (Hv) (GPa)
15.0 14.5 14.0
12 11
Indentation Toughness (SPS) Hardness (SPS)
10 9
13.5
8 7
13.0 12.5
6 5
12.0 11.5
4 3
11.0 10.5 10.0 1.5
2.0
2.5
3.0
Yttria Content (mol %)
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Indentation Toughness (KIC) (MPa·m1/2)
20.3 Tribological Properties
2 1 3.5
Figure 20.2â•… Comparison of the variation of hardness and indentation toughness with mol% of Y2O3 for Y-TZP nanoceramic synthesized by the SPS route.27
0.5
COF
0.4 0.3
T3.0 TM2.5 TM2.0
0.2 0.1 0.0
0
20,000 40,000 60,000 80,000 100,000 No. of Cycles
Figure 20.3â•… Evolution of frictional behavior of nanostructured Y-TZP ceramics against bearing steel under reciprocatory sliding conditions. The wear test conditions include a normal load of 10â•›N, relative displacement of 50â•›µm, frequency 8â•›Hz, and testing duration of 100,000 cycles.27
20.3 TRIBOLOGICAL PROPERTIES Figure 20.3 plots the evolution of friction during sliding of the TM grade samples against steel at 10 N, while keeping all other test parameters constant, as shown in Figure 20.4a (frequency 8â•›Hz, displacement 50â•›µm, duration 100,000 cycles). For all the materials, COF attains a high value of 0.3–0.5 within the running-in period of 3000–8000 cycles. Subsequently, COF either remains constant or goes through a transition before it reaches a steady-state value, which is maintained throughout the entire fretting test of 100,000 cycles. For 3Y-TZP, the friction transition occurs
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before reaching the steady state over the first 25,000 cycles. An interesting observation is made for TM2.5 ceramics: COF decreases from 0.45 to 0.35 and a steady-state COF of 0.35 is maintained for the first 60,000 cycles before going through a transition and reaching a final steady-state value of ∼0.5 at 80,000 cycles. For TM2 ceramics, a little variation in steady-state COF around 0.5 is observed under similar operating conditions. On the basis of laser profilometer measurement of wear volume, the specific wear rate is calculated. The specific wear rate is of the order of 10−6â•›mm3/Nâ•›m and the maximum wear depth is in the range of 4–16â•›µm, depending upon the toughness of the materials. No consistent correlation between wear rate/depth and toughness could be recorded with nanostructured Y-TZP. Considering the finer grain size of the materials and wear scar depth, it is obvious that extensive wear leads to material removal from the contacting interface. To evaluate the tribological performance, the wear results of SPS-processed samples are compared with those of the conventionally sintered materials in Figure 20.4b. The wear rate varies on the order of 10−6â•›mm3/Nâ•›m. From Figure 20.4b, it is clear that larger wear volume is observed with coarser grain sizes. A critical observation of the wear data clearly shows that SPS-processed samples experience less wear than the conventionally sintered material. Among the SPS-processed samples, ZrO2 ceramic, SPS at 1200°C, ceramic shows the best wear resistance. The material removal of the flat samples predominantly occurs by abrasive wear and the SPSprocessed sample, being harder than the conventionally sintered material, suffers less wear at the fretting contacts. The improved wear resistance is reported to be achieved with nanoceramic composites. For example, the wear resistance is improved with the addition of nanoreinforcements such as SiC or TiO2 to the alumina matrix.21,23
20.4 TRIBOMECHANICAL WEAR OF YTTRIASTABILIZED ZIRCONIA NANOCERAMIC WITH VARYING YTTRIA DOPANT Scanning electron microscopy (SEM) images showing the details of surface topographical features on the worn surface of various TZP nanoceramics are presented in Figures 20.5–20.7. For 3â•›mol% Y-TZP nanoceramic, the severe wear under the experimental fretting conditions leads to formation of abrasive grooves of larger depth and width varying in the range of 10–20â•›µm. A closer look at Figure 20.5 shows a significant intergranular fracture and grain pullout dominating the wear process. For 2.5Y-TZP, significant microcracking occurs along the fretting direction as well as perpendicular to fretting direction, compared with that of 3Y-TZP (see Fig. 20.6). For 2Y-TZP, the delaminated tribolayer is observed at the central region of the wear scar (Fig. 20.7). A closer look at the worn surface further shows that the wear process is dominated by the process of grain pullout as well as delamination due to propagation of larger surface cracks, having typical length of 100â•›µm. The wear debris particles are found to accumulate around the wear scar. From this observation, it should be clear that the intergranular fracture and delamination wear dissipates a larger amount of friction energy and hence results in higher COF.
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20.4 TRIBOMECHANICAL WEAR OF YTTRIA-STABILIZED ZIRCONIA NANOCERAMIC
8N
Steel
Y-TZP
150 µm, 8 Hz, 10,000 cycles (a) Average grain sizes 80 nm
150 nm
200 nm
300 nm
SPS-1250 SPS-1300 Sintering conditions
CS-1450
Wear rate (× 10–6), mm3/Nm
8
6
4
2
0 SPS-1200
(b)
Figure 20.4â•… (a) Schematic of the fretting test conditions and (b) variation of wear rates (unlubricated sliding) of the 3Y-TZP nanoceramics and conventionally sintered ceramic as a function of grain sizes. Abbreviations: SPS, spark plasma sintered; CS, conventionally sintered (hot pressed).27 The number in the sample designation indicates corresponding spark plasma sintering temperature: that is, the CS-1450 sample is hot pressed at 1450°C for 1 hour.
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10 µm
Figure 20.5â•… SEM images revealing the severe wear of 3Y-TZP nanoceramic after sliding against bearing steel. The evidence of plowing along and around abrasive grooves can be observed. The wear test conditions include a normal load of 10â•›N, relative displacement of 50â•›µm, frequency 8â•›Hz, and testing duration of 100,000 cycles.27
Figure 20.6â•… SEM images revealing the severe wear of 2.5Y-TZP nanoceramic after being worn against bearing steel. The evidence of microcracks on a wider abrasive groove can be observed. The wear test conditions include a normal load of 10â•›N, relative displacement of 50â•›µm, frequency 8â•›Hz, and testing duration of 100,000 cycles.27
It is important to consider contact stress in the context of description of tribomechanical wear. Taking Poisson’s ratio equal to 0.3 for both Y-TZP and steel, elastic modulus of the steel as 210â•›GPa, and that of Y-TZP as 210â•›GPa, the mean and maximum Hertzian contact stresses are calculated to be 674 and 1010â•›MPa, respectively, at 10-N load. The corresponding contact radius on Y-TZP surface is measured as 68.75â•›µm at 10 N. It is easy to understand that the contact spots, under stresses as high as 674â•›MPa or more, can undergo severe wear.
â•… 333
20.4 TRIBOMECHANICAL WEAR OF YTTRIA-STABILIZED ZIRCONIA NANOCERAMIC
100 µm (a)
100 µm (b)
Figure 20.7â•… SEM images revealing the overview of fretting damage (a) as well as details of the wear features (b) of 2Y-TZP nanoceramic after being worn against bearing steel. The evidence of delamination (b) and grain pullout of the severely worn region, located mostly at the central region of wear scar, can be noted. The wear test conditions include a normal load of 10â•›N, relative displacement of 50â•›µm, frequency 8â•›Hz, and testing duration of 100,000 cycles.27
In general, the intergranular fracture followed by grain pullout coupled with severe abrasion were observed to be the major wear mechanisms (Figs. 20.5–20.7). The presence of microcracks and intergranular fracture along with the occurrence of grain pullout show the tribomechanical wear, that is, microcutting, to be the major micromechanism of material removal for the Y-TZP nanoceramics. The severity of wear along with the observation of rough worn surfaces was commensurate with high COF (∼0.4–0.5). Material removal can occur by the extension of the lateral
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n ctio Fri ction e dir
Potential grains to be pulled out Grain boundary microcrack
Figure 20.8â•… Schematic illustrating the possible wear micromechanisms of Y-TZP nanoceramics.
cracks when they intersect each other and propagate to the surface. During sliding, the contact stress for the 3Y-TZP/steel tribocouple exceeds the critical stress for the t-ZrO2 to m-ZrO2 transformation and microcracks nucleate preferentially at the grain boundaries. In an earlier chapter, the results of Raman spectroscopy are provided to confirm such a phase transformation on as-worn ZrO2 surface. The microcracks nucleate because of grain sliding and strain incompatibility. Our present understanding is that the grain boundary microcracks develop in the tensile stress field ahead of the slider at the fretting contact (see Fig. 20.8). During repeated fretting strokes, these microcracks move either through the grain boundary or the interior of grains, depending on the relative toughness of grain boundary with respect to grain interior. Observation of the detailed worn surface morphology shows that the microcracks propagate along the grain boundary, leading to grain pullout, and they coalesce together to form a longer crack. Hence, delamination wear has a large contribution toward severe wear behavior. This process leads to the formation of coarser wear debris as observed in the case under discussion here. From the preceding discussion, the tribological properties of nanostructured ZrO2 clearly illustrate two contrasting features: (1) higher hardness corroborates well with lower wear rate measured with SPS-processed ceramics and (2) coarser grain size in conventionaly sintered ZrO2 results in large wear rate (see Fig. 20.4). Since intergranular fracture and grain pullout are dominant wear mechanisms in nanostructured ZrO2, it is quite probable that SPS-processed ZrO2 could have less wear resistance than conventionally sintered materials. However, that is not the case. This aspect, however, corroborates well with the experimental observations that the wear damage in nanoceramics is restricted to the central region of the wear pit in the case of high-toughness ZrO2 (TM2) and that such localized wear, although driven by intergranular fracture, eventually results in less measurable wear compared with conventionally sintered ZrO2.
20.6 Concluding Remarks
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20.5 COMPARISON WITH OTHER STABILIZED ZIRCONIA CERAMICS A wide range of wear rates for zirconia ceramics has been reported in the available literature. Very low wear rates in the range of 10−7–10−8â•›m3/Nâ•›m were found for selfmated partially stabilized zirconia (PSZ) under dry conditions with very low speeds.11 A polished PSZ–PSZ couple suffered very low wear loss at room temperature.13 It has been observed that the wear behavior of a particular tribological pair depends on the nature and properties of the tested materials, the environment (humidity), and the test conditions (mild or severe wear regime). He et al.26 demonstrated that the overall wear resistance of Y-TZP, tested on a pin-on-disk tester against SiC, under N2 atmosphere at room temperature, increased with the decreasing grain size of Y-TZP. A Hall–Petch type of relation is observed in TZP (grain sizeâ•›≤â•›0.7â•›µm), where the main wear mechanisms were plastic deformation and microcracking. For relatively coarser Y-TZP (grain sizeâ•›≥â•›0.9â•›µm), the wear resistance was found to be proportional to the inverse of grain diameter and the wear is mainly dominated by delamination and grain pullout. Marshall et al.14 showed that coarser grained alumina ceramics exhibited a low grinding resistance. Nanocrystalline ceramics with little or no porosity may be expected to have high wear resistance. Fischer et al.15 showed that the wear resistance of yttria-doped zirconia, measured in air at room temperature under slow sliding speed (1â•›mm/s and 9.8-N load), increased by a factor of 1200 from brittle (6Y-PSZ, KIC ∼2.5â•›MPaâ•›m1/2) to the toughest (3Y-TZP, KIC ∼11.6â•›MPaâ•›m1/2) materials. Also, the wear rate was found to be proportional to the fourth power of toughness. The wear behavior of TZPs is related to the phase transformation at the frictional interface.16 It was found that chemical effects played a dominant role in determining the wear rate, which could change over four orders of magnitude depending on changes in environment and properties of ceramics. Comparison of the wear of zirconia ceramics with that of Si3N4 (K ∼4â•›MPaâ•›m1/2), shows that, for the same fracture toughness value, ZrO2 (hardness ∼11â•›GPa, E ∼250â•›GPa) would exhibit 10 times lower wear rate than that of silicon nitride (hardness ∼30â•›GPa, E ∼255â•›GPa).15 The higher hardness and stiffness of Si3N4 than ZrO2, implies higher contact stresses of Si3N4 and, consequently, the higher tendency for fracture and mechanical wear is higher. This comparison suggests that ceramics with greater hardness and E-modulus could show inferior wear resistance.
20.6 CONCLUDING REMARKS Better wear resistance is exhibited by SPS-processed TZPs, in comparison with the conventionally sintered Y-TZP. The wear tests conducted on the developed nanoceramics of varying yttria stabilization level show almost similar frictional behavior with steady-state COF of ∼0.5. The wear rate lies around 10−6â•›mm3/Nâ•›m. Detailed topographical observation shows that tribomechanical wear, that is, material loss, occurs via intergranular fracture and grain pullout occurs in TZP nanoceramics. Microcracking around transforming tetragonal grain boundaries contributes to wear loss of nanoceramics. Deeper and wider abrasive grooves are observed on worn
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3Y-TZP nanoceramic; on the other hand, the wear damage is highly localized at the central region of worn surface on the high-toughness TZP ceramic with overall yttria content of 2â•›mol%.
REFERENCES ╇ 1â•… K. Niihara. New design concept of structural ceramics—Ceramic nanocomposites. J. Jpn. Ceram. Soc. 99(10) (1991), 974–982. ╇ 2â•… B. Basu, J.-H. Lee, and D.-Y. Kim. Development of WC-ZrO2 nanocomposites by spark plasma sintering. J. Am. Ceram. Soc. 87(2) (2004), 317–319. ╇ 3â•… D. J. Chen and M. J. Mayo. Rapid rate sintering of nanocrystalline ZrO2-3 mol% Y2O3. J. Am. Ceram. Soc. 79(4) (1996), 906–912. ╇ 4â•… C. Suryanarayana. Nanocrystalline materials. Int. Mater. Rev. 40 (1995), 41–64. ╇ 5â•… Z. Shen, M. Johnsson, Z. Zhao, and M. Nygren. Spark plasma sintering of alumina. J. Am. Ceram. Soc. 85(8) (2002), 1921. ╇ 6â•… S. W. Wang, L. D. Chen, and T. Hirai. Densification of Al2O3 powder using spark plasma sintering. J. Mater. Res. 15(4) (2000), 982. ╇ 7â•… L. Gao, J. S. Hong, H. Miyamoto, and S. D. D. L. Torre. Bending strength and microstructure of Al2O3 ceramics densified by spark plasma sintering. J. Eur. Ceram. Soc. 20 (2000), 2149. ╇ 8â•… J. Rodriguez, A. Martin, J. Y. Pastor, J. Llorca, J. F. Bartolome, and J. S. Moya. Sliding wear of alumina/silicon carbides nanocomposites. J. Am. Ceram. Soc. 82(8) (1999), 2252–2254. ╇ 9â•… J. Hong, L. Gao, S. D. D. L. Torre, H. Miyamoto, and K. Miyamoto. Spark plasma sintering and mechanical properties of ZrO2 (Y2O3)–Al2O3 composites. Mater. Lett. 43 (2000), 27. 10â•… E. P. Butler. Transformation-toughened zirconia ceramics. Mater. Sci. Technol. 1 (1985), 417–432. 11â•… R. H. J. Hannink, M. J. Murray, and H. G. Scott. Friction and wear of partially stabilized zirconia— Basic science and practical application. Wear 100 (1984), 355–366. 12â•… N. Gane and R. Beardsley. Measurement of the friction and wear coefficients of PSZ and other engineering materials using pin-on-disc machine. Proc. Int. Tribology Conf., Melbourne, 1987, Inst. Eng. Aust., Nat. Conf. Publ. 87/18, 1987, 187–192. 13â•… C. S. Yust and F. J. Carigan. Observations on the sliding wear of ceramics. ASLE Trans. 28 (1985), 245–252. 14â•… D. B. Marshall, B. R. Lawn, and R. F. Cook. Microstructural effects on grinding of alumina and glass-ceramics. J. Am. Ceram. Soc. 70(6) (1987), C-139–C-140. 15â•… T. E. Fischer, M. P. Anderson, and S. Jahanmir. Influence of fracture toughness on the wear resistance of yttria-doped zirconia oxide. J. Am. Ceram. Soc. 72(2) (1989), 252–257. 16â•… B. Basu, R. G. Vitchev, J. Vleugels, J. P. Celis, and O. Van Der Biest. Influence of humidity on the fretting wear of self-mated tetragonal zirconia ceramics. Acta Mater. 48 (2000), 2461–2471. 17â•… Y. He, L. Winnubst, A. J. Burggraaf, H. Verweij, P. G. Th. Vander Varst, and W. Bert. Influence of porosity on friction and wear of tetragonal zirconia polycrystal. J. Am. Ceram, Soc. 80(2) (1997), 377–380. 18â•… G. W. Stachowiak and G. B. Stachowiak. Unlubricated friction and wear behavior of toughened zirconia ceramics. Wear 132 (1989), 151–171. 19â•… A. G. Evans and D. B. Marshall. Wear mechanisms in ceramics, in Fundamentals of Friction and Wear of Materials, D. A. Rigney (Ed.). American Society for Metals, Metals Park, OH, 1980, 439–452. 20â•… H. Chen, Y. Zhang, and C. Ding. Tribological properties of nanostructured zirconia coatings deposited by plasma spraying. Wear 253 (2002), 885–893. 21â•… R. W. Davidge, P. C. Twigg, and F. L. Riley. Effects of silicon carbide nano-phase on the wet erosive wear of polycrystalline alumina. J. Eur. Ceram. Soc. 16 (1996), 700–802. 22â•… M. Kašiarová, E. Rudnayová, J. Dusza, M. Hnatko, P. Šajgalík, A. Merstallinger, and L. Kuzsella. Some tribological properties of a carbon-derived Si3N4/SiC nanocomposite. J Eur. Ceram. Soc. 24 (2004), 3431–3435.
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23â•… S. W. Lee, C. Morillo, J. L. Olivares, S. H. Kim, T. Sekino, K. Niihara, and B. J. Hockey. Tribological and microstructural analysis of Al2O3/TiO2 nanocomposites to use in the femoral head of hip replacement. Wear 255 (2003), 1040–1044. 24â•… B. Basu, J.-H. Lee, and D.-Y. Kim. Development of nanocrystalline wear resistant Y-TZP ceramics. J. Am. Ceram. Soc. 87(9) (2004), 1771–1774. 25â•… A. Mukhopadhyay and B. Basu. Consolidation-microstructure-property relationships in bulk nanoceramics and ceramic nanocomposites: A review. Int. Mater. Rev. 52(5) (2007), 257–288. 26â•… Y. He, L. Winnubst, A. J. Burggraaf, H. Verweij, P. G. Th. Van der Varst, and B. de With. Grain-size dependance of sliding wear in tetragonal zirconia polycrystals. J. Am. Ceram. Soc. 79(12) (1996), 3090–3096. 27â•… A. Mukhopadhyay, T. Venkateswaran, R. Verma, and B. Basu. Processing-structure-property correlation of bulk ceramic nanomaterials. Trans. Indian Inst. Met. 59(3) (2006), 303–319.
CH A P T E R
21
CASE STUDY: NANOSTRUCTURED TUNGSTEN CARBIDE– ZIRCONIA NANOCOMPOSITES In this chapter, the wear behavior of tungsten carbide–zirconia (WC–ZrO2) nanocomposite is reported. A clear transition in friction and wear behavior with load is observed under selected testing conditions. The extremely low wear rate (10−8â•›mm3/Nâ•›m) coupled with low wear depth (<1â•›µm) shows the high wear resistance of nanoceramic composites. The observed tribological properties of the nanocomposites are discussed in terms of material properties, abrasion, and tribochemical wear.
21.1 INTRODUCTION In the last few decades, nanoceramic composites received increasing attention of researchers in the materials community. Nanocrystalline materials or nanoceramic composites are defined as materials in which at least one of their phases has dimension in the range of nanometers.1–7 Although the consolidation of nanoceramic composites without excessive grain growth is achievable by spark plasma sintering (SPS), other sintering routes are also reported.4,5,8,9 A major advantage of the SPS process is that the sintering temperature is typically around 200°C below that of conventional sintering.10,11 Rapid rates of densification and heating, reduced sintering time, and retention of finer grained microstructure are the other significant characteristics of SPS.7–10,12,13 The nanograined second phase particles dispersed within the matrix or along the grain boundary of the micron- and/or submicron-sized matrix material lead to significant improvement in strength and fracture toughness by the order of two to four times greater than conventional composite materials.1,2 Following the early work of Niihara et al.,1–3 on the structural nanoceramic composites, investigations have been carried out on alumina- and silica-based nanoceramic composites. Mukherjee and coworkers have developed alumina nanocomposites with reinforcements of carbon nanotubes and measured three times higher toughness (9.7â•›MPaâ•›m1/2) Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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â•… 339
compared with that of conventional alumina.8 Also, some of our work on the development of yttria-stabilized zirconia polycrystalline (Y-TZP) nanoceramics with high hardness (∼14â•›GPa) and improvement in wear resistance is reported.10 Better mechanical properties of nanoceramic composites make them potential candidates for various structural and wear-resistance applications.4 Reports in the literature also indicate the potential of ceramic nanocomposites in machining and tribological applications. For example, a comparison of the Si3N4/ BN nanocomposites with the microcomposite shows higher fracture strength and better machinability in the former, as revealed by superior surface finish of the machined component.13 Also, nanocrystalline WC–Co composites are reported to have four times higher wear resistance, and more than twice the lifetime in cutting applications compared with conventional coarse-grained composites.7 This is due to superior mechanical properties (hardness, toughness). From Chapters 19 and 20, it is evident that considerable work has been carried out to develop nanoceramics and their composites,14,15 as well as to evaluate their potential in engineering applications.16,17 However, research work to understand the tribological properties of the nanocomposites are rather limited.18–20 Kasiarova et al.18 studied the tribological properties of Si3N4/SiC nanoceramic composites against Si3N4 using a pin-on-disk wear tester. The measured coefficient of friction (COF) varied in the range of 0.4–0.6 under varying load (10, 15, and 20â•›N), sliding distance (600, 900, and 1200â•›m), and sliding speed (0.1, 0.2, and 0.3â•›m/s).18 The erosive wear behavior of Al2O3/SiC nanocomposites (against Al2O3) was studied by Davidge et al.19 The dispersion of secondary SiC nanoparticles dispersed in polycrystalline alumina significantly reduced the wear rate, and smooth transgranular fracture paths were observed in the worn nanocomposites. Rodriguez et al.20 studied the sliding wear properties of Al2O3/SiC nanocomposites with variation in the grain size of SiC. Intergranular fracture followed by grain pullout is observed to be the dominant wear mechanism. The tribological behavior of the Al2O3/TiO2 nanocomposites against Si3N4 was performed with a ball-on-disk tribometer. The highest wear resistance was observed for 10â•›mol % of TiO2 reinforcement in which the governing wear mechanisms were abrasive and plastic deformation.21 In this chapter, some of the important results illustrating the tribological properties of a newly developed WC–(6â•›wtâ•›%)ZrO2 nanocomposite is discussed and more details can be found elsewhere.22
21.2 MATERIALS AND EXPERIMENTS In this case study, SPS-processed WC–(6â•›wtâ•›%)ZrO2 nanocomposite is used as the flat material. A commercial-grade bearing (SAE 52100) steel ball, 8â•›mm in diameter with mirror finish (surface roughness 0.02â•›µm, according to supplier) was used as the material for the counterbody. The commercially available coprecipitated ZrO2 (Tosoh grade TZ-3Y, starting particle size 27â•›nm) and commercial WC powder (99.5% pure, particle size around 0.2â•›µm, H.C. Starck, Germany) were used as starting powders. The SPS experiments were performed under a vacuum of 50–60â•›mTorr at 1300°C for 5 minutes under a pressure of 30â•›MPa, resulting in a final density of 98–99% ρth (theoretical density). The nanocomposite shows ultrahigh hardness
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(∼23â•›GPa) and moderate toughness (∼5â•›MPa m1/2). X-ray diffraction (XRD) investigation of the WC–(6â•›wtâ•›%)ZrO2 nanoceramic composites showed tetragonal ZrO2 and WC as the predominant phases. The microstructure and mechanical properties of WC–ZrO2-based materials are published elsewhere.23–25A representative lower magnification transmission electron microscopy (TEM) image (Fig. 21.1a) of the sample spark plasma sintered for 5 minutes is shown in Figure 21.1a. An overall fine microstructure with WC grains varying in the range of 0.3–0.5â•›µm along with ZrO2 particles is observed in this figure. The ZrO2 particles are indicated by black arrows. Almost all the observable WC grains have equiaxed morphology, and no abnormal grain growth of WC is observed in the spark plasma-sintered microstructure. A higher magnification TEM image is presented in Figure 21.1b to show the distribution of ZrO2 particles in the WC matrix. These nanosized ZrO2 particles are characterized by the faceted morphology. In contrast, much finer ZrO2 particles (∼30â•›nm) are observed inside WC grains, as shown in Figure 21.1b. Considering the average grain sizes of the starting powders—27â•›nm for ZrO2 and 0.2â•›µm for WC—these observations certainly indicate that limited grain growth occurred during SPS and, hence, led to the development of a nanocomposite microstructure. Motivated by the work that nanocrystalline materials can exhibit good frictional properties,26 a ball-on-flat reciprocating sliding-type wear tester was used to study the tribological behavior of the nanoceramic composite in a mode I fretting configuration.22 The WC–(6â•›wtâ•›%)ZrO2 nanoceramic composites were used as a flat sample and held mechanically on a translation table oscillating with the preset displacement and frequency. The wear experiments were conducted over constant frequency (8â•›Hz) and constant displacement stroke (50â•›µm) by having load and cycles as the external variables of the tribosystem. The load and cycles were varied from 2 to 10â•›N and from 10,000 to 100,000 cycles, respectively (Fig. 21.2). The average sliding velocity was around 0.008â•›m/s. Under the selected experimental conditions, the contact displacement in fretting mode I establishes gross slip sliding between mating surfaces across the whole contact area. It can be mentioned that mode I, among three modes of fretting wear, is widely investigated,27,28 while the research results obtained with mode II and mode III are rather limited. For a better understanding of the wear mechanisms, worn surfaces were investigated using a scanning electron microscope (SEM; JEOL-JSM840) and an electron probe microanalyzer (EPMA; JEOL-JXA8600). The depth of the wear scars on the ultrasonically cleaned worn surfaces was measured using a stylus surface profilometer (Tencor α-Step-100™). From the transverse wear scar diameter, the wear volume of the flat sample was calculated using Klaffke’s formula.29 Justifying that the normalized diameter is greater than the order of 2, that is, the wear scar diameter is greater than twice the Hertzian contact diameter, our calculation provides less than 8% error as was the case in our experiments.30
21.3 FRICTION AND WEAR CHARACTERISTICS The influence of variation in load (2, 5, and 10â•›N) on the frictional behavior of WC–(6â•›wtâ•›%)ZrO2 nanoceramic composites fretted against steel for 100,000 cycles
21.3 Friction and Wear Characteristics
â•… 341
WC
Z
200 nm (a)
WC
z WC WC
z
WC 50 nm (b)
Figure 21.1â•… Representative bright field TEM image of WC–ZrO2 nanocomposite (spark plasma sintered for 5 minutes at 1300°C) showing distribution of nanocrystalline ZrO2 particles (indicated by arrows) in submicron WC matrix (a). ZrO2 particles are distributed in the WC matrix at the triple junction and interior of a WC grain (b). Z stands for ZrO2.24
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Load (2 N, 5 N, 10 N)
Steel
WC-6 wt % ZrO2
50 µm, 8 Hz, 10,000–100,000 cycles
Figure 21.2â•… Schematic representation of contact and operation parameters.22
0.6 10 N 0.5
COF
0.4 5N
0.3 0.2 0.1
2N
0.0 0
20,000
40,000
60,000
80,000
100,000
CYCLES
Figure 21.3â•… The evolution of frictional behavior of WC–(6â•›wtâ•›%)ZrO2/steel contact during 100,000 sliding cycles with varying load at a constant frequency of 8â•›Hz and a constant displacement stroke of 50â•›µm.22
is illustrated in Figure 21.3. From Figure 21.3, it is clear that the evolution of frictional behavior strongly depends on normal load as well as fretting cycles. Under the minimum load of 2â•›N, the COF increases from a low value to 0.1 in the initial 500 cycles and steady-state COF of 0.1 is maintained throughout the entire test period of 100,000 cycles. A clear transition in the frictional behavior is noted with increase in load from 2 to 5â•›N. In the case of 5-N load, the steady-state COF of ∼0.1–0.15 is maintained up to 20,000 cycles followed by a slow increase in COF
21.3 Friction and Wear Characteristics
1.4
10,000 Cycles 50,000 Cycles 100,000 Cycles
1.2 Wear Rate (10–8 mm3/N m)
â•… 343
1.0 0.8 0.6 0.4 0.2 1
2
3
4
5
6 7 Load (N) (a)
8
9
10
11
10,000 Cycles 50,000 Cycles 100,000 Cycles
0.6
Wear Depth (µm)
0.5
0.4
0.3
0.2
0.1
1
2
3
4
5
6 7 Load (N) (b)
8
9
10
11
Figure 21.4â•… Wear rates (a) and the maximum depth of the wear scar (b) on the WC–(6â•›wtâ•›%)ZrO2 nanocomposite (flat). The maximum depth of wear scar is measured using stylus profilometer. The different testing conditions are described. Typically, 10% deviation around the average wear data was observed in our experiments.22
from 0.15 to 0.50 between 40,000 and 55,000 cycles of sliding, and a steady-state COF of 0.5 is maintained beyond 60,000 cycles. At the maximum load of 10 N, the transition to higher COF of 0.5 takes place at an early stage of around 20,000 cycles. Beyond 40,000 cycles COF is maintained at steady state (0.5) throughout the test. The wear rate with increasing testing time (10,000–100,000 fretting cycles) is plotted against normal load in Figure 21.4a. The wear rate is quite low and on the
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10 N,100K
530 nm
10 N,50K
250 nm
10 N,10K
171 nm
5 N,100K
241 nm
5 N,50K
221 nm
5 N,10K
202 nm
2 N,100K
176 nm
500 nm
Figure 21.5â•… Surface profile characteristics of worn pit on WC–ZrO2 composite, as traced by stylus profilometer for different contact conditions: load, 2–10â•›N; test duration, 10–100,000 cycles. Right side marker indicates the vertical scale.22
order of ∼10−8â•›mm3/Nâ•›m. It may be mentioned here that dense WC–(6â•›wtâ•›%)ZrO2 composite, pressureless sintered at 1600°C, shows wear rate on the order of ∼10−6â•›mm3/Nâ•›m.31 After testing for 10,000 cycles, low wear rate is measured on the worn flats that does not vary much with normal load. The wear rate measured after 50,000 cycles increases with increasing load from 2 to 5â•›N. However, no further increase in wear rate is noted at the highest load of 10â•›N. The wear rate measured after 100,000 sliding cycles shows a monotonic increase with normal load varying between 2 and 10â•›N. The maximum depth of the wear scar was measured on the ultrasonically cleaned worn surfaces using a stylus profilometer. The representative profilometer traces of worn surfaces for different contact conditions are shown in Figure 21.5. The wear depth data is plotted against testing variable in Figure 21.4b. At the lowest load of 2 N, the depth of the wear scar is not measurable after 10,000 and 50,000 cycles. The wear depth is measured as 0.15â•›µm after sliding for 100,000 cycles at 2-N load. The increase in wear depth from ∼0.20 to 0.24â•›µm with increase in the number of cycles (10,000–100,000 cycles) at an intermediate load of 5â•›N is noted. However, this increase is significant at the highest load of 10â•›N. A maximum depth of ∼0.55â•›µm was measured after 100,000 cycles at 10-N load. Summarizing, the wear data clearly show that the severity of the wear of the investigated ceramic nanocomposites, that is, the amount of material removed at the fretting contacts, increases with increase in load and test duration.
21.4 Wear Mechanisms
â•… 345
21.4 WEAR MECHANISMS The surface topography and details of the wear scar on the nanocomposite after testing under load 2â•›N for 100,000 cycles are shown in Figure 21.6. The smooth appearance of the wear scar corroborates well with the measured wear depth (Fig. 21.6a). The observation of mild abrasive scratches also indicates mild wear at low load (Fig. 21.6b). It may be noted here that the topographical features correlate well with the profilometer trace of the wear scar.
100 µm (a)
10 µm (b)
Figure 21.6â•… SEM images showing the overall topography (a) and the evidence of mild abrasive scratches (b) on the worn surface of the WC–ZrO2 nanoceramics after the test at load of 2â•›N, frequency of 8â•›Hz, and duration of 100,000 cycles. The double-pointed arrow indicates the sliding direction.22
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5 N, 100K
10 N, 100K
10 µm
10 µm
(a)
(b)
Figure 21.7â•… SEM images revealing the details of the worn surfaces on the WC–(6â•›wtâ•›%) ZrO2 nanoceramic composites. The contact conditions are described in the corresponding microstructure. Double-pointed arrow indicates sliding direction.22 2 N, 100K
10 N, 100K
10 µm (a)
10 µm (b)
Figure 21.8â•… SEM images showing the details of the wear debris generated during experiments with the nanoceramic composites. The operating parameters are mentioned in the corresponding microstructure.22
The formation of a tribochemical layer is only observed under normal load of 5â•›N and above. Observation of the difference in contrast between the tribolayer and the underlying worn surface of the nanocomposite makes it clear that the tribochemical layer (dark contrast) is rich with transferred iron oxide (see Fig. 21.7a). The smeared transfer layer (white contrast) WO3 is heavier than WC (gray contrast). It is quite probable that WC undergoes oxidation at the fretting contacts, as was also observed in our earlier work on the wear of TiCN–(WC)–Ni cermets (under similar fretting conditions).32 The formation of WO3, resulting from the tribo-oxidation of WC, was confirmed by Raman spectroscopy (not shown). However, the severe cracking of the tribolayer is indicative of its nonprotective nature and justifies the higher wear loss. Severe cracking followed by spalling of the tribolayer is much more pronounced at the higher load of 10 N after 100,000 cycles. The morphology of the wear debris after experiments at low (2â•›N) and high (10â•›N) load is shown in Figure 21.8. A striking difference in wear debris particles is
21.5 Explanation of High Wear Resistance of Ceramic Nanocomposites
â•… 347
distinctly observed. After subjecting contacts to 2-N load, the debris particles are mostly submicron in size and spherical in nature. Occasionally, it is observed that wear debris particles form agglomerates of size up to 2â•›µm. At the highest load of 10â•›N, a large fraction of sheetlike and platelet-like wear debris particles are formed. The size of the wear debris is typically around 50–100â•›µm. Submicron-sized debris particles are also observed in smaller fraction.
21.5 EXPLANATION OF HIGH WEAR RESISTANCE OF CERAMIC NANOCOMPOSITES One of the important aspects of tribology research is to establish the wear mechanisms. In the case of ceramic nanocomposites, two important issues need to be addressed: (1) whether the improved hardness results in any appreciable improvement in wear resistance and (2) what role will be played by a finer microstructure in the process of material removal from the contacting interfaces? Based on our experimental results, it is evident that higher hardness (23â•›GPa) leading to lower COF and wear is only realized at low load of 2â•›N. Additionally, the measured lower wear depth (<1â•›µm) can be correlated with higher hardness (23â•›GPa) of the nanocomposite. It also indicates that the material damage induced by wear of the nanocomposite is limited to a submicron region below the tribological surface. As the normal load increases from 2 to 5â•›N, a transition in friction and wear of the WC–(6â•›wtâ•›%)ZrO2 nanocomposite occurs. Mild wear coupled with low friction is observed at the lowest load of 2â•›N while high COF coupled with severe wear is observed at the intermediate load of 5â•›N. More severe wear is observed at the highest load of 10â•›N. At higher load (5â•›N or more), the predominant mechanism for material removal is tribochemical wear. The observation of the smooth worn surfaces and extremely fine wear debris (brighter contrast), makes it evident that WC grains are probably pulled out from the tribological interface following the intergranular fracture caused by repeated fretting strokes. Following the pullout of WC grains from the matrix, WO3 is formed by the oxidation of WC grains. Subsequently the debris particles are ejected from the worn surface. Hence, mild oxidative wear is proposed to be the major wear mechanism. This physicochemical process, involved in wear at low load (2â•›N), is highly probable and supported by the fact that the size of the wear debris particles is similar to that of WC grains. It may be noted here that a similar wear mechanism, that is, intergranular fracture followed by grain pullout, is also reported for Al2O3/ SiC nanoceramic composite20 and nanocrystalline Y-TZP.10 At higher load (5â•›N or more), the wear of the nanocomposite depends on the formation and behavior of the tribochemical layer at the contacting interface. The oxidized wear debris agglomerates at higher load and forms the tribochemical layer, which also contains iron oxide transferred from the worn steel ball. After the tribochemical layer spreads over the worn surface, the contact is established between the worn steel ball (iron oxide layer) and the tribolayer on the flat. This provides an explanation of the transition in friction leading to higher COF of 0.5, maintained in the high load regime (5â•›N or more). It may be noted here that the transition in the
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degree of wear is not observed in Cu–TiB2 nanocomposites, when tested against steel on a pin-on-disk tribometer with varying loads (20–140â•›N). The dominant mechanism of material removal is plastic deformation with formation of flakelike wear debris.33 The transition from mild to severe wear mechanism can be discussed34 on the basis of debris analysis. First the cause for submicron debris formation at low load is discussed. Based on the concept of indentation fracture mechanics, Roberts observed the minimum load required for fracture due to abrasion35: 3
P* =
54.47β K IC K IC, πη2 θ H
(21.1)
where P* is the minimum load required to produce fracture from a point contact (in newtons), β is the constant relating hardness to diagonal (2.16 for Vickers indentation), η is a constant, θ is the geometrical constant (≈0.2), KIC is the fracture toughness of the material indented (MPa m1/2), and H is the hardness of the material indented (GPa). Incorporating the material properties in Equation 21.1, it has been found that a minimum load of ∼2.15â•›N is required for fracture induced by abrasion of the investigated nanocomposites. Hence, it is quite plausible that formation of the finer wear debris (≤2â•›µm) due to intergranular fracture of the nanocomposites occurs even at such a low load (2â•›N). After their formation, the finer debris particles can get trapped between the mating counterfaces and some debris particles can eventually be ejected out to the edges of the wear scars. This indicates that the low COF of ∼0.15 and lower wear rate at 2-N load are primarily due to three-body abrasion, that is, due to rolling of the finer debris particles between the mating surfaces. However, coarser wear debris, in the form of sheets, is generally observed at higher load (≥5â•›N). Based on the adhesional theory of friction, Rabinowicz36 proposed that the average diameter of a wear debris particle is dependent on work of adhesion:
d = 60,000
Wad , H
(21.2)
where Wad is the work of adhesion, H is the hardness of the material, and d is the debris diameter. From Equation 21.2, it is clear that the size of wear debris particles, being directly proportional to the work of adhesion, should depend on normal load and tangential force. In fact, COF undergoes linear increase with Wad.35,36 This can be explained on the basis that, in order to maintain relative motion at frictional contact, higher frictional force is necessary to break the adherent bonds between interlocking asperities at higher loads. In our experiments, it is observed that COF increases with load and therefore work of adhesion is expected to be higher at the frictional contact. From the preceding discussion, it should now be evident that formation of coarser debris should be expected at higher load (≥5â•›N). Also, higher COF at higher load (≥5â•›N) implies the dissipation of greater frictional energy. It is quite probable that more energy is spent in the formation and removal of coarser wear debris from the
REFERENCES
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tribosurface. Another probable reason is the nonprotective nature of the tribochemical layer (see Fig. 21.7a), which gets fractured due to repeated small-stroke reciprocating sliding. This results in the formation of sheetlike wear debris, which is subsequently ejected around the wear scar (see Fig. 21.8b). Once the tribolayer is removed, a similar process gets repeated during the fretting tests. Since the size of the wear debris is also suggestive of the wear rate, the observation of coarser wear debris indicates higher wear rate of nanocomposites.
21.6 CONCLUDING REMARKS The WC–ZrO2 nanoceramic composites, processed by means of SPS, show lower COF (0.1) at lower load (2â•›N). Frictional behavior undergoes a clear transition leading to high COF (0.5) at higher load (5â•›N). This interesting observation also elucidates a transition from mild to severe wear at higher loads. Under the selected experimental conditions, the newly developed WC–ZrO2 nanoceramic composite exhibits low wear depth (maximum ∼0.6â•›µm) and high wear resistance (wear rate ∼10−8â•›mm3/Nâ•›m). This is due to the high hardness of the nanocomposite (∼23â•›GPa). The formation of mild abrasive scratches along with finer (submicron-sized) wear debris particles indicates better wear resistance of the nanocomposite at lower load (2â•›N). Tribochemical wear is the major mechanism of material removal at higher load (5 and 10â•›N). The observation of spalling induced by cracking of the nonprotective tribolayer and the formation of sheetlike agglomerated wear debris suggests severe wear at high load.
REFERENCES ╇ 1â•… K. J. Niihara. New design concept of structural ceramics—ceramic nanocomposites. Jpn. Ceram. Soc. 99 (1989), 974–982. ╇ 2â•… K. Niihara. Nanostructure design and mechanical properties of ceramic composites. J. Jpn. Soc. Powder Powder Metall. 37(2) (1990), 348–351. ╇ 3â•… K. Niihara, K. Izaki, and A. Nakahira. Nano-composite materials based on Si3N4 and SiC. Key Eng. Mater. 56 (1991), 319–326. ╇ 4â•… M. J. Mayo. Processing on nanocrystalline ceramics from ultra fine particles. Int. Mater. Rev. 41(3) (1996), 85–115. ╇ 5â•… H. Gleiter. Nanostructured materials: Basic concepts and microstructure. Acta Mater. 48 (2002), 1–29. ╇ 6â•… M. Sternitzke. Review: Structural ceramic nanocomposites. J. Eur. Ceram. Soc. 17 (1997), 1061–1082. ╇ 7â•… C. Suryanarayana. Nanocrystalline materials. Int. Mater. Rev. 40(2) (1995), 41–64. ╇ 8â•… G.-D. Zhan, J. D. Kuntz, J. Wan, and A. K. Mukherjee. Single-wall carbon nanotubes as attractive toughening agents in alumina based nanocomposites. Nature Materials 2 (2003), 38–42. ╇ 9â•… G. J. Gonzalez, B. Hockey, and G. J. Piermarini. High pressure compaction and sintering of nanosize γ-Al2O3 powder. Mater. Manuf. Processes 11 (1996), 951–967. 10â•… B. Basu, J.-H. Lee, and D.-Y. Kim, Development of nanocrystalline wear resistant Y-TZP ceramics. J. Am. Ceram. Soc. 87(9) (2004), 1771–1774. 11â•… M. Omorio. Sintering, consolidation, reaction and crystal growth by spark plasma sintering (SPS). Mater. Sci. Eng. A 287 (2002), 183–188.
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12â•… J. Hung, L. Gao, S. D. D. L. Torre, H. Miyamoto, and K. Miyamoto. Spark plasma sintering and mechanical properties of ZrO2(Y2O3)-Al2O3 composites. Mater. Lett. 43 (2000), 27–31. 13â•… T. Kusunose, T. Sekino, Y.-H. Choa, and K. Niihara. Machinability of silicon nitride/boron nitride nanocomposites. J. Am. Ceram. Soc. 85(11) (2002), 2689–2695. 14â•… S. Komarneni. Nanocomposites. J. Mater. Chem. 2 (1992), 1219–1230. 15â•… R. Roy, A. R. Roy, and M. D. Roy. Alternative perspectives on “quasi-crystallinity”: Non-uniformity and nanocomposites. Mater. Lett. 4 (1986), 323–328. 16â•… P. F. Becher. Microstructural design of toughened ceramics. J. Am. Ceram. Soc. 72 (1991), 255–269. 17â•… S. Yamanaka. Design and synthesis of functional layered nanocomposite. Ceram. Bull. 70 (1991), 1056–1058. 18â•… M. Kašiarová, E. Rudnayová, J. Dusza, M. Hnatko, P. Šajgalík, A. Merstallinger, and L. Kuzsella. Some tribological properties of a carbon-derived Si3N4/SiC nanocomposite. J. Eur. Ceram. Soc. 24 (2004), 3431–3435. 19â•… R. W. Davidge, P. C. Twigg, and F. L. Riley. Effects of silicon carbide nano-phase on the wet erosive wear of polycrystalline alumina. J. Eur. Ceram. Soc. 16 (1996), 700–802. 20â•… J. Rodriguez, A. Martin, J. Y. Pastro, J. L. Lorce, J. F. Bartlome, and J. S. Moya. Sliding wear of alumina/silicon carbide nanocomposites. J. Am. Ceram. Soc. 82(8) (1999), 1252–1254. 21â•… S. W. Lee, C. Morillo, J. L. Olivares, S. H. Kim, T. Sekino, K. Niihara, and B. J. Hockey. Tribological and microstructural analysis of Al2O3/TiO2 nanocomposites to use in the femoral head of hip replacement. Wear 255 (2003), 1040–1044. 22â•… T. Venkateswaran, D. Sarkar, and B. Basu. Tribological properties of WC-ZrO2 nanocomposites. J. Am. Ceram. Soc. 88(3) (2005), 691–697. 23â•… A. Mukhopadhyay, D. Chakrabarty, and B. Basu. Spark plasma sintered WC-ZrO2-Co nanocomposites with high fracture toughness and strength. J. Am. Ceram. Soc. 93(6) (2010), 1754–1763. 24â•… K. Biswas, A. Mukhopadhyay, B. Basu, and K. Chattopadhyay. Densification and microstructure development in spark plasma sintered WC-6 wt. % ZrO2 nanocomposites. J. Mater. Res. 22(6) (2007), 1491–1501. 25â•… B. Basu, J.-H. Lee, and D.-Y. Kim. Development of WC-ZrO2 nanocomposites by spark plasma sintering. J. Am. Ceram. Soc. 87(2) (2004), 317–319. 26â•… R. Mishra, B. Basu, and R. Balasubramaniam. Effect of grain size on the tribological behavior of nanocrystalline nickel. Mater. Sci. Eng. A 373(1–2) (2004), 370–373. 27â•… R. B. Waterhouse. Fretting wear, ASM handbook. ASM Int. 18 (1992), 242–256. 28â•… S. R. Brown (Ed.). Materials Evaluation under Fretting Conditions. ASTM Special Technical Publication 780, Warminster, PA, 1982, 780. 29â•… D. Klaffke. Fretting wear of ceramics. Tribol. Int. 22 (1989), 89–101. 30â•… M. Kalin and J. Vizintin. Use of equations for wear volume determination in fretting experiments. Wear 237 (2000), 39–48. 31â•… D. Sarkar, T. Venkateswaran, and B. Basu. Pressureless sintering and tribological properties of WCZrO2 composites. J. Eur. Ceram. Soc. 25 (2005), 1603–1610. 32â•… D. Sarkar, S. Ahn, S. Kang, and B. Basu. Fretting wear of TiCN-Ni cermet: Influence of secondary carbide content. P/M Sci. Technol. Briefs 5(2) (2003), 5–11. 33â•… J. P. Tu, W. Rong, S. Y. Guo, and Y. Z. Yang. Dry sliding wear behavior of in situ Cu-TiB2 nanocomposites against medium carbon steel. Wear 255 (2003), 832–835. 34â•… B. Bhushan. Principles and Applications of Tribology. John Wiley & Sons, New York, 1999, 487. 35â•… S. G. Roberts. Depths of cracks produced by abrasion of brittle materials. Scr. Mater. 40(1) (1999), 101–108. 36â•… E. Rabinowicz. Friction and Wear of Materials, 2nd ed. John Wiley & Sons, New York, 1995, 178.
SECTION
V
LIGHTWEIGHT COMPOSITES AND CERMETS
CHAPTER
22
OVERVIEW: LIGHTWEIGHT METAL MATRIX COMPOSITES AND CERMETS In this section of the book, the focus of discussion is on the friction and wear properties of lightweight ceramics and some new cermet materials. The results of systematic experiments involving fretting of Mg–SiCp composites are presented in Chapter 23. At each stage the operating mechanisms are identified and correlated with changes that occurred in frictional behavior. In addition, this section also summarizes an extensive experimental study being carried out on TiCN–Ni-based cermets, reinforced with various secondary carbides (WC, NbC, TaC, Hfc). Various aspects of such studies include (1) the role played by different types of secondary carbides in influencing the tribochemical wear behavior of the cermets, (2) the influence of varying amounts of WC content on load-dependent sliding wear of TiCN–Ni–WC cermets, and (3) high-temperature sliding wear behavior of cermets. In a different chapter, the tribological properties of a new-generation cermet, that is, mixed carbide cermet [(W,Ti)C–Co], are discussed. Both the TiCN–Ni-based and (W,Ti)C–Co cermets are being developed as potential substitutes for conventional WC–Co cermets, and therefore, their results are significant to the tribology community.
22.1 DEVELOPMENT OF METAL MATRIX COMPOSITES Increasing demand for stiffer, stronger, and lightweight materials has created a large interest in the development of metal matrix composites (MMCs). The selection of proper matrix, reinforcement, and processing technique are necessary for realizing the advantages of MMCs. It is possible to tailor the specific properties of the MMCs such as stiffness and strength by reinforcing with a ceramic material. MMCs are suitable for several engineering applications in the aerospace and automotive sectors, for example connecting rods, pistons, and cylinder liners.1 This is because particulatereinforced MMCs show distinct advantages such as low density, high load-carrying capacity, and good wear resistance. High wear resistance of particle-reinforced Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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MMCs is mainly due to the presence of ceramic particles, which prevent the continuous removal of the metallic matrix at the tribocontact. The wear and friction behavior of MMCs depends on the type of matrix alloy, the reinforcement chemistry and volume, the counterbody material, and the testing conditions.2 A literature review shows that extensive work has been done to understand the dry sliding wear behavior of Al–SiCp composites.3–12 In the recent past, lighter Mg-alloys-based MMCs have received greater attention. Mg-alloys are potentially useful where higher specific mechanical properties are desirable.13 The elastic modulus, which is low for Mg and its alloys, can be improved by using harder and stiffer ceramic particle reinforcements. However, there are limited studies on the tribological properties of Mg-based MMCs.13–17 In many automotive and aerospace applications, Mg-based composites are preferred over Al-based composites due to the attractive combination of light weight, superior strength, and better performance. In one of the early works by Alahelisten et al.2 on sliding, abrasion, and erosion behavior of alumina fiber-reinforced Mg and Mg–9Al–Zn matrix composites, it is reported that wear resistance can be increased with an increase in volume fraction of fiber. An increase in wear resistance was observed in two-body abrasion, whereas it decreased in three-body abrasion and erosion. Although the wear mechanisms were not explicitly explored, Saravanan and Surappa13 reported an increase in wear resistance of Mg–(30 vol%)SiCp composites during adhesive wear, compared with base Mg. Sharma et al.15 investigated the sliding wear behavior of Mg-alloy–feldspar composite against EN24 steel. The wear rate decreased with increasing amount of feldspar particles. It was indicated that the interparticle distance of SiC particles affect the tribological properties of composites.16 A 2003 study by Lim et al.17 gives an account of the severe adhesive wear contacts on Mg–SiCp composites against steel. An important observation of this study was that the wear rate, measured on Mg–(9â•›wt %)Al–(3 vol%)SiCp using a pin-on-disk tribometer, initially decreased and then increased with increasing sliding speed. It was shown that, at lower load (10â•›N), the composite reveals a slight increase in wear resistance due to its improved load-bearing capacity and its ability to maintain a stable oxide film. At 30â•›N, the improvement in wear resistance was observed only at moderate speeds (1 and 2â•›m/s), but higher speeds resulted in delamination and abrasion, resulting in a decrease in the wear resistance of the composite. The useful range of these composites was observed to be around 5â•›m/s, where frictional heating results in melting of the composite and gross plastic deformation of the pin surface is observed. Unlike in the case of Al–SiCp composites,4,9,11 the earlier research results indicate that the decrease in wear rate is not substantial in Mg-based composites with the addition of SiC particulates. The occurrence of tribochemical reactions at the tribocontact and their subsequent effect on friction and wear properties have been well understood in the case of ceramics,18–22 but not much information is available in the case of composites. Most of the reported work on the tribology of ceramics was carried out using pinon-disk setups, that is, under unidirectional sliding conditions. However, no work has reported the fretting wear of MMCs. Another important aspect of fretting is the development of fatigue cracks in the damaged region, resulting in decreased fatigue strength of a cyclically loaded
â•… 355
22.1 Development of Metal Matrix Composites
component. Extensive studies had been carried out in the past to understand the fretting wear and fretting fatigue behavior, both at room temperature and at high temperature, of a number of metallic alloys, including Ti,23,24 Ni–Cr,25 Ti–6Al–4V26, and steels.27,28 Since 2000, the fretting wear behavior of engineering ceramics and their composites have begun to be investigated.29–33 In 2006, the fretting wear behavior of a newly developed Al-based in situ MMC reinforced with iron aluminide particles was studied.34 The phase composition and in situ reinforcement content were varied by creating variation in the hot pressing temperature and volume percent of precursor powders (nanosized Fe2O3). The wear experiments were carried out in the gross slip fretting regime to understand the composites’ tribological behavior against bearing steel in the ambient conditions of temperature (22–25°C) and humidity (50–55% relative humidity). The tribological data obtained indicate a reduction in the coefficient of friction (COF) of the in situ composites with 20 vol% reinforcement, irrespective of the processing temperature. The most interesting results show that a significantly low value of COF (0.25) is recorded for 20 vol%-reinforced composite, hot pressed at 700°C. The wear data show that the wear volume, estimated mainly from the wear scar diameter, decreases considerably (20 vol%) and then undergoes further increase with the increase in the reinforcement content under the investigated fretting condition. The important result of the study is that the 20 vol%-reinforced Al composite exhibits the highest wear resistance when hot pressed at 800°C. The experimental results suggest that both the reinforcement content and the hot pressing temperature should be optimized to obtain a superior combination of COF and wear rate. It was observed that extensive plastic deformation causes more wear in unreinforced Al. The wear induced by deformation is explained on the basis of formation of persistent slip bands and subsequent cracking as a result of cyclic deformation during a large number of fretting cycles. The improvement in hardness due to harder aluminide reinforcements results in an increase in wear resistance of 20 vol%reinforced composite, and this has been observed to lead to better wear resistance. For 40 vol% reinforcement, the wear mechanism undergoes a transition from predominantly abrasion and deformation to predominantly abrasion and tribochemical wear. This resulted in larger wear loss in the composites with 40 vol% in situ reinforcement. In another study,35 the tribological behavior of Al-based composites with reinforcement of in situ TixAly and Al2O3 particles are reported. An important result is that Al-based composites with 20 vol% reinforcement show a very low COF of 0.2 under unlubricated conditions. Also, the wear volume measured is around five times lower with 20 vol% composites compared with unreinforced Al. The experimental results indicate that the volume of reinforcement should be optimized in order to obtain lower COF in in situ reinforced Al composites. While the damage induced by deformation is the major fretting wear mechanism for unreinforced Al, transfer layer formation and tribochemical wear are the dominant wear mechanisms for the in situ composites. The formation of in situ intermetallic particulates lowers the severity of wear damage induced by plastic deformation. Also, the formation of a protective tribochemical layer results in higher wear resistance of 20 vol%-reinforced composites.
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22.2 DEVELOPMENT OF CERMETS Ceramic-based materials are being used to a large extent to design wear-resistant components for automobiles and aircrafts. This is especially true for engine parts, for which achieving a more efficient combustion and significant fuel savings are possible with these materials. Among various tribological materials, cermets are known for their excellent wear resistance under unlubricated conditions. It is recognized that microstructural and compositional aspects such as grain size, binder content, and amount of secondary carbide influence the wear behavior of cermets. In particular, WC–Co hardmetals are being considered as traditional cutting tool inserts, because of their high hardness and wear resistance.36–39 Recent attempts have been made to use TiC-based composites for high-temperature applications, because of their favorable material properties.40–43 It has been reported44 that with increasing VC the wear resistance of (Ti,V)C–Co cermets decreases. The conclusion drawn was that microstructure, rather than hardness, is the dominant factor in determining wear rate of such cermets. Kameo et al.,45 in their investigation on the sliding wear of cermets and ceramic composites, reported that wear of the steel counterbody is minimal against cermets, compared with that against ceramic composites containing ceramic reinforcements. Since the 1990s, work has been under way to develop TiCN-based cermets as cutting inserts for consideration as an alternative to WC–Co cermets.46–50 The microstructural features of the cermets and the bonding between different microstructural phases play a crucial role in controlling the wear of different materials. An improvement in structural homogeneity leads to an increase in the wear resistance of ultrafine TiCN-based cermets.50 The adhesion and diffusion between rubbing surfaces increase with increase in load and speed, resulting in an increase in wear of TiCN from sliding against steel in dry conditions.51 The development of microstructure during sintering for a range of TiCN-based cermets has been relatively well reported in the literature.52–58 Typical core–rim structure resulting from the dissolution and reprecipitation during sintering of the cermets significantly influences the material properties. The rim phase consists of Ti(CN) cores and transitional metal carbides. Undissolved Ti(CN) cores can result in excellent wear resistance. On the other hand, crack interactions with the solid solutions, such as (Ti,Nb)(C,N) rim structures and Ni binder phase, determine a material’s toughness and integrity. Further, the addition of secondary carbides, such as WC, TaC, HfC, NbC, and Mo2C, increases the density, high-temperature strength, and resistance to hightemperature deformation by improving hot hardness and thermal shock resistance,46,48,49,59–68 as schematically shown in Figure 22.1. The addition of WC is indispensable, in many cases, for achieving densification and fracture toughness,46 whereas NbC improves the interrupted cutting performance by increasing hot hardness and thermoshock resistance at elevated temperatures.49,54 The addition of HfC increases high-temperature strength and wear resistance,65,66 and addition of TaC improves the resistance to high-temperature deformation.62 The addition of Mo2C also results in an increase in toughness, a decrease in particle size, suppression of graphite formation, and a contribution toward high-temperature strength.64
22.2 Development of Cermets
â•… 357
Co/Ni to aid in densification at TiN/WC/TiC addition to enhance hardness, wear resistance
lower temperature via liquid phase sintering as well as to improve toughness property
Development of TiCN-based cermets
HfC addition to improve hot hardness/thermal shock resistance
Addition of secondary carbides to be tailored to maintain a balance between wear resistance and other mechanical properties
Figure 22.1â•… Summary of various aspects to be considered while developing TiCN-based cermets.
The superior properties such as high melting point, hardness, thermal conductivity, wear resistance, and chemical stability make TiCN-based cermets potential materials for cutting tools.46,49,59,60 This class of materials has better surface finishing, excellent chip and tolerance control, and dimensional accuracy of the workpiece, compared with WC–Co materials. The nitrogen content affects the properties of the titanium carbonitride-based cermets in the same manner, as discussed for titanium carbonitride ceramics. For example, the increase in thermal conductivity with nitrogen content resulted in improved thermal shock resistance.69,70 The fracture toughness of cermets increases with nitrogen content [up to N/(C╯+╯N)╯=╯0.5]. For N/(C╯+╯N)╯>╯0.5, a decreasing trend in fracture toughness is observed due to grain growth.71,72 The wear resistance of the cermet increases with nitrogen content, thus making high-speed machining possible.73,74 Also, the increase in hardness and strength with decreasing grain size of TiC-rich TiCN cermets (due to increasing nitrogen content), result in longer tool life and better wear resistance.73,75 Furthermore, the modulus of elasticity and oxidation resistance increase,71,76 while the steady-state creep rate decreases77,78 with nitrogen content in the TiCN-based cermets. Based on the reported literature, the following general observations about Tibased cermets compared with WC–Co hard materials can be listed60: 1. They have lower density (around 6.0â•›g/cm3) compared with WC–Co (around 13.0â•›g/cm3). 2. They have similar hardness (Hv30╯=╯14–15â•›GPa) but lower strength (1.4– 1.9â•›GPa against 3.5â•›GPa) and toughness values (7–11â•›MPa m1/2 against 20– 24â•›MPa m1/2). 3. The coefficient of thermal expansion is higher for WC–Co alloys (6╯×╯10−6â•›K−1) as compared with that of Ti-based cermets (9.0╯×╯10−6â•›K−1).
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4. Their thermal conductivity is lower (10â•›W/m k) compared with that of WC–Co alloys (15â•›W/m K). 5. The oxidation resistance of TiCN is much higher than that of WC. As part of ongoing efforts to refine the material composition for improvement in high-temperature machining performance, refractory carbides such as WC, NbC, and TaC are added to baseline TiCN–Ni cermets to enhance high-temperature strength and hardness properties.62 Ni binder (typically up to 20â•›wt %) is commonly used for the densification of these cermets. The microstructural characterization of these novel cermets have been discussed in more detail elsewhere.12,52,79,80 Typical sintered microstructure of these cermets consists of a core–rim grain structure, formed during a dissolution-and-reprecipitation process. Besides the aspects of the microstructural investigation in various compositionally designed cermets, the tribological study of TiCN cermets has not been considered to any significant degree. More recently, the performance of this significant class of materials has been studied under different environments and tribological conditions, such as fretting,81 sliding,82 erosive,83 corrosive,84 and die-sinking electrodischarge machining85 by our research group. The results on the tribological properties of TiCN–(20â•›wt %)Ni–(xâ•›wt %)WC (where x is 5, 10, 15, 20, 25) cermets are discussed in Chapter 24. In this section, various wear mechanisms, acting as a function of cermet microstructure, under varying operating conditions are discussed. In a separate chapter (Chapter 25), the background of the development of a new class of mixed carbide cermets, that is, (W,Ti)C–Co, is discussed along with their mechanical and tribological properties.
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38â•… J. Pirso, S. Letunovitš, and M. Viljus. Friction and wear behaviour of cemented carbides. Wear 257 (2004), 257–265. 39â•… K. Jia and T. E. Fischer. Sliding wear of conventional and nanostructured cemented carbides. Wear 203–204 (1997), 310–318. 40â•… T. Hisakado and N. Hashizume. Effect of normal loads on the friction and wear properties of metals and ceramic against cermet in vacuum. Wear 237 (2000), 98–106. 41â•… C. C. Degnan, P. H. Shipway, and J. V. Wood. Elevated temperature sliding wear behaviour of TiC–reinforced steel matrix composites. Wear 251 (2001), 1444–1451. 42â•… J. Pirso, M. Viljus, and S. Letunovitš. Sliding wear of TiC–NiMo cermets. Tribol. Int. 37 (2004), 817–824. 43â•… M. Komac and S. Novak. Mechanical and wear behaviour of TiC-cemented carbides. Int. J. Refract. Met. Hard Mater. 4 (1985), 21–26. 44â•… F. Arenas, C. Rondón, and R. Sepu′lveda. Friction and tribological behavior of (Ti,V)C–Co cermets. J. Mater. Process. Technol. 143–144 (2003), 822–826. 45â•… K. Kameo, K. Friedrich, J. F. Bartolome, M. Diaz, L.-E. Sonia, and J. S. Moya. Sliding wear of ceramics and cermets against steel. J. Eur. Ceram. Soc. 23 (2003), 2867–2877. 46â•… S. Zhang. Titanium carbonitride–based cermets: Process and properties. Mater. Sci. Eng. A 163 (1993), 141–148. 47â•… G. E. D’Errico, E. Buglisosi, and E. Guglielmi. Tool-life reliability of cermet inserts in milling tests. J. Mater. Process. Technol. 77 (1998), 337–343. 48â•… W. T. Kwon, J. S. Park, S.-W. Kim, and S. Kang. Effect of WC and group IV carbides on the cutting performance of Ti(CN) cermet tools. Int. J. Mach. Tools and Manuf. 44(4) (2004), 341–346. 49â•… P. Ettmayer, H. Kolaska, W. Lengauer, and K. Dreyer. Ti(CN) cermets: Metallurgy and properties. Int. J. Refract. Met. Hard Mater. 13 (1995), 343–351. 50â•… E. T. Jeon, J. Joardar, and S. Kang. Microstructure and tribo-mechanical properties of ultrafine Ti(CN) cermets. Int. J. Refract. Met. Hard Mater. 20 (2002), 207–211. 51â•… X. Zhao, J. Liu, B. Zhu, Z. Luo, and H. Miaon. Effects of lubricants on friction and wear of Ti (CN)/ l045 steel sliding pairs. Tribol. Int. 30(3) (1997), 177–182. 52â•… S. Ahn and S. Kang. Effect of various carbides on the dissolution behavior of Ti(C0.7 N0.3) in a Ti (C0.7 N0.3)-30 Ni system. Int. J. Refract. Met. Hard Mater. 19 (2001), 539–545. 53â•… S.-Y. Ahn and S. Kang. Formation of core/rim structure in Ti(C,N)–WC–Ni cermets via a dissolution and precipitation process. J. Am. Ceram. Soc. 83–86 (2000), 1489–1494. 54â•… S.-Y. Ahn, S.-W. Kim, and S. Kang. Microstructure of Ti(C,N)–WC–NbC–Ni cermets. J. Am. Ceram. Soc. 84(4) (2001), 843–849. 55â•… P. Ettmayer, H. Kolaska, and K. Dreyer. Effect of the sintering atmosphere on the properties of cermets. Powder Metall. Int. 23 (1991), 224–230. 56â•… M. G. Gee, M. J. Reece, and B. Roebuck. High resolution electron microscopy of Ti(C, N) cermets. J. Hard Mater. 3 (1992), 119–141. 57â•… S. Kim, K.-H. Min, and S. Kang. Rim structure in Ti(C0.7N0.3)-WC-Ni cermets. J. Am. Ceram. Soc. 86(10) (2003), 1761–1766. 58â•… F. Monteverde, V. Medri, and A. Bellosi. Microstructure of hot-pressed Ti(C,N)- based cermets. J. Eur. Ceram. Soc. 22 (2002), 2587–2593. 59â•… H. Pastor. Titanium-carbonitride-based hard alloy for cutting tools. Mater. Sci. Eng. A 105–106 (1998), 401–409. 60â•… E. B. Clark and B. Roebuck. Extending the application areas for titanium carbonitride cermets. Refract. Met. Hard Mater. 11 (1992), 23–27. 61â•… T. Laoui and O. Van der Biest. Effect of TiC addition on the microstructure and properties of Ti(C,N)–WC–Co–Ni cermet. J. Mater. Sci. Lett. 13 (1994), 1530–1532. 62â•… S. Kang. Some issues in Ti(CN)–WC–TaC cermets. Mater. Sci. Eng. A 209 (1996), 306–312. 63â•… P. Lindahl, P. Gustafson, U. Rolander, L. Stals, and H.-O. Andren. Microstructure of model cermets with high Mo or W content. Int. J. Refract. Met. Hard Mater. 17 (1999), 411–421. 64â•… M. Komac and D. Lange. The influence of MoC and NbC additions on microstructure and mechanical Properties of TiC based cemented carbides. Int. J. Powder Met. Powder Tech. 18(4) (1982), 313–322. 65â•… Y. Ozaki and R. H. Zee. Coarsening of hafnium carbide particles in tungsten. J. Mater. Sci. 30(13) (1995), 3421–3428.
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66â•… S. Mun and S. Kang. Effect of HfC addition of microstructure of Ti (CN)-Ni system. Powder Met. 42(3) (1999), 251–256. 67â•… D. S. Park, C. Park, and Y. D. Lee. Oxidation of Ti(C,N)-based ceramics exposed at 1373 K in air. J. Am. Ceram. Soc. 83(3) (2000), 672–674. 68â•… H. Pastor. Present status and development of tool materials: Part 1 Cutting tools. Int. J. Refract. Met. Hard Mater. 6(4) (1987), 196–209. 69â•… Y. Shimizu, M. Tabioka, N. Kitagawa, and T. Nomura. Development of new T110A and T130A tough cermet. Met. Powder Rep. 44(12) (1989), 827–834. 70â•… H. Tanaka. Relationship between thermal, mechanical properties and cutting performance of TiNTiC cermet, in Cutting Tool Materials, F. W. Gorsler (Ed.). Conference Proceedings Amer Soc for Metal, 1980, 349–361. 71â•… J. Suzuki and H. Tanaka. TiN-base Cermet Cutting Tool, SME Technical Paper, 1989, TE89-345. 72â•… H. Yasui, H. Tanaka, and T. Kobayashi. Effect of N/C ratio on mechanical properties in Ti(C, N) base cermet, in Metal and Ceramic Matrix Composites: Processing, Modeling and Mechanical Behavior, R. B. Bhagat, A. H. Clauer, P. Kumar, and A. M. Ritter (Eds.). The Minerals, Metals and Materials Society, 1990, 67–74. 73â•… D. Moskowitz, L. L. Terner, and M. Humenik Jr. Some physical and metal-cutting properties of titanium carbonitride base materials, in Science of Hard Materials, A. Almond, C. A. Brookes, and R. Warren (Eds.). Proc. 2nd Int. Conf. on the Science of Hard Materials Rhodes, September 23–28, 1984, Inst. Phys. Conf. Ser. 75, 1987, 605–617. 74â•… J. Larsen-Basse. Abrasive wear of some titanium carbonitride-based cermets. Mater. Sci. Eng. A 105–106 (1988), 395–400. 75â•… M. Fukuhara and H. Mitani. Effect of nitrogen content on grain growth in Ti(C,N)-Ni-Mo sintered alloy. Powder Met. Int. 14 (1982), 196–200. 76â•… H. Suzuki, H. Matsubara, and A. Matsuo. High temperature mechanical properties of Ti(C, N)–Mo sintered compacts. Powder Powder Met. 32(5) (1985), 24–27. 77â•… H. Matsubara and T. Sakuma. Microstructure and Mechanical Properties of Titanium Carbonitride Base Cermets, 10th ed., Sintering ‘87 Vol. 2, Elsevier, 1988, 1269–1274. 78â•… M. Fährmann. High temperature creep of Ti(C, N)-based cermets. Int. J. Refract. Met. Hard Mater. 8(4) (1989), 219–223. 79â•… S. Ahn and S. Kang. Formation of core/rim structures in Ti(C,N)-WC-Ni Cermets via a dissolution and precipitation process. J. Am. Ceram. Soc. 83(6) (2000), 1489–1494. 80â•… F. Qi and S. Kang. A study on microstructural changes in Ti(CN)–NbC–Ni cermets. Mater. Sci. Eng. A 251 (1998), 276–285. 81â•… D. Sarkar, B. V. Manoj Kumar, S. Ahn, S. Kang, and B. Basu. Fretting wear behavior of Ti(CN)based advanced cermets. Key Eng. Mater. 264–268 (2004), 1115–1118. 82â•… B. V. Manoj Kumar, B. Basu, M. Kalin, and J. Vizintin. Load dependent transition in sliding wear properties of TiCN-WC-Ni cermets. J. Am. Ceram. Soc. 90(5) (2007), 1534–1540. 83â•… B. V. Manoj Kumar, B. Basu, S. Kang, and J. Ramkumar. Erosion wear behavior of TiCN-Ni cermets containing secondary carbides (WC/NbC/TaC). J. Am. Ceram. Soc. 89(12) (2006), 3827–3831. 84â•… B. V. Manoj Kumar, R. Balasubramaniam, and B. Basu. Electrochemical behavior of TiCN-Ni based cermets. J. Am. Ceram. Soc. 90(1) (2007), 205–210. 85â•… B. V. Manoj Kumar, J. Ramkumar, B. Basu, and S. Kang. Electro-discharge machining performance of TiCN-based cermets. Int. J. Refract. Met. Hard Mater. 25 (2007), 293–299.
CH A P T E R
23
CASE STUDY: MAGNESIUM– SILICON CARBIDE PARTICULATE-REINFORCED COMPOSITES In striving toward achieving lightweight wear-resistant composite materials, considerable effort has been put forward to develop Mg–SiC composites. Considering the importance of lightweight metal matrix composites (MMCs), this chapter summarizes the experimental results obtained while investigating the tribological properties of Mg–SiC composites.1,2 In particular, the roles of SiC addition and load on the fretting wear of Mg–SiCp composites are analyzed. The tribological properties of Mg–SiCp/steel contacts are reported, with varying SiC particulate volume fraction and load, and then compared with those of pure Mg/steel contacts. The tribochemical reactions are observed to be dominant. In the early stage of fretting, the dominant wear mechanism is that of abrasive wear due to the formation of MgO. The harder SiC particulates present in the composites cause severe abrasion with the progress of fretting. An important observation has been the formation of a dense hydrous magnesium silicate (DHMS), clinoenstatite layer as the tribochemical reaction product. The experimental results reveal that the SiC particulate loading (especially higher volume fractions) has a more significant effect on the coefficient of friction (COF) compared with the normal load applied during fretting.
23.1 INTRODUCTION In the overview given in Chapter 22, a brief background to the development of Mgbased MMCs is provided. As far as the tribological properties of these composites are concerned, most of the published research on Mg-based composites is focused on severe conditions (load╯>╯10â•›N and sliding speed╯>╯0.2â•›m/s) in adhesive wear contacts with limited studies in the mild wear regime. Also, since the tribological performance of these materials is greatly influenced by the fretting contacts in various applications,1–12 it is important to understand the properties of Mg–SiCp composites with variation of different fretting parameters. The first part of the Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
362
23.3 Load-Dependent Friction and Wear Properties
â•… 363
chapter reports the performance of Mg–SiCp composites with varying load (2–10â•›N) and SiC particulate content (9.8 and 26.3â•›wtâ•›%) at very low sliding speed (0.0016â•›m/s) with an effort to critically understand the dominant wear mechanisms. The second part discusses the fretting duration dependent wear behavior and more details reported elsewhere.1,2
23.2 MATERIALS AND EXPERIMENTS Samples of Mg–SiCp composites containing 9.8 and 26.3â•›wtâ•›% (or 5.6 and 16.3â•›volâ•›%) SiC particulate reinforcement were used. These composites were processed by a novel mechanical disintegration deposition technique, in which the composite melt of Mg and α-SiC (25â•›µm) particulates were disintegrated into multiple streams by a mechanical device. The disintegrated slurry was deposited into a metallic substrate. Subsequently, 10-mm rods were produced by extrusion of the ingots. The detailed description of this process can be found elsewhere.3 The comparisons of the mechanical properties of these composites with those of pure Mg are given in Table 23.1. Typical microstructures of the Mg base matrix and its composites studied in this investigation are presented in Figure 23.1. The optical micrographs show SiC particulates (about 25â•›µm in size) uniformly distributed in the recrystallized Mg matrix having equiaxed grain (31╯±â•¯11â•›µm) morphology. There appears to be good bonding of the SiC particles with the matrix. The intercept method was used to determine the interparticle distance between SiC particulates. The average values of interparticle distance in the case of Mg–SiCp (9.8â•›wtâ•›%) and Mg–SiCp (26.3â•›wtâ•›%) composites are 81.7 and 44.1â•›µm, respectively. It is also noted that the reinforcing of base Mg matrix with SiCP increases the hardness from HR 15T 39 (base Mg) to HR 15T 50 for Mg–SiCp (9.8â•›wtâ•›%) and HR 15T 53 for Mg–SiCp (26.3â•›wtâ•›%) composites (see Table 23.1).
23.3 LOAD-DEPENDENT FRICTION AND WEAR PROPERTIES The tribological properties of pure Mg and its composites were studied under fretting conditions. The effect of load was studied while the other fretting variables (10,000 cycles, 100-µm amplitude, and 8-Hz frequency) were kept constant. The variation in COF for Mg and SiCp-reinforced Mg at different loads are plotted in Figure 23.2. During the running-in period (first 500 cycles), COF values increase significantly TABLE 23.1â•… Mechanical Properties of Base Mg Matrix and Mg–SiCp Composites2
Material Mg Mg–(9.8â•›wtâ•›%)SiC Mg–(26.3â•›wtâ•›%)SiC
E (GPa)
0.2%YS (MPa)
UTS (MPa)
Ductility (%)
Hardness (HR15T)
40.7 45.7 53.0
125╯±â•¯13 128╯±â•¯3 144╯±â•¯6
216╯±â•¯3 189╯±â•¯14 216╯±â•¯3
9╯±â•¯5 4╯±â•¯1 3╯±â•¯1
39╯±â•¯1 50╯±â•¯1 53╯±â•¯2
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CHAPTER 23â•… MAGNESIUM–SILICON CARBIDE PARTICULATE-REINFORCED COMPOSITES
50 mm (a)
(b)
(c)
Figure 23.1â•… Microstructures of (a) base Mg, (b) Mg–SiCp (9.8â•›wtâ•›%), and (c) Mg–SiCp (26.3â•›wtâ•›%) composites.2
and then decrease to attain a steady-state value. Pure Mg exhibits a COF value of 0.23, which gradually increases to 0.31 with load increasing from 2 to 10â•›N, respectively. The following important observations are made from the COF data obtained on Mg and its composites: (1) In pure Mg at lower loads, maximum COF value during running-in period is around 0.3, which increases with load to more than 0.5. Under similar conditions at lower loads, the composite materials exhibit higher COF values (0.33–0.35). At higher loads, COF values do not vary much for the Mg–SiCp (9.8â•›wtâ•›%) composites. However, peak COF values decrease for composites with large volume fraction (26.3â•›wtâ•›%) of the reinforcement. (2) In all the materials, COF values gradually decrease (from running-in period) and attain a steady-state value. The steady-state COF values increase with SiC particulates (from 9.8 to 26.3â•›wtâ•›%) at very low loads (2â•›N), but the reverse effect was observed at high loads (8 or 10â•›N). The steady-state COF values decrease to 0.24 with the volume fraction of SiC particulates and load reaching a maximum. It may be noted that the difference in steady-state COF between base Mg and Mg–SiCp (26.3â•›wtâ•›%) is minimal. This indicates that the SiC particulates present in the composite do not significantly affect the friction behavior after reaching steady state. (3) In the fretting wear process, formation of debris and tribochemical products followed by removal of debris and tribochemical layer results in the observed fluctuations in friction curves as noticed in Figure 23.2. In other words, the wear mechanisms undergo change starting from the running-in period to the steady-state period.
23.3 Load-Dependent Friction and Wear Properties
0.8
â•… 365
Base Mg
0.7 0.6 COF
0.5
10 N 8N 6N
0.4 0.3 0.2
2N
4N
0.1 0
2000
4000 6000 8000 Number of Cycles
10,000
(a) 0.8
0.8
Mg–SiCp(9.8 wt. %)
0.7
0.6
0.6 0.5 4N
0.4
0.5
6N 10 N 8N
0.3 0.2
0
2000
4N
2N
0.4 0.3 0.2
2N
0.1
COF
COF
Mg–SiCp(26.3 wt. %)
0.7
8 N 10 N
6N
0.1
4000 6000 8000 10,000 Number of Cycles
(b)
0
2000
4000 6000 8000 Number of Cycles
10,000
(c)
Figure 23.2â•… The frictional behavior of monolithic Mg and its composites: (a) pure Mg, (b) Mg–SiCp (9.8â•›wtâ•›%) composite, and (c) Mg–SiCp (26.3â•›wtâ•›%) composite, during fretting against bearing steel. Fretting conditions: load, 2–10â•›N; duration, 100,000 cycles; frequency, 8â•›Hz; and stroke length, 100â•›µm.2
Detailed microstructural investigation using a scanning electron microscope with energy-dispersive spectrometry (SEM-EDX), an electron probe microanalyzer (EPMA), and Raman spectroscopy was carried out on the worn surfaces of Mg and its composites to understand the mechanisms involved in the fretting wear. The details of the worn surfaces were observed and analyzed to understand the operating mechanisms (Fig. 23.3). At lower loads (2â•›N), the oxide debris in the range of 2–20â•›µm is seen to accumulate at the edge of the scar (marked as A). It is also observed that the base Mg matrix undergoes deformation on the surface (marked as B). The beginning of tribolayer formation, smearing, and removal over the worn surfaces are observed at higher loads (4–6â•›N), as shown in Figure 23.3b,c. At still higher loads (8â•›N), removal of the tribochemical layer and further oxidation take place underneath the surface (Fig. 23.3d). From these observations, it is clear that the tribochemical product formed during fretting is soft and spreads easily on the surfaces.
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A
B
20 mm
2N
20 mm
20 mm
4N (b)
(a)
6N
20 mm
(c)
8N (d)
Figure 23.3â•… Fretted surfaces of pure Mg flats after fretting at different loads. Arrow marks indicate fretting direction. Fretting conditions: duration, 100,000 cycles; frequency, 8â•›Hz; stroke length, 100â•›µm; counterbody, bearing steel.1
Compared with fretted surfaces on base Mg, the additional feature that is occasionally observed on Mg–(9.8%)SiC composite surfaces is the formation of needle-like particles (Fig. 23.4). It has been reported that needle-like wear debris particles are formed in ceramics (both oxides and nonoxides), when they were subjected to wear in a moist atmosphere.4,7–9 Important features such as accumulation of (hydrous) oxide at the periphery of a scar, smearing, and fracture and removal of the tribochemical layer (as sheets) are illustrated in Figure 23.4. Raman spectroscopy was employed for the analysis of the chemical composition of the triboproducts. Raman analysis (Fig. 23.5) shows that the chemistry of the reaction products matches with that of DHMS, which is likely to be clinoenstatite (MgSiO3).10 The presence of pyroxene [(Mg,Fe)SiO3] was not detected when analyzing the Raman spectroscopic data. This suggests that the amount of iron oxide debris particles transferred to the worn flat sample is minimal. Also, a weaker Raman peak at 263â•›cm−1 shows the presence of SiO2 (see Fig. 23.5).
23.4 FRETTING-DURATION-DEPENDENT TRIBOLOGICAL PROPERTIES The wear rates were computed from the laser profilometry analysis and are shown in Figure 23.6. The wear rate data show that the wear rate of base Mg and Mg–SiC
23.4 Fretting-Duration-Dependent Tribological Properties
5 mm
10 mm
8N
8N (b)
(a)
10 mm
â•… 367
10 mm
8N (c)
8N (d)
Relative intensity
Figure 23.4â•… Worn surfaces of Mg–SiCp (9.8â•›wtâ•›%) composites at 8-N load: (a) formation of needle-like particles; (b) accumulation of hydrous oxide; (c) smearing; and (d) fracture of tribochemical layer. Arrows indicate the fretting direction. Fretting conditions: duration, 100,000 cycles; frequency, 8â•›Hz; stroke length, 100â•›µm; counterbody, bearing steel.1
250
300
350 400 Raman shift cm–1
450
500
Figure 23.5â•… Raman peaks showing the presence of DHMS (at 379 and 512â•›cm−1) and silica (at 263â•›cm−1) on the Mg–SiCp (26.3â•›wtâ•›%) composite worn surface. The exact peak positions are marked below with arrows. Fretting conditions: load, 8â•›N; duration, 5000 cycles; frequency, 8â•›Hz; and stroke length, 100â•›µm.2
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Specific wear rate (10–2 mm3/N m)
5.0 Base Mg Mg–SiCp (9.8 wt.%) Mg–SiCp (26.3 wt.%)
4.5 4.0 3.5 3.0 2.5 2.0 1.5 1.0 0.5 0.0
1K
2K Number of cycles
5K
Figure 23.6â•… Variation of specific wear rate for base Mg and Mg–SiCp composites against bearing steel at different number of fretting cycles. Fretting conditions: load, 8â•›N; frequency, 8â•›Hz; and stroke length, 100â•›µm.1
composites vary over the range of 10−2–10−4â•›mm3/Nâ•›m. It can be noted here that the fretting wear rate of ceramics and ceramic composites typically varies over the range of 10−6–10−7â•›mm3/Nâ•›m.5 For base Mg and Mg–SiCp (26.3â•›wtâ•›%) composite, there was a systematic decrease in the wear rate with fretting duration. However, for composites with 9.8â•›wtâ•›% SiCp, the specific wear rate decreased until 2000 cycles, followed by an increase. For 2000 and 5000 fretting cycles, wear rate that is lower by consistently one order of magnitude was achieved with Mg–SiCp (26.3â•›wtâ•›%) composites compared with that of Mg–SiCp (9.8â•›wtâ•›%) composites. The wear mechanisms were understood with the help of detailed microstructural investigation using SEM and Raman spectroscopy. Figure 23.7 shows the SEM images (backscattered contrast, in backscattered electron [BSE] mode) of base Mg after fretting for different numbers of cycles. Under the investigated experimental conditions, oxidation of the worn surface begins as soon as the fretting gets started. Oxide debris can be seen at 100 cycles in Figure 23.7a. The worn surface is completely covered with an oxide (or a hydroxide) layer after 500 and 1000 cycles as shown in Figure 23.7b,c. The BSE images show a clear difference in contrast after 2000 cycles, indicating the formation of a tribochemical layer with a chemistry differing from that of the base Mg (Fig. 23.7d). With the increase in fretting cycles (as in Fig. 23.7e,f), the smearing and removal of the tribochemical layer is clearly visible. Additionally, the deformation of the underlying Mg surface along with some surface fatigue cracks can also be noticed in Figure 23.7f. It is interesting to note that for metallic materials, surface fatigue occurs simultaneously with fretting beyond a certain threshold number of cycles.6,11,12 Closer observation of Figure 23.7f shows that fatigue crack growth dominates wear at the highest number of cycles (5000 cycles) under dynamic fretting conditions. It is apparent from the preceding
23.4 Fretting-Duration-Dependent Tribological Properties
100 mm
0.1K
100 mm
0.5K
100 mm
(b)
100 mm
5K (e)
1K (c)
2K (d)
(a)
100 mm
â•… 369
10 mm
5K (f )
Figure 23.7â•… Fretted surfaces of pure Mg flats after different fretting cycles. Arrow marks indicate fretting direction. Fretting conditions: load, 8â•›N; frequency, 8â•›Hz; stroke length, 100â•›µm; counterbody, bearing steel.1
observations that, under the experimental fretting conditions studied, tribochemical layer formation and its subsequent removal continue to occur with increasing number of fretting cycles. The effect of SiCp addition to Mg matrix on wear can be observed during comparison of worn surfaces of base Mg with those of Mg–SiCp (Fig. 23.7 vs. Figs. 23.8–23.9). More extensive oxide film formation and its spreading are observed at early stages of fretting, at 200 and 500 cycles (Fig. 23.8a,b). During experiments, a soft tribochemical layer is formed and gets smeared as fretting continues. The chemistry of the tribolayer is analyzed with the help of Raman spectroscopy, as mentioned earlier.
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100 mm
0.2K
100 mm
100 mm
0.5K
10 mm
(b)
10 mm
2K (e)
1K (c)
2K (d)
(a)
100 mm
5K (f )
Figure 23.8â•… Worn surfaces of Mg–SiCp (9.8â•›wtâ•›%) composite after (a) 200, (b) 500, (c) 1000, (d,e) 2000, and (f) 5000 cycles. Arrows indicate the fretting direction. Fretting conditions: load, 8â•›N; frequency, 8â•›Hz; stroke length, 100â•›µm; counterbody, bearing steel.1
The effect of the addition of more of SiCp on fretting wear was critically investigated for Mg–SiCp (26.3â•›wtâ•›%) composite. Despite having similarities with base Mg, subtle differences can be observed in the wear behavior of these composites (Fig. 23.7 vs. Fig. 23.9). A dense oxide layer is observed on the worn surface of Mg–SiCp (26.3â•›wtâ•›%) at a lower number of cycles due to the presence of the increased amount of SiCp and also due to transferred oxide debris from the steel ball (Fig. 23.9a,b). The distinct feature observed on the worn surfaces is the smearing of brighter contrast tribolayer at larger fretting duration (Fig. 23.9c,d). Selective fretting of Mg–SiCp (26.3â•›wtâ•›%) composite at lower number of cycles (i.e., 5–80) shows certain interesting features. Figure 23.10a shows the formation of wear debris (oxidized) at five cycles itself. After 50 cycles, these small oxide wear
â•… 371
23.5 TRIBOCHEMICAL WEAR OF MAGNESIUM–SILICON CARBIDE
Mg Si Fe
100 mm
0.5K
50 mm (c)
(a)
100 mm
1K (b)
2K
100 mm
5K (d)
Figure 23.9â•… Worn surfaces of Mg–SiCp (26.3â•›wtâ•›%) composite after (a) 500, (b) 1000, (c) 2000, and (d) 5000 cycles. The EDX analysis on the tribolayer is shown in (c). Arrows indicate the fretting direction. Fretting conditions: load, 8â•›N; frequency, 8â•›Hz; stroke length, 100â•›µm; counterbody, bearing steel.1
debris particles are dispersed around the periphery of the wear scar (Fig. 23.10b). However, the formation of a thicker tribolayer is clearly observed in Figure 23.10c. At 100 cycles, this layer is spread over the worn surface (see Fig. 23.10d).
23.5 TRIBOCHEMICAL WEAR OF MAGNESIUM– SILICON CARBIDE PARTICULATE-REINFORCED COMPOSITES In the early stages of fretting experiments, Mg, being a soft metal (compared with bearing steel ball), undergoes extensive plastic deformation under repeated rubbing action, resulting in fine debris generation at the contact interface. As the activation energy for the oxidation of Mg is very low (ΔGâ•›≈â•›−599â•›kJ/mol at 25°C13), the debris readily oxidizes (due to frictional heating) and abrades the steel ball. It is well known that under fretting conditions, oxidation plays a major role in causing the wear damage of metallic materials.14 The iron oxide debris generated from the steel ball gets transferred to the worn surface and combines with magnesia. However, it should be noted that the material removal from the steel ball is less, as the bearing steel is much harder than Mg. The fine grooves formed on the worn surfaces of Mg indicate the occurrence of abrasive wear mechanism at very early stages of fretting. The
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100 mm
5k
100 mm (c)
(a)
10 mm
80k
50k
100 mm
(b)
0.1K (d)
Figure 23.10â•… Typical worn surfaces of Mg–SiCp (26.3â•›wtâ•›%) composites after fretting at lower number of cycles: (a) 5, (b), 50, (c) 80, and (d) 100 cycles. Arrows indicate the fretting direction. Fretting conditions: load, 8â•›N; frequency, 8â•›Hz; stroke length, 100â•›µm; counterbody, bearing steel.1
deformation of asperities and oxide formation at contacting surfaces cause a sharp increase in the COF during the running-in period (i.e., within 1500 cycles). Beyond 1500 cycles, the COF values decrease irrespective of load and attain a steady-state value. Careful observation shows the smearing of oxidation product on the worn surface (especially at higher loads). This indicates the reaction of oxide wear debris and formation of a soft reaction product at the contacting interface. The MgO formed during the early stages of fretting undergoes tribochemical reaction with moisture, resulting in the formation of a soft magnesium hydroxide tribolayer according to the following reactions:
2 Mg + O2 → 2 MgO; MgO + H 2 O → Mg (OH)2.
(23.1)
The hydrated magnesia, being viscous, easily spreads over the surface. Consequently, the friction and wear rate decrease significantly. The oxidative wear undergoes gradual transformation to tribochemical wear with the formation of hydrated magnesia [Mg(OH)2]. A similar reaction was reported in the adhesive wear of self-mated alumina under humid conditions.15 The formation of tribochemical film on worn alumina was observed at relative humidity (RH) levels of more than 10%. Below 10% RH, no film formation was observed, resulting in a large wear rate.
â•… 373
23.5 TRIBOCHEMICAL WEAR OF MAGNESIUM–SILICON CARBIDE
In the case of composites, the generation of Mg debris becomes severe, owing to the abrasive action of SiC particulates. However, the peak COF values for Mg and its composites (at any given load) are different. COF values obtained with Mg–SiCp (9.8â•›wtâ•›%) composites are higher than that of pure Mg, but this trend is not maintained after incorporation of higher volume fraction of SiCp. The peak COF values decrease to almost half (COF∼0.3) in Mg–SiCp containing 26.3â•›wtâ•›% SiCp, compared with the composites with 9.8â•›wtâ•›% reinforcement (COF ∼0.7). As the fretting progresses, COF values decrease and reach a steady-state value in all the composites. However, the steady-state values in the case of 9.8â•›wtâ•›% SiCP composites are somewhat scattered between 0.26 and 0.42 depending on the load, whereas in the case of 26.3â•›wtâ•›% composites the steady-state values vary between 0.23 and 0.31. This observation shows that the volume fraction of reinforcement has a large contribution to the fluctuations of the COF values. A detailed analysis was performed on the worn surfaces, and the chemistry of the tribolayer provides interesting information about the differences. The scars formed on composite materials show features that are almost similar to those seen on Mg metal. The only major difference is that the abrasive wear occurs with more severity due to SiC particulates. However, the reaction product formed in the composites appears brighter than that in pure Mg, indicating that these films have different chemistry from that of pure Mg samples. Raman spectroscopy results indicate that the triboproducts contain a considerable amount of SiO2 (peak at 263â•›cm−1). Second, the needle-like particles formed on the worn surfaces (Fig. 23.4) have similarity to the debris formed on the worn surfaces of SiC or Si3N4-based ceramic materials under humid conditions. The formation of needle-like particles has been detected in both oxides and nonoxides.16,17 Considering the influence of band gap energy (Eg) on the solubility of different covalent oxides in water, Fischer and Mullins observed that SiO2 (Eg ∼8.5â•›eV) has a higher tendency for tribochemical reaction with water18 compared with MgO (Eg ∼7.5â•›eV). Under humid conditions (≥30% RH, which is similar to our experiments), along with Mg(OH)2, hydrated silica forms by the following reactions:
2SiC + 3O2 → 2 SiO2 + 2 CO↑; SiO2 + 2 H 2 O → Si (OH )4.
(23.2)
The hydrated silica then reacts with Mg(OH)2 and forms DHMS, as evidenced by the Raman spectroscopy:
Mg (OH)2 + Si (OH)4 → MgSiO3 ⋅ 3H 2 O.
(23.3)
The DHMS is probably hydrated clinoenstatite.10 Although the Gibbs free energy for the formation of forsterite (Mg2SiO4) is lower than enstatite (MgSiO3),13 enstatite has higher stability than forsterite in the presence of silica. Consequently, enstatite is formed as the oxidation product of SiC formed during our fretting experiments. The iron transferred from the steel ball oxidizes and gets uniformly distributed in the viscous magnesium silicate layers and smears on the surface. The low value of COF in Mg–SiCp (26.3â•›wtâ•›%) composites is attributed to the formation of
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In Pure Mg Stage 1
Stage 2
Ball FeO FeO
Stage 3 Accumulated reaction products
Tribochemical reaction product MgO
Fe
Wear scar
Wear scar
Wear scar Wear grooves
• 2Mg + O2 → 2MgO • Transfer of Fe to tribosurface and oxide formation 4Fe + 3O2 → 2Fe2O3 • Oxidative wear
• Beginning of tribochemical reaction MgO + H2O → Mg(OH)2
Fatigue
• Formation and spreading of hydrated MgO on the tribosurface • Surface fatigue
In Mg–SiCp composites Stage 1 Fractured SiCp
Stage 2 SiCp
Ball
Stage 3 Mg(OH)2
Underneath SiCp
Fe FeO Wear scar
FeO
Wear scar
Wear scar Hydrated magnesium silicate tribolayer with Fe
• 2Mg + O2 → 2MgO • Fracture of SiC particulate • Transfer of Fe and Fe2O3 formation
• Tribochemical reactions MgO + H2O → Mg(OH)2 2SiC + 3O2 → 2SiO2 + 2CO SiO2 + 2H2O → Si(OH)4
• Formation of DHMS Fe2O3 + Si(OH)4+ Mg(OH)2 → MgSiO3·3H2O with Fe2O3 • Continuous smearing of triboproducts on the worn surface to form a tribolayer
Figure 23.11â•… Schematic of wear mechanisms operating during fretting of pure Mg and Mg–SiCp composites with steel.2
a magnesium silicate layer. These observations are in agreement with the adhesive results reported by Lim et al.19 In their results, the initial decrease in wear rate is due to the formation of the tribochemical layer. The mechanisms operating at different stages of fretting are schematically shown in Figure 23.11. To maintain clarity, the operating mechanisms are shown to be occurring in different stages. However, in practice, there is an overlap of the operating wear mechanisms (with increase in either load or number of cycles). Two aspects of the tribochemical wear of Mg–SiC have emerged. One is loaddependent behavior; the other is fretting-duration-dependent behavior. The smearing of the tribolayer is particularly affected by two factors: (1) volume fraction of SiC particulates and (2) load. When the volume fraction of SiC particulates is high, the formation of tribochemical layers sets in at early stages of fretting wear and the
23.6 Concluding Remarks
â•… 375
contact surfaces experience low COF from a very early stage. In 26.3â•›wtâ•›% SiC particulate composite, there is a very small difference in COF values during the running-in period and at steady state. On the contrary, similar steady-state values were obtained after 2000 cycles in the presence of small volume fraction of SiC particulates (9.8â•›wtâ•›%). Further, the steady-state COF differs largely depending on the load. Furthermore, the load also plays an important role. At lower loads, there is very small difference between peak values in running-in period and steady-state COF values. With the load increasing from 2 to 10â•›N, the difference between the two COF values becomes larger and the time (or number of cycles) taken to reach steady-state value is much longer. This suggests that, at higher loads, abrasive wear remains the dominant mechanism for quite a long time. From the topographical observation of worn surfaces using SEM, it is evident that the tribochemical layer smears fully and more uniformly under high load (>6â•›N). The tribolayer formation and its smearing over the worn surface are affected by the volume fraction and the number of cycles. At high volume fraction of SiCp (26.3â•›wtâ•›%), a dense and viscous tribochemical layer forms, which is responsible for decreased wear and friction without any fluctuations. On the contrary, the composite with lower SiCp (9.8â•›wtâ•›%) content experiences fluctuations in friction and wear. There exist a threshold number of cycles for the formation of a soft tribolayer and the achievement of steady-state friction condition. There is a small difference in steady-state COF value for Mg–SiCp (9.8â•›wtâ•›%) and Mg–SiCp (26.3â•›wtâ•›%) composites (0.05) compared with that between base Mg and composites (around 0.16). The implication of this work has significance in real-life applications. In Mg–SiCp composites (especially with higher volume fractions of SiC), the COF undergoes reduction (hence wear rate) with the formation of a viscous tribochemical layer. Conversely in other composites (e.g., Cu–Al2O3 or Al–Al2O3 composites), the formation of nonhygroscopic oxides leads to severe abrasive action, finally resulting in rapid failure of components due to high wear loss. Against the backdrop of the preceding considerations and based on the experimental results obtained under the selected fretting conditions, Mg–(26.3%)SiCp composites are considered to be a better choice for a tribosystem with the potential to replace Al–SiCp composites.
23.6 CONCLUDING REMARKS During the early stage of fretting, abrasive wear is dominant due to oxide formation. In composites, more abrasion occurs due to the presence of SiC particulates. After approximately 1500 fretting cycles, the wear mechanism transforms from abrasive wear to tribochemical wear with the formation of hydrous oxides. In pure Mg, MgO reacts with moisture, resulting in the formation of hydrous magnesium oxide, Mg(OH)2. In the case of composites, the tribochemical reaction of Mg(OH)2 and hydrated silica, Si(OH)4, results in the formation of DHMS, which is identified as clinoenstatite (MgSiO3·3H2O). The formation of hydrous magnesium silicate significantly decreases the COF values. The friction and wear of composites is observed to decrease only at higher reinforcement content (26.3â•›wtâ•›%). At lower SiCP content (9.8â•›wtâ•›%), the properties of base Mg matrix dominate and, hence, the beneficial role
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of reinforcement addition is not realized. Furthermore, the addition of larger volume fraction of SiCp beyond 9.8â•›wtâ•›% considerably improves the wear resistance under the investigated fretting conditions. During initial stages of fretting, oxidation is dominant in base Mg, while formation of a tribochemical layer is responsible for low friction and wear beyond 2000 cycles. At longer fretting duration, the underlying surface experiences fretting fatigue in base Mg, while the tribochemical layer (DHMS, i.e., MgSiO3·3H2O) provides protection to the surface in the case of composites. The experimental results summarized in this chapter clearly indicate that the amount of SiCp addition should be optimized in order to obtain better tribological properties.
REFERENCES ╇ 1â•… B. V. Manoj Kumar and B. Basu. Evolution in friction and wear of fretting wear of Mg-SiCp composites: Influence of fretting duration. J. Mater. Res. 20(4) (2005), 801–812. ╇ 2â•… B. V. Manoj Kumar, B. Basu, V. S. R. Murthy, and M. Gupta. The role of tribochemistry on fretting wear of Mg-SiC particulate composites. Composites Part A 36 (2005), 13–23. ╇ 3â•… M. K. K. Oo, P. S. Ling, and M. Gupta. Characterization of Mg-based composites synthesized using a novel mechanical disintegration and deposition technique. Metall. Mater. Trans. 31A (2000), 1873–1881. ╇ 4â•… T. E. Fischer and H. Tomizawa. Interaction of tribochemistry and microfracture in the friction and wear of silicon nitride. Wear 105 (1985), 29–45. ╇ 5â•… B. Basu, T. Venkateswaran, and D. Sarkar. Pressureless sintering and Tribological properties of WC-ZrO2 composites. J. Eur. Ceram. Soc. 25 (2005), 1603–1610. ╇ 6â•… R. C. Bill. Review of factors that influence fretting wear. Am. Soc. Test. Mat. (1982), 164–182. ╇ 7â•… P. Anderson, J. Juhanko, A.-P. NIkkila, and P. Lintula. Influence of topography on the running-in of water-lubricated silicon carbide journal bearings. Wear 201 (1996), 1–9. ╇ 8â•… M. G. Gee. The formation of aluminium hydroxide in the sliding wear of alumina. Wear 153 (1992), 201–227. ╇ 9â•… H. Tomizawa and T. E. Fischer. Friction and wear of silicon nitride and silicon carbide in water: Hydrodynamic lubrication at low sliding speed obtained by tribochemical wear. ASLE Trans. 30(1) (1986), 41–46. 10â•… A. M. Hofmeister, H. Cynn, P. C. Burnley, and C. Meade. Vibrational spectra of dense, hydrous magnesium silicates at high pressure: Importance of the hydrogen bond angle. Am. Miner. 84 (1999), 454–464. 11â•… R. C. Bill. Fretting wear and fretting fatigue—how are they related? Trans. ASME 105 (1983), 230–238. 12â•… D. W. Hoeppner and G. L. Goss. A fretting-fatigue damage threshold concept. Wear 27 (1974), 61–70. 13â•… M. W. Chase. NIST-JANAF thermochemical tables. J. Phys. Chem. Ref. Data Monogr. 9 (1998), 1–1951. 14â•… R. B. Waterhouse. Fretting wear. Wear 100 (1984), 107–118. 15â•… A. J. Perez-Unzueta, J. H. Beynon, and M. G. Gee. Effects of surrounding atmosphere on the wear of sintered alumina. Wear 146 (1991), 179–196. 16â•… M. Chen, K. Kato, and K. Adachi. The difference in running-in period and friction coefficient between self-mated Si3N4 and SiC under water lubrication. Tribo. Lett. 11 (2001), 23–28. 17â•… J. Xu and K. Kato. Formation of tribochemical layer of ceramics sliding in water and its role for low friction. Wear 245 (2000), 61–75. 18â•… T. E. Fischer and W. M. Mullins. Relation between Surface Chemistry and Tribology of Ceramics. Friction and Wear of Ceramics. S. Jahanmir (ed.). Marcel Dekker, New York, 51–60. 19â•… C. Y. H. Lim, S. C. Lim, and M. Gupta. Wear behaviour of SiCp-reinforced magnesium metal matrix composites. Wear 255 (2003), 629–637.
CHAPTER
24
CASE STUDY: TITANIUM CARBONITRIDE–NICKELBASED CERMETS Currently, TiCN-based cermets are being developed as alternative materials to WC–Co hardmetals for cutting tool inserts due to their superior properties, such as high hardness with improved toughness, wear resistance with chemical stability, and thermal shock resistance. Several secondary carbides, such as WC, NbC, HfC, and TaC, are added to TiCN–Ni cermets to enhance high-temperature mechanical and wear-resistance properties. The results discussed in this chapter will attempt to answer a number of relevant issues: (1) How does the addition of various secondary carbides (WC, NbC, TaC, HfC) influence friction and wear behavior and whether such influence is dependent on load? (2) How does the secondary carbide addition affect the stability of the tribochemical layer? In a separate section in this chapter (Section 24.6), the influence of WC addition (5–25â•›wtâ•›%) on the sliding wear behavior of TiCN-20â•›wtâ•›% Ni cermets is analyzed. Efforts have been made to discuss the possible reaction pathways to explain the formation of a tribochemical layer. Another important aspect of this chapter is to discuss the high-temperature wear properties of TiCN–Ni cermets containing different secondary carbides (WC, NbC, TaC, or HfC).
24.1 INTRODUCTION Titanium carbonitride (TiCN)-based cermets, in view of their lower density and good specific properties (specific hardness/modulus), have been in development since the mid-1990s.1 The typical core–rim structure produced as a result of a dissolutionand-reprecipitation process greatly influences the material properties of cermets.2 The influence of secondary carbide addition on core–rim structure has also been investigated from a microstructural point of view.3–8 While the addition of WC is indispensable in many cases to achieve densification and satisfactory fracture toughness, NbC or TaC addition enhances the interrupted cutting performance by retaining
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
377
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CHAPTER 24â•… Case Study: Titanium Carbonitride–Nickel-Based Cermets
hardness at high temperature and increasing thermal shock resistance at elevated temperatures.3,4 The addition of HfC reduces strain misfit among different microstructural constituents.5 From the compositional design aspect of TiCN-based cermets, the influence of various secondary carbide additions (WC, NbC, etc.) on sliding wear properties needs to be investigated. Different materials as well as operating parameters influence wear mechanisms such as adhesion, surface fatigue, tribochemical reactions, and abrasion, among others.9 While the sliding of metallic materials is dominated by plastic deformation and oxidative wear, the wear of ceramics and their composites is influenced by brittle fracture and tribochemical wear.9–11 However, the simultaneous occurrence of several mechanisms often complicates the wear of different materials and this requires systematic tribological experiments and analyses. Attempts have been made to assess TiC-based composites for high-temperature applications.12–14 A transition in wear mechanism with load was observed for TiCbased composites. In another study, formation of transfer film and deformation were reported during sliding of TiC–NiMo cermets against steel.12 Peterson and Lee15 suggested that the characteristic of transferred film could significantly affect the unlubricated frictional behavior of metal–ceramic and metal–cermet tribocouples. Engqvist et al.16 also recorded the formation of a thick tribofilm on the worn surface of WC–Co hardmetals during self-mated sliding conditions. It was experimentally observed that WC debris particles rapidly cemented together with the Co binder phase to form a mechanically mixed tribolayer at a sliding contact.17 Zhao et al.19 observed that the cooling effects of lubricants could enhance the friction and wear of TiCN ceramics against steel. The beneficial effects of an oil lubricant under boundary lubrication conditions were primarily due to the formation of carbon films on the rubbing surfaces. Kameo et al.,20 in their sliding wear study on cermets and ceramic composites, observed that wear of the steel counterbody is minimal against cermets, compared with that against ceramic composites. In particular, the wear of a steel counterbody was much higher in the case of both ZrO2 and mullite, when compared with competing cermet material. A number of researchers investigated the various aspects of the wear of TiCN-based cermets.21–25 For example, Jeon et al.21 investigated the effect of TiCN particle size and WC content on the wear loss of TiCN–xWC–Ni (x, 5–25) cermets, slid against a Si3N4 ball at 600°C. For TiCN particle sizes of 0.7–0.95â•›µm, the addition of WC (up to 25%) could not influence the severity of wear, as wear depth varied around 5.5–5.8â•›µm, independent of WC content. In contrast, a linear increase in wear depth with WC content was measured, when coarser TiCN particles (3–5â•›µm) were used to develop such cermets. Sarkar et al.22 reported that the fretting wear resistance of the TiCN–WC–Ni cermets against steel increased when the amount of WC added was above 5â•›wtâ•›%. The relevance of friction and wear can be better understood in the context of the machining applications23,24,26,27. During machining, the tool–chip contact experiences both normal and shear stresses and high temperatures.26,28 Under such severe tribological situations, simultaneous occurrence of different wear mechanisms is possible, including (1) adhesion and material transfer due to cold welding and breaking, (2) abrasion due to hard protuberances on either or both tool and workpiece, and (3) tribochemical wear as a result of diffusion of chemical species and rubbing,
24.2 Materials and Experiments
â•… 379
among others. Such mechanisms could be well understood by maintaining intimate surface contact during cutting operation.28 Xingzhong et al.29 investigated high-temperature wear properties of TiC–Ni against high-speed steel. The formation of an oxide film was found to be the major mechanism at 600°C. The addition of Mo increased the lubricity of the tribochemical layer with the formation of MoO3, whereas the formation of WO3 degraded the friction and wear. Meng et al.30 also observed tribo-oxidation at high temperature (600°C) for TiCN–Ni–Mo, as well as TiCN–Al2O3–Ni–Mo cermets, when slid against Si3N4 in a ball-on-disk configuration. Concerning TiCN particle size, cermets made of ultrafine TiCN powders (0.4â•›µm) exhibited an enhanced sliding wear resistance of TiCN–WC–Ni cermets against Si3N4 at 600°C.31 Degnan et al.13 reported mild wear at 500°C for TiC-reinforced steel composites, and this was attributed to the rapid formation and stability of oxide layers and compacted glazes of crystalline oxide particles. However, Akhtar and Guo32 observed severe wear due to extensive microplowing and rapid material removal at high loads and low TiC content against high-speed-steel counterbody. In assessing high-temperature wear behavior, oxidation properties are important.33 While the addition of NbC or TaC could lead to an increase in the oxidation resistance, the addition of WC is reported to degrade oxidation properties due to the volatilization of WO3.34 Against this backdrop, the tribological properties of TiCN–(20â•›wtâ•›%)Ni– (10â•›wtâ•›%)X (where X is WC, NbC, TaC, HfC) cermets are being investigated under sliding–fretting at ambient and high-temperature conditions and results are reported elsewhere.35–39 This chapter reports some important aspects of the sliding wear and friction properties with a major focus on understanding various wear mechanisms.
24.2 MATERIALS AND EXPERIMENTS Samples of TiCN-based cermets were sintered at 1510°C for 1â•›hour in vacuum using Ti(C0.7N0.3) powders with various 10â•›wâ•›t% secondary carbides (WC, NbC, TaC, and HfC) and Ni as the binder. Some representative scanning electron microscope (SEM; in backscattered electron [BSE] mode) images of TiCN–20Ni–10WC cermet in Figure 24.1 reveals the characteristic core–rim structure. Typical SEM images revealing polished microstructures of unworn TiCN-based cermets are shown in Figure 24.1a–c. A typical energy-dispersive x-ray spectrometry (EDS) line scan profile of TiCN–20Ni–25WC showing compositional variation is presented in Figure 24.1d. Such a qualitative analysis with respect to Ti, W, and Ni (see inset in Fig. 24.1d) confirms the predominant presence of Ti in the core (black) region and Ti, W, and Ni in the rim (gray) region. The binder (white) region consists mainly of Ni with dissolved Ti and W. A schematic of the phase assemblage is also provided in Figure 24.1e. The physical and mechanical properties are summarized in Table 24.1. To obtain hardness and toughness, the length of the radial cracks propagating from the Vickers indent edges was measured and Shetty’s formula40 was used to estimate the fracture toughness (KIc). From Table 24.1, it is clear that hardness varies over a wide range, with the lowest hardness (∼10â•›GPa) observed for TiCN–20Ni and
380â•…
CHAPTER 24â•… Case Study: Titanium Carbonitride–Nickel-Based Cermets
20 mm
20 mm
(b)
(a)
Ti
Ni 5 mm
20 mm 0
(c)
W 1
2
3
4
µm
(d)
Outer rim (Ti, M) CN
Inner rim (Ti, M) CN
Ni-rich binder
Core TiCN (e)
Figure 24.1â•… Representative SEM (BSE) images of unworn polished surfaces of TiCN–Nibased cermets: (a) TiCN–(20â•›wtâ•›%)Ni–(5â•›wtâ•›%)WC; (b) TiCN–(20â•›wtâ•›%)Ni–(15â•›wtâ•›%)WC; (c) TiCN–(20â•›wtâ•›%)Ni–(25â•›wtâ•›%)WC. A representative line scan–EDS analysis of a selected length on the microstructure of TiCN–(20â•›wtâ•›%)Ni–(25â•›wtâ•›%)WC cermet (inset) is shown in (d). A schematic of the microstructural phase assemblage is also shown (e).35,37
TiCN–20Ni–10NbC/HfC and the highest (∼12â•›GPa) for TiCN–20Ni–10WC/TaC. Also, the highest KIc (∼15â•›MPaâ•›m1/2) was recorded for TiCN–20Ni cermet, whereas TiCN–20Ni–10WC/NbC exhibits a KIc of ∼14â•›MPaâ•›am1/2. The literature reports indicated that a much higher toughness (18.5â•›MPaâ•›m1/2) could be measured with TiC– NiMo cermets.41 The operating parameters for room-temperature fretting experiments performed on cermets against steel balls were: varying load (P) of 2, 6, and 10â•›N at 4-Hz oscillating frequency and 100-µm linear stroke for a duration of 100,000
â•… 381
24.3 Energy Dissipation and Abrasion at Low Load
TABLE 24.1â•… Summary of the Density and Mechanical Properties of Various TiCN-Based Cermets35–37
Cermet composition (numerals in wtâ•›%)
Density (g/cm3)
HV30 (GPa)
Fracture toughness MPaâ•›m1/2
Cermets with different type of secondary carbides Ti(CN)–20Ni Ti(CN)–20Ni–10WC Ti(CN)–20Ni–10NbC Ti(CN)–20Ni–10TaC Ti(CN)–20Ni–10HfC
5.33 5.87 5.52 5.48 5.33
9.9╯±â•¯0.5 11.8╯±â•¯0.6 10.9╯±â•¯0.5 12.5╯±â•¯0.6 9.3╯±â•¯0.8
15.5╯±â•¯0.8 14.8╯±â•¯0.7 11.2╯±â•¯0.5 14.0╯±â•¯0.7 11.7╯±â•¯0.6
Cermets with varying amount of WC addition Ti(CN)–20Ni–5WC Ti(CN)–20Ni–15WC Ti(CN)–20Ni–20WC Ti(CN)–20Ni–25WC
5.58 5.99 5.89 5.76
11.2 10.6 11.7 11.5
13.9 15.2 13.2 18.0
cycles. Correspondingly, the initial Hertzian contact pressure varied from 710â•›MPa at 2â•›N to 1020â•›MPa at 6â•›N and 1210â•›MPa at 10â•›N. Two sets of materials were used for room-temperature sliding tests on a pinon-flat tribometer. In the first set of experiments, TiCN–Ni cermets with 10â•›wtâ•›% of various secondary carbides were evaluated; the second set of sliding tests were performed on TiCN–Ni cermets with varying amounts of WC (up to 25â•›wtâ•›%). Each cermet was slid against a steel counterbody at loads of 5, 20, and 50â•›N, corresponding to initial Hertzian contact pressure of 640, 1020, and 1390â•›MPa, respectively. The tests were carried out at a constant oscillation frequency of 5â•›Hz and a stroke length of 2.4â•›mm. The total sliding distance in each test was 100.8â•›m (21,000 cycles). Such a combination of operating parameters covers broad spectrum of stress conditions in understanding the friction and wear response. High-temperature wear results were obtained from low-amplitude oscillations at 550°C. Studies were conducted on samples of TiCN–20Ni–10XC (where X is W, Nb, Ta, Hf) cermets using a ball-on-flat tribometer; the use of such a test setup is reported elsewhere.42 The operating parameters were a 20-N load at 6-Hz oscillating frequency and 0.1-mm linear stroke for a duration of 4â•›hours and 37 minutes. The corresponding linear velocity and sliding distance were 1.2â•›mm/s and 19.94â•›m, respectively.
24.3 ENERGY DISSIPATION AND ABRASION AT LOW LOAD A typical three-dimensional fretting log, revealing the variation of frictional force versus displacement stroke length, is provided in Figure 24.2a. The shape of the hysteresis loops remains almost a parallelogram and this indicates the gross slip
Tangential Force (N)
382â•…
CHAPTER 24â•… Case Study: Titanium Carbonitride–Nickel-Based Cermets
8 7 6 5 4 3 2 1 0 –1 –2 –3 –4 –5
0 25,000 Num 50,000 ber of C 75,000 ycle 100,000 s
60 40 m) 0 t (m n e –20 –40 lacem p –60 s Di 20
TiCN–20Ni–10TaC @10 N (a) 5 2N 6N 10 N
4
Tangential Force (N)
3 2 1 0 –1 –2 –3 –4 –5
–60
–40
–20 0 20 Displacement (mm) (b)
40
60
Figure 24.2â•… (a) Representative 3D fretting log for TiCN–20Ni–10TaC cermet/steel tribocouple at 10-N load and (b) variation of tangential frictional force with displacement at 6000 cycles for TiCN–20Ni–10TaC cermet/steel tribocouple at different loads.35
regime. Figure 24.2b plots the variation of tangential force with load for a TiCN– 20Ni–10TaC cermet/steel couple at 6000 cycles. The area of the loop increased with increase in load from 2 to 10â•›N and this indicates a corresponding increase in dissipated energy. The material removal during sliding can be analyzed on the basis of the frictional energy dissipation at the sliding contact.43,44 For the ball-on-flat configuration, the cumulative dissipated energy (Ed) can be determined from the area of the tangential friction (F) versus displacement (d) loop:
â•… 383
24.3 Energy Dissipation and Abrasion at Low Load
5.0
TiCN–20Ni TiCN–20Ni–10WC TiCN–20Ni–10NbC TiCN–20Ni–10TaC TiCN–20Ni–10HfC
Wear volume (¥ 10–4 mm3)
4.5 4.0 3.5 3.0 2.5 2.0 1.5 1.0
R2 = 0.99 Slope: 3.9 × 10–6 mm3/J
0.5 0.0
0
10
20
30
40
50
60
70
80
90
100
Dissipated energy (J)
Figure 24.3â•… Linear relationship between wear volume and dissipated energy during fretting of investigated cermets after fretting against steel. Slope of the linear plot was approximated with the least squares fit.35
Ed =
∑ Fd.
(24.1)
The dissipated energy can be calculated using the following relationship: Ed = µWVt ,
(24.2)
where μ is the average coefficient of friction (COF), W is the normal load, v is the fretting velocity, and t is the total testing duration.43,45 The variation in wear volume against dissipated energy for TiCN–Ni-based cermets is plotted in Figure 24.3. The fact that the wear volume increases linearly with dissipated energy indicates that the compositional variation of cermet does not change the major material removal mechanism. From the slope of the linear plot, the wear rate is found to be 3.9╯×╯10−6â•›mm3/J, which lies in the same range as profilometric wear rate calculations (see Table 24.2). The tribological response is expected to be the combined response of binder metal and ceramic particulates. It is known that the abrasive wear in brittle solids is similar to the damage, resulting from sliding of a sharp indenter on a soft surface. When sharp asperity or indenter slides over a softer solid, the lateral cracks grow upward from the subsurface, removing the material as platelets, and the wear volume per unit sliding distance (V) can be considered to be given by the following relation:46
V=
αN ( E / H )W 9 / 8 , 1/ 2 K IC H 5/8
(24.3)
384
1.33
1.34
1.35
1.23
1.42
5.31
5.05
5.75
5.66
Wear volume (×10−4) mm3
5.88
Max. wear depth (µm)
2â•›N
All measurements within ±5% scatter in values.
TiCN– 20Ni TiCN– 20Ni– 10WC TiCN– 20Ni– 10NbC TiCN– 20Ni– 10TaC TiCN– 20Ni– 10HfC
Material
3.55
3.07
3.38
3.31
3.33
Wear rate (×10−6â•›mm3/Nâ•›m)
6.57
6.45
6.36
6.58
6.45
Max. wear depth (µm)
3.02
2.67
2.80
2.63
2.49
Wear volume (×10−4) mm3
6â•›N
2.52
2.18
2.33
2.19
2.08
Wear rate (×10−6â•›mm3/Nâ•›m)
8.28
7.40
8.19
7.06
7.14
Max. wear depth (µm)
3.99
3.32
3.92
3.26
3.49
Wear volume (×10−4) mm3
10â•›N
1.99
1.66
1.96
1.63
1.75
Wear rate (×10−6â•›mm3/Nâ•›m)
TABLE 24.2â•… Profilometry Measurements, Wear Volume, and Wear Rate of Investigated Cermets After Fretting at Different Loads Under Ambient Conditions35
â•… 385
24.3 Energy Dissipation and Abrasion at Low Load
5.0
Wear Volume (¥ 10–4mm3)
4.5 4.0 3.5 3.0 2.5 2.0 1.5
TiCN–20Ni TiCN–20Ni–10WC TiCN–20Ni–10NbC TiCN–20Ni–10TaC TiCN–20Ni–10HfC
1.0 0.5 0.0 0.0
0.2
0.4
0.6
0.8
1.0
1.2
Abrasion Parameter ([W9/8/(Kc1/2 ¥ H5/8)])
Figure 24.4â•… The variation of wear volume with the abrasion parameter for the investigated cermets.35
where W╯=╯normal contact force, KIC╯=╯fracture toughness, E╯=╯elastic modulus, H╯=╯hardness, N╯=╯total number of contacting asperities, and α╯=╯materialindependent constant. Since the ratio (E/H) is insensitive to a greater extent for different brittle solids,8,47 the wear volume should be proportionate to the abrasion term as: W 9/8 (24.4) Vα 1 / 2 5 / 8 . K IC H The experimentally measured wear volume against the abrasion parameter reveals a linear relationship (Fig. 24.4), indicating abrasion is the major wear mechanism. Nevertheless, it has been reported elsewhere that the wear of the TiCN–Nibased cermets under severe sliding conditions is primarily characterized by the transfer of iron oxide to the cermet surface with increase in load, velocity, or surface temperature. The influence of secondary carbides on the wear of TiCN–Ni cermets can be explained as follows. The fracture toughness decreased with the addition of HfC or NbC. The lower hardness and toughness of the TiCN–20Ni–10HfC/NbC cermets can lead to easy grain pullouts and increased amount of debris, which is in agreement with the wear data (Table 24.2). The influence of load on the wear behavior of TiCN–Ni-based cermets can be discussed. The abrasion is primarily caused by the hard oxides of different phases of cermets and steel. The possible reactions of oxides of various phases are explained elsewhere.36 With increase in load, the formation of hard oxide debris and the corresponding abrasion increased, and this leads to an increase in wear volume (Table 24.2). This corroborates well with Archard’s wear theory.48 The wear rate decreased at 10-N load with the domination of threebody abrasion.
386â•…
CHAPTER 24â•… Case Study: Titanium Carbonitride–Nickel-Based Cermets
24.4 INFLUENCE OF TYPE OF SECONDARY CARBIDES ON SLIDING WEAR OF TITANIUM CARBONITRIDE– NICKEL CERMETS The frictional behavior of TiCN–Ni-based cermets with different secondary carbides (WC, NbC, TaC, HfC) was investigated at varying load from 5 to 50â•›N. The average steady-state COF values are in the range 0.64–0.75 at 5-N load, while lower COF values from 0.46 to 0.52 can be recorded at 20- or 50-N load. Further, at a low load (5â•›N), TiCN–20Ni–10WC/steel exhibited steady-state COF of 0.64, whereas a higher COF of 0.75 was measured with TiCN–20Ni–10HfC cermet/steel. In general, the friction decreased with increase in load from 5 to 20â•›N and such a decrease in COF is independent of the presence of secondary carbides. The wear rates of cermets, as well as ball surfaces, as plotted in Figure 24.5a,b, reveal that HfC exhibited the highest wear volume (20╯×╯10−4â•›mm3) at 50â•›N, whereas a minimum wear volume (3╯×╯10−4â•›mm3) was measured with TaC-containing cermet at 5-N load. The mechanical properties of the cermet materials are expected to influence the wear behavior. We can observe from Table 24.1 and Figure 24.5a that at lower load (5â•›N), the wear volume has an inverse relationship with hardness, usually true for abrasive-wear-dominated conditions. That is, baseline TiCN–20Ni cermet and TiCN–20Ni–10HfC cermet, having lower hardness, experience severe wear,
(c)
TiCN–20Ni–10HfC
0.0
TiCN–20Ni–10TaC
5.0
TiCN–20Ni–10WC
10.0
TiCN–20Ni–10NbC
15.0
TiCN–20Ni–10HfC
TiCN–20Ni–10TaC
TiCN–20Ni–10NbC
0.0
TiCN–20Ni–10WC
10.0
(b)
25.0
5N 20 N 50 N
20.0 15.0 10.0 5.0 0.0
TiCN–20Ni–10HfC
20.0
20.0
TiCN–20Ni–10TaC
25.0
5N 20 N 50 N
30.0
TiCN–20Ni–10NbC
30.0
Specific wear rate of ball (¥ 10–5 mm3/N m)
(a)
40.0
TiCN–20Ni–10WC
TiCN–20Ni–10HfC
TiCN–20Ni–10TaC
0.0
Wear volume of steel ball (¥ 10–2 mm3)
TiCN–20Ni–10WC
TiCN–20Ni
5.0
TiCN–20Ni–10NbC
10.0
5N 20 N 50 N
50.0
TiCN–20Ni
15.0
60.0
TiCN–20Ni
20.0
Specific wear rate (¥ 10–7 mm3/N m)
5N 20 N 50 N
TiCN–20Ni
Wear volume (¥ 10–4 mm3)
25.0
(d)
Figure 24.5â•… The wear volume and wear rates of various cermets (a,b) and steel balls (c,d) with varying load. The sliding conditions include frequency, 5â•›Hz; stroke length, 2.4â•›mm; and duration, 21,000 cycles.36
â•… 387
24.4 INFLUENCE OF TYPE OF SECONDARY CARBIDES ON SLIDING WEAR
while TiCN–20Ni–10TaC cermet with higher hardness exhibits lower wear at 5-N load. With increase in load to 20 and 50â•›N, such influence is diminished primarily due to the formation of a potential tribolayer. The thicker tribolayer at higher load (50â•›N) protects the cermet surface, and thus the influence of hardness is minimal.
24.4.1â•…Wear Mechanisms As far as the worn surface topography is concerned, it is severely deformed at 15-N load with few occurrences of layer formation noted on TiCN–20Ni cermet (Fig. 24.6a). The removal of thin layers as fragments can be observed after sliding at a higher load of 20â•›N (Fig. 24.6b); with further increase in load to 50â•›N, the worn surface is covered by a thick tribolayer (Fig. 24.6c). The EDS analysis essentially reveals major peaks of Ti, Fe, and O, indicating that the tribolayer is rich in iron oxides and the formation of such layer is responsible for decreased wear as well as friction (see Fig. 24.3). The lubricity of the layer might have reduced COF (Table 24.2). Figure 24.7 provides representative SEM images of worn surface of TiCN–Ni cermet containing WC. Compared with TiCN–20Ni, the TiCN–20Ni–10WC cermet exhibits lower wear at 5â•›N (Fig. 24.7a). The increased tribolayer formation can be observed at 50-N load in Figure 24.7b. A similar wear behavior with mild tribolayer formation is observed at 5-N load for TiCN–Ni–NbC cermet (see Figure 24.8a). Interestingly, the debris entrapped in the sliding contact leads to severe abrasion at 20â•›N (shown in Figs. 24.8b and 24.8c). This corroborates well with the increased wear rate for NbC-containing cermet (see Fig. 24.5). SEM images of TiCN–20Ni– 10TaC cermet worn surface (Fig. 24.9) also show similar features. With increase in load to 50â•›N (see Fig. 24.9b), the formation and removal of tribolayer is observed. TiCN–20Ni–10HfC cermet worn surface, shown in Figure 24.10, reveals the characteristic grain pullout as well as the presence of severe abrasion grooves. X-ray mappings (not shown) indicate that Fe and O are present at the tribolayer, while Ti, Ni, and Hf are at the worn cermet surface, which is not covered with tribolayer.
24.5 TRIBOCHEMICAL WEAR OF TITANIUM CARBONITRIDE–BASED CERMETS First, the evolution of tribochemical reactions is discussed, followed by a discussion on the effect of microstructural features on the friction and wear behavior of cermets at varying load.
24.5.1â•… Evolution of Tribochemistry and Contact Temperature Due to higher hardness of cermets, the abrasion of steel ball by harder cermet asperities leads to the formation of iron debris:
xFe + ( y/ 2)O2 → Fe x O y.
(24.5)
388â•…
CHAPTER 24â•… Case Study: Titanium Carbonitride–Nickel-Based Cermets
10 mm
10 mm (b)
(a)
O Ti Fe
Ti Fe Ti
10 mm 2
4
(c)
Fe 6
8
keV
(d)
O Ti Fe
Fe Ti 2
Fe
Ti
4
6
8 keV
(e)
Figure 24.6â•… Secondary electron SEM images revealing the characteristic topographical features of worn TiCN–20Ni cermet after sliding against steel ball at loads (a) 5â•›N, (b) 20â•›N, and (c) 50â•›N. The EDS analysis of tribolayer at loads 5â•›N (shown with arrow) and 50â•›N are shown as (d) and (e), respectively. Sliding conditions include frequency, 5â•›Hz; stroke length, 2.4â•›mm; and duration, 21,000 cycles. Double-pointed arrows indicate sliding direction.36
Subsequent sliding in the presence of iron oxides causes abrasion of the cermet surface. When the deformed binder phase is smeared away during sliding, stressinduced cracking leads to pullout of ceramic grains. Consequently, various ceramic phases in cermets are subjected to oxidation, as is evident from EDS results presented in Figures 24.6b,e, 24.7c,d, 24.8d, and 24.9c Thus, the oxidation of rim solid solution (Ti,X)CN (where X is W, Ta, Nb, Hf) occurs according to the following reaction pathways:
WC + O2 → WO3 + CO,
(24.6)
24.5 Tribochemical Wear of Titanium Carbonitride–Based Cermets
10 mm
10 mm
(b)
(a) O Cr Ti Fe
Ti Fe Cr O Ni
Ti
Fe
Ni W W 2
â•… 389
Ti Cr Cr 4
6
Fe
Ti Ni W Ni Fe W W W 8 keV
Ti 2
(c)
CrCr
4
6
Fe Ni Ni 8
keV
(d)
Figure 24.7â•… SEM images showing the details of the worn surfaces of TiCN–20Ni–10WC cermet, after sliding against steel ball at loads (a) 5â•›N and (b) 50â•›N. The EDS analyses acquired from the indicated regions (bright spots in (a) and (b)) of the tribolayer at 5- and 50-N loads are shown in (c) and (d), respectively. Sliding conditions include frequency, 5â•›Hz; stroke length, 2.4â•›mm; and duration, 21,000 cycles. Double-pointed arrows indicate sliding direction.36
NbC + O2 → Nb2 O5 + CO, TaC + O2 → Ta 2 O5 + CO, HfC + O2 → HfO2 + CO.
(24.7) (24.8) (24.9)
Further, TiCN can be simplistically considered as a solid solution of TiC and TiN and, therefore, the oxidation of the TiCN can be described by their respective oxidation reactions: (24.10) TiC + (3 / 2)O2 → TiO2 + CO, TiN + O → TiO + ( 1 / 2 ) N . (24.11) 2 2 2 Therefore, sliding in the presence of this debris at the tribocontact will result in probable hard contacts between various oxides of cermet constituents and oxides of iron, resulting in high COF (see Table 24.3). However, with increase in load, the compaction of oxide wear debris results in tribolayer formation, which potentially protects the underlying material from further wear. In addition to the severe mechanical aspect, with the increase in amount and compaction of debris, the chemical aspect—namely, diffusion or material transfer—will be enhanced at higher loads due to temperature rise at the contact.
390â•…
CHAPTER 24â•… Case Study: Titanium Carbonitride–Nickel-Based Cermets
100 mm
10 mm
(b)
(a) Nb
Ni Fe Ti
10 mm
2
(c)
Ti
Nb
Ti
4
Fe FeNi Ni 6 8 keV
(d)
Figure 24.8â•… SEM images of TiCN–20Ni–10NbC worn surface, after sliding against steel ball at loads (a) 5â•›N and (b,c) 20â•›N. The EDS analysis of debris–tribolayer fragment [indicated with an arrow in (b)] at 20-N load is shown as (d). Sliding conditions include frequency, 5â•›Hz; stroke length, 2.4â•›mm; and duration, 21,000 cycles. Double-pointed arrows indicate sliding direction.36
Figure 24.11 reveals the formation of loose debris on TiCN–20Ni–25WC cermet during initial stages (at 3000 cycles in Fig. 24.11a), followed by their agglomeration with further sliding (at 21,000 cycles in Fig. 24.11c). The composition of the collected debris, analyzed using x-ray diffraction, indicates the formation of Fe9TiO15 and Fe2O3 (Fig. 24.11d). The formation of Fe9TiO15 can be described by the following tribochemical reaction:
FeTiO3 + 4 Fe 2 O3 → Fe 9 TiO15.
(24.12)
The following tribochemical reaction pathway can explain the formation of an iron titanate layer and can be expressed as:
FeO + TiO2 → FeTiO3. 49,50
(24.13)
According to existing literature reports, the FeO–Fe2O3–TiO2 system is characterized by three major solid solution series: the orthorhombic ferropseudobrookite–pseudobrookite (FeTi2O5–Fe2TiO5) series, the rhombohedral illmenite–hematite (FeTiO3–Fe2O3) series, and the spinel ulvospinel–magnetite (Fe2TiO4–Fe3O4) series. The illmenite–hematite series forms a solid solution and develops a miscibility gap below 800°C. At temperatures below 400°C, both Fe2O3
24.5 Tribochemical Wear of Titanium Carbonitride–Based Cermets
â•… 391
10 mm
10 mm
(b)
(a)
O Ti Fe
Fe Ti Ti 2
4
Fe 6
8 keV
(c)
Figure 24.9â•… The worn surface morphology of TiCN–20Ni–10TaC cermet after sliding against steel ball at loads (a) 5â•›N and (b) 50â•›N. EDS analysis of tribolayer at 50-N load is shown as (c). Sliding conditions include frequency, 5â•›Hz; stroke length, 2.4â•›mm; and duration, 21,000 cycles. Double-pointed arrows indicate sliding direction.36
10 mm
10 mm (a)
(b)
Figure 24.10â•… Typical SEM images of TiCN–20Ni–10HfC worn surface after sliding against steel ball at loads (a) 5â•›N and (b) 20â•›N. Sliding conditions include frequency: 5â•›Hz; stroke length, 2.4â•›mm; and duration, 21,000 cycles. Double-pointed arrows indicate sliding direction.36
392â•…
CHAPTER 24â•… Case Study: Titanium Carbonitride–Nickel-Based Cermets
O Cr Ti Fe
Fe
Ni W W
1 mm
2 (a)
Fe Ti Ni W Cr Ti Cr W Ni 4 6 8 keV (b)
Relative Intensity (a.u)
350
20 mm
300
# *
250 200 150
# #
100 50 0
(c)
* : Fe2O3 # # : Fe9TiO15 *
20
40 60 2q (degrees) (d)
80
100
Figure 24.11â•… Typical SEM images of debris formed after sliding TiCN–20Ni–25WC against steel ball at 50-N load for (a) 3000 and (c) 21,000 cycles. The EDS analysis (b) of debris after 3000 cycles indicates the presence of a small amount of W, whereas the x-ray diffraction (XRD) spectra [shown as (d)] of debris produced after 21,000 cycles reveals the crystalline phases of debris to be Fe2O3 and/or Fe9TiO15. The largely diffused background intensity was due to the fact that a plastic tape was used to fix wear debris particles for XRD investigation. Other sliding conditions include frequency, 5â•›Hz; and stroke length, 2.4â•›mm.36
and FeTiO3 can potentially form. The formation of Fe2O3 and FeTiO3 in the Fe–Ti–O system was confirmed by transmission electron microscope (TEM) observation of debris collected from worn TiN surface (after sliding against steel)51 and during TiC sputtering on steel in the presence of oxygen.52 The change in Gibbs free energy (ΔG) for various reactions, using a commercial software,53 indicates that the formation of FeTiO3 is not feasible beyond 500°C, whereas the change in free energy of formation of TiO2, Fe2O3, or FeO is negative, up to a very high temperature. Therefore, the contact temperature, although not measured experimentally, should be lower than 500°C. The rise in contact temperature, ΔTmax, can be estimated using the classical Archard’s model,54
∆Tm = µ
(πPm )1 / 2 1 / 2 W V, 8k
(24.14)
24.6 INFLUENCE OF TUNGSTEN CARBIDE CONTENT ON SLIDING WEAR PROPERTIES
â•… 393
where μ is the COF, Pm is the Hertzian contact stress at yield, W is the applied load, V is the sliding velocity, k is the thermal conductivity of the cermet flat. The maximum contact temperature is calculated according to Archard and Rowntree55 as: ∆Tmax = 1.67∆Tm.
(24.15)
Thermal conductivity of various cermet compositions can be determined by applying the rule of mixture. Assuming the real contact area to be 10% of the nominal contact area,55 the maximum temperature rise is estimated to be in the range 250–340°C for various cermets. This will enable the formation of a tribochemical layer that is rich in iron oxide or iron titanate.
24.6 INFLUENCE OF TUNGSTEN CARBIDE CONTENT ON LOAD-DEPENDENT SLIDING WEAR PROPERTIES The TiCN–20Ni– (5–25â•›wtâ•›%)WC cermet/steel sliding couples exhibit similar behavior: COF increases abruptly during the first 1000–2000 cycles up to ∼0.7 and reach steady state after about 5000 cycles (Fig. 24.12). All the WC-containing cermets experience stable frictional behavior, except for the cermet containing in 20â•›wtâ•›% WC under lower load of 5â•›N (not shown). Figure 24.13 reveals that the wear volume increases with the load, for TiCN– WC–Ni cermet, as could be expected (Fig. 24.13a). In contrast, an opposite trend is noticed, if wear rate is considered instead of wear volume (see Fig. 24.13b). In general, wear rate is highest at lower load (5â•›N) and systematically decreases to its
1.0
TiCN–20Ni–5WC
0.8
0.8
0.6
0.6
COF
COF
1.0
0.4 5N 20 N 50 N
0.2 0.0
5000
10,000 15,000 Number of cycles (a)
20,000
TiCN–20Ni–25WC
0.4 0.2 0.0
5N 20 N 50 N 5000
10,000 15,000 Number of cycles (b)
20,000
Figure 24.12â•… The COF plots of (a) TiCN–20Ni–5WC and (b) TiCN–20Ni–25WC cermets against steel as function of load and number of cycles.61
394â•…
CHAPTER 24â•… Case Study: Titanium Carbonitride–Nickel-Based Cermets
Wear volume (¥ 10–4 mm3)
16
12
5WC 10WC 15WC 20WC 25WC
8
4
0 5
20
50
Load (N) (a)
Wear rate (¥ 10–7 mm3/N m)
35
5WC 10WC 15WC 20WC 25WC
30 25 20 15 10 5 0 5
20
50
Load (N) (b)
Figure 24.13â•… (a) Wear volume and (b) wear factor of investigated cermets with different WC content at 5, 20, and 50â•›N against steel.37
lowest value with increase in load up to 50â•›N. As a summary, a minimum wear rate of 2.7╯×╯10−7â•›mm3/Nâ•›m is measured at 50â•›N with the lowest WC-content cermet (TiCN–20Ni–5WC), whereas a maximum wear rate of 34╯×╯10−7â•›mm3/Nâ•›m at 5â•›N is recorded with the cermet with the highest WC content (TiCN–20Ni–25WC). It can therefore be concluded that friction and wear of the TiCN–Ni–WC cermets are significantly influenced by both WC addition and applied load. It can be noted here
24.6 INFLUENCE OF TUNGSTEN CARBIDE CONTENT ON SLIDING WEAR PROPERTIES
â•… 395
20 mm (a)
O Ti Fe Ti Ti
20 mm
10 mm (b)
Fe Fe
(c)
Figure 24.14â•… SEM image revealing the presence of a thick tribolayer on worn surface of (a) TiCN–20Ni–5WC cermet at 50-N load, (b) TiCN–20Ni–15WC at 50-N load, and (c) TiCN–20Ni–25WC at 50-N load.37 The EDS analysis of the tribolayer is also shown as an inset in (c).35,37,39,61
that wear rates of 10−5–10−6â•›mm3/Nâ•›m have been reported for TiC-based cermets against steel.7–9 The effect of load on the wear damage of TiCN-based cermets can be realized in Figure 24.14. The worn surface after sliding under a load of 5â•›N is characterized by abrasion and mild tribo-oxidation. With increase in load from 5 to 50â•›N, a thicker tribolayer is observed to form, as seen in Figure 24.14 and it can be suggested that higher oxidation is related to enhancing frictional heating, which is certainly promoted at higher loads. A transition in material removal mechanism from abrasion and mild tribo-oxidation at low loads to extensive tribo-oxidation at high loads can therefore explain a decrease in friction and wear rate. Two major influences of WC content on the tribolayer properties can be proposed, depending on the load: 1. Larger amounts of WC in the cermet provide a large number of interface boundaries, which, under sliding, would provide an increased number of sites for grain attachment. At 5-N load, where the oxide layer was thin and only partially covered the surfaces, the increased WC content obviously increased the abrasive actions, as evident from localized abrasive scratches and pits (Fig. 24.9). Moreover, as the number of detached hard WC particles increases (with
396â•…
CHAPTER 24â•… Case Study: Titanium Carbonitride–Nickel-Based Cermets
higher WC content in cermet), the wear also increases, leading to higher wear loss. Furthermore, increased number of hard WC particles promotes the tribolayer removal and thus reduces the extent of adhered tribolayer at the surface. 2. With increase in load, the oxidation as well as the tribolayer thickness increased (Fig. 24.14). As observed in Figure 24.14, higher amounts of WC clearly resulted in an increase in the fracture of the tribolayer and delaminated regions (Fig. 24.14). Namely, the boundaries of (Ti,W)CN rim phase with TiCN core or Ni-rich binder can cause even more fracture at high loads. This will result in poor oxide layer support, leading to the fracture of the (brittle) oxide layer to a larger extent in cermets containing higher amounts of WC, compared with those with low WC content. The increased wear with WC amount at higher loads (see Fig. 24.13a) correlates well with this suggestion. In summary, the results of the free energy change for various tribochemical reactions proposed are plotted in Figure 24.15, and a schematic of the dominant wear mechanisms for various cermet compositions is presented in Figure 24.16. For clarity, the operating mechanisms are qualitatively illustrated in the form of wear maps, as shown in Figure 24.17. The formation and consequent failure of oxide tribochemical layer are affected by two factors: (1) load and (2) amount of WC. At low loads, the severity of wear is much higher due to a combination of abrasion and grain detachment. 50
Free energy change (kcal)
0 –50 –100 –150 –200 –250 0
200
400
600
800
1000
Temperature (°C) Fe + 0.5O2 = FeO
2Fe + 1.5O2 = Fe2O3
TiC + 1.5O2 = TiO2 + CO
TiN + O2 = TiO2 + 0.5N2
FeO + TiO2 = FeTiO3
Figure 24.15â•… Free energy change (ΔG) as a function of temperature for various possible tribochemical reactions, which were discussed to explain the formation of tribolayer and wear debris at the sliding contacts of TiCN–(20â•›wtâ•›%)Ni–xWC(x╯=╯5–25â•›wtâ•›%)/steel.37
24.7 High Temperature Wear of Titanium Carbonitride–Nickel Cermets
FexOy
TiO2
Wear scar TiCN–20Ni–5WC @5 N
Fe2O3
Underneath WC
Fe9TiO15
TiCN–20Ni–5WC @50 N
WC
FexOy
WC
â•… 397
TiO2
TiCN–20Ni–25WC @5 N
Fe2O3
Fe9TiO15
TiCN–20Ni–25WC @50 N
Figure 24.16â•… Schematic of dominant wear mechanisms operating during sliding of TiCN–Ni–WC cermets at different loads.
WC content
Abrasion and Mild tribo-oxidation and formation and high removal rate of thin oxide layer
Intensive tribo-oxidation and thick oxide layer (Fe2O3, Fe9TiO15), severe fracture of the layer
Mild tribo-oxidation and slow formation and low removal rate of thin oxide layer
Intensive tribo-oxidation and thick oxide layer (Fe2O3, Fe9TiO15), moderate fracture of the layer Load
Figure 24.17â•… Wear map (qualitative) of the TiCN–Ni–WC cermet with varying %WC content, when subjected to sliding against steel at different loads.37
24.7 HIGH TEMPERATURE WEAR OF TITANIUM CARBONITRIDE–NICKEL CERMETS The effect of the addition of different secondary carbides (WC, NbC, TaC, HfC) on the friction properties of TiCN–20Ni-based cermets against steel at high temperature is summarized in Table 24.3. It is evident from Table 24.3 that the COF varies in a narrow range (0.43–0.46) with the change in cermet composition. This suggests that
398â•…
CHAPTER 24â•… Case Study: Titanium Carbonitride–Nickel-Based Cermets
TABLE 24.3â•… Summary of Friction and Wear Data, Measured When TiCN–Ni-Based Cermets were Fretted Against Steel at 550°C Under the Selected Test Conditions, as Reported in Reference 39
Cermet composition
Steady state COF
Wear rate of cermets (×10−6â•›mm3/Nâ•›m)
Wear rate of steel ball (×10−6â•›mm3/Nâ•›m)
TiCN–20Ni TiCN–20Ni–10WC TiCN–20Ni–10NbC TiCN–20Ni–10TaC TiCN–20Ni–10HfC
0.46 0.46 0.46 0.43 0.43
3.2 5.0 5.2 3.5 3.8
7.1 14.1 16.9 12.4 6.5
the presence of WC, NbC, TaC, or HfC does not affect the frictional behavior of TiCN–Ni cermet to any noticeable extent. The average COF of TiCN–Ni–Mo and TiCN–Al2O3–Ni–Mo varied in the range of 0.8–0.85 at 600°C during sliding against silicon nitride balls.30 However, a lower range of COF (0.2–0.3) was recorded when TiC-based ceramics slid against high speed steel at 600°C.29 The COF of TiN coating on steel against steel was around 0.55 at 200°C,56 whereas the COF in the steady state was around 0.4 during fretting tests of TiN coatings at 500°C against corundum balls.57 The wear volumes of TiCN–Ni-based cermets are summarized in Table 24.3. The wear rate of TiCN–20Ni cermets increased with WC and NbC addition, while the addition of TaC and HfC does not have any effect. It has been reported in the literature that TiC-based and TiCN-based ceramics experience wear rates of 10−5–10−7â•›mm3/Nâ•›m.29,30
24.7.1â•…Wear Mechanisms As representation of high temperature wear mechanisms, the overall wear scar of TiCN–20Ni cermet after fretting at 550°C is shown in Figure 24.18a. Figure 24.18b provides the three-dimensional profile of the wear scar of TiCN–20Ni cermet. The adhesion of debris resulted in a rather rough surface (Fig. 24.18c). Moreover, the debris particles are compacted and adhered at the center of the wear scar. The smearing or extrusion of the thin layer can also be observed along the fretting direction in Figure 24.18a. Interestingly, the debris particles have different shapes, such as flake type and round type (see Fig. 24.18c). Also, the size of the spherical debris particles varies over a range from submicron to around 5â•›µm. Such wear features can be correlated with the repeated sliding-induced fatigue stress. The EDS analysis reveals that TixOy, FexOy, or their mixed oxides form during fretting at 550°C against steel. Like the worn cermet, the worn surface of the steel ball, shown in Figure 24.19, reveals similar features after fretting at 550°C. The formation and ductile failure of the tribolayer can be noted in Figure 24.19b. As in the earlier case, the EDS analysis (inset of Fig. 24.19b) indicated the presence of Ti, O, Ni, and Fe only, implying material transfer from cermets.
24.7 High Temperature Wear of Titanium Carbonitride–Nickel Cermets
â•… 399
3.00
2.00
1.00 [V]
0.00 0.00
1.00[x]
(a)
2.00
3.00
(b)
Ti O Ni
Fe Ni
(c)
Figure 24.18â•… (a,c) Representative SEM images demonstrating worn surface features of TiCN–20Ni cermet after fretting at 550°C against steel. (b) Three-dimensional topographical view of the worn surface. The EDS spectrum of the tribolayer is shown as an inset in (c). Arrows indicate fretting direction.39,61
Ti
O
(a)
Fe Ni
(b)
Figure 24.19â•… Typical SEM images revealing worn surface features of steel ball after fretting at 550°C against TiCN–20Ni cermet. The EDS spectrum of the tribolayer is shown as an inset in (b). Arrows indicate fretting direction.39,61
400â•…
CHAPTER 24â•… Case Study: Titanium Carbonitride–Nickel-Based Cermets
Fe O
Ti W
(a)
Fe Ni
(b)
Figure 24.20â•… Representative SEM images demonstrating worn surface features of (a) TiCN–20Ni–10WC cermet and (b) corresponding steel ball after fretting at 550°C. The EDS spectrum of layer on ball surface is shown in (b). Arrows indicate fretting direction.38
The worn surface of TiCN–20Ni–10WC reveals severe abrasion (Fig. 24.20) with a number of characteristic grooves. Compared with the worn surface of TiCN– 20Ni cermet, it is observed that the periphery of the wear scar is covered by an increased amount of debris in the case of TiCN–20Ni–10WC cermets. The worn surface of the steel ball, shown in Figure 24.20b, reveals the presence of a tribolayer. However, the nonprotectiveness of the layer appears to be due to fatigue and deformation (see Fig. 24.20a,b). This indicates that the adhesion of the layer to the worn surface is poor, which corroborates with the support offered by the surface. EDS analysis of the steel ball tribolayer after fretting against TiCN– 20Ni–10WC indicated the presence of major peaks of Ti, Fe, and O as well as a minor W peak (see inset of Fig. 24.20b). Figure 24.21 presents the wear features of TiCN–20Ni–10NbC cermet. EDS analysis of the tribolayer revealed the presence of Nb along with Ti, Fe, Ni, and O (see inset of Fig. 24.21a). Figure 24.21b shows the evidence of a cracked tribolayer on the worn surface of the ball, and EDS analysis indicated that the layer contains Fe, Ti, O, and Ni along with a small amount of Nb (inset of Fig. 24.21b). In contrast, cermet containing TaC addition exhibits a different behavior. Figure 24.22a reveals deformation of the tribolayer as well as entrapment of wear debris and microcracks (Fig. 24.22a). The EDS analysis (inset of Fig. 24.22a) indicates the predominant presence of Fe and O with a small presence of Ti on worn TiCN–20Ni–10TaC. Figure 24.22b illustrates that the worn surface of the steel ball is covered by a thick nonadherent tribolayer, when fretted against TiCN– 20Ni–10TaC cermet. As in the earlier case, EDS analysis of the worn ball surface after fretting against TiCN–20Ni–10TaC cermet indicated the presence of Fe, Ti, and O (inset of Fig. 24.22b). The worn surface of the TiCN–20Ni–10HfC cermet is entirely covered by a thick tribolayer (Fig. 24.23a) and closer observation reveals cracking perpendicular to the fretting direction. Such characteristic features indicate surface fatigue as the
24.7 High Temperature Wear of Titanium Carbonitride–Nickel Cermets
Ti
Fe O Nb
â•… 401
O
Ti
Fe
Nb
Ni
(a)
Ni
(b)
Figure 24.21â•… Typical SEM images illustrating worn surface features of (a) TiCN–20Ni– 10NbC cermet and (b) corresponding steel ball. The respective EDS analyses of tribolayers are shown as insets. Arrow indicates fretting direction.38
O
Fe
Fe Ti
O Ti
(a)
(b)
Figure 24.22â•… Representative SEM images revealing worn surfaces of (a) TiCN–20Ni– 10TaC cermet and (b) corresponding steel ball. The respective EDS analyses of tribolayers are shown as insets. Arrows indicate fretting direction.38
dominant mechanism in the case of TiCN–20Ni–10HfC cermet (Fig. 24.23a). The EDS analysis of the tribolayers of TiCN–20Ni–10HfC (inset of Fig. 24.23a) cermets reveals the predominant presence of Fe and O and minor presence of Ti. Interestingly, EDS analysis (inset of Fig. 24.23b) could not detect the presence of Hf.
24.7.2â•… Discussion of High-Temperature Oxidation and Its Relation to Material Removal During high-temperature wear experiments, the cermet samples as well as balls were kept in the heating chamber and allowed to heat up to 550°C and, when the
402â•…
CHAPTER 24â•… Case Study: Titanium Carbonitride–Nickel-Based Cermets
O
Fe O
Fe
Ti
(a)
(b)
Figure 24.23â•… SEM images of the worn surfaces of (a) TiCN–20Ni–10HfC cermet and (b) corresponding steel ball. The respective EDS analyses of tribolayers are shown as insets. Arrows indicate fretting direction.38
temperature had stabilized, the wear tests were performed. Therefore, the tribological surfaces were affected due to initial oxidation, prior to fretting tests.
24.7.3â•… Thermal Oxidation The oxidation of iron (steel ball) during high-temperature exposure results in the formation of multilayered scale containing FeO, Fe3O4, and/or Fe2O3.58 The possible oxidation reactions with corresponding change in Gibbs free energy of formation (ΔGf) at 550°C are summarized here53:
Fe + (1/ 2)O2 → FeO ∆G f = −50 kcal/mol, 2Fe + (3/ 2)O2 → Fe 2 O3 ∆G f = −34.44 kcal/mol 3Fe + 2O2 → Fe 3 O 4 ∆G f = −48.32 kcal/mol.
(24.16) (24.17) (24.18)
According to the Fe–O phase diagram, the wustite phase FeO can form above 570°C; therefore FeO, Fe3O4, and F2O3 can form in equilibrium with FeO on the tribological surfaces. As far as the cermet oxidation is concerned, the following thermodynamically feasible reactions can be proposed:
Ni + O2 → NiO ∆G f WC + O2 → WO3 + CO ∆G f NbC + O2 → Nb2 O5 + CO ∆G f TaC + O2 → Ta 2 O5 + CO ∆G f HfC + O2 → HfO2 + CO ∆G f
= −39 kcal/mol, = −186 kcal/mol, = −391 kcal/mol, = − 423 kcal/mol, = −226 kcal/mol.
(24.19) (24.20) (24.21) (24.22) (24.23)
Following an earlier literature report, oxidation of the binder phase results in outward migration of respective elements, followed by inward diffusion of oxygen.33 Also, the spontaneous thermal oxidation of TiN on exposure to ambient air at 500°C
24.8 Summary of Key Results
â•… 403
is sufficient for crystallization of Ti–O.59 As TiCN can be simplistically described as a solid solution of TiC and TiN, the oxidation of the TiCN can be described by their respective oxidation reactions:
TiC + (3/ 2)O2 → TiO2 + CO ∆G f = −192 kcal/mol, TiN + O 2 → TiO 2 + (1/ 2)N 2 ∆G f = −128 kcal/mol.
(24.24) (24.25)
Also, the oxygen dissolution in Hf is reported to be more than other metals (20â•›at%).58 The oxidation of W at moderate and high temperatures (>800°C) results in the formation of the volatile WO3.58 During oxidation of Ta in the range 450– 600°C, the initial parabolic oxidation is followed by a rapid breakaway oxidation. On the other hand, the Nb forms stable NbO and NbO2 and Nb2O5. However, at high temperature (>650°C), Nb2O5 scale cannot offer any resistance to oxidation.58
24.7.4â•… Influence of Different Secondary Carbide Addition The frictional heat, generated at the sliding contacts, oxidizes the debris particles during room temperature testing and the hard oxide debris particles abrade, resulting in severe wear of surfaces. Thus, sliding of surfaces covered with tribochemical layers results in mild wear. Therefore, the rate of tribochemical layer formation must dominate over the rate of removal in order to maintain lower wear. SEM-EDS analysis reveals that, under sliding conditions, the compaction of wear debris results in the tribolayer formation and their subsequent transfer to the countersurface. The subsequent chemical diffusion of oxide products through the interface potentially leads to the formation of mixed oxides of iron, titanium, and nickel, as observed in the EDS analysis (see Fig. 24.18). Therefore, it can be suggested that initial thermal oxidation of the surface, followed by frictional oxidation at the contact area, results in the formation of a tribolayer containing iron, titanium, and nickel oxides and subsequent diffusion to the countersurfaces. In the context of high-temperature wear behavior, the properties or stability of the tribolayer particularly play an important role.60 The stability of the tribolayer depends on the nature of oxides as well as support by the underlying material and also on the operating parameters (speed, temperature, etc.). Sufficient resistance to plastic deformation or fracture is necessary to support oxide layers during sliding. In the case of TiCN–Ni-based cermets, the nature and stability of layers is influenced by the addition of secondary carbides: WC, NbC, HfC, or TaC.
24.8 SUMMARY OF KEY RESULTS Summarizing the discussion of fretting properties at low load, it can be said that the fretting wear behavior of TiCN–Ni cermets is influenced by two factors: (1) load and (2) cermet composition. At high load, three-body abrasion and material transfer dominate and the mating materials experience a decrease in wear. On the other hand, the resistance to abrasion for cermets with lower hardness is decreased. Against the backdrop of the preceding considerations, the addition of HfC should not be considered for the TiCN–Ni tribosystem.
404â•…
CHAPTER 24â•… Case Study: Titanium Carbonitride–Nickel-Based Cermets
The discussion in this chapter also implies that the addition of WC to TiCN–Ni cermet must be limited to 5â•›wtâ•›% for tribological applications involving sliding contact (up to a load of 50â•›N). In the future, similar studies to probe into the influence of the addition of other secondary carbides (NbC, TaC, HfC) should be undertaken. Such comprehensive work would provide guidelines for designing appropriate cermet compositions for a wider range of tribological applications. Summarizing the high-temperature tribological behavior, the effect of secondary carbides appears to have negligible influence on the frictional behavior of TiCN–Ni-based cermets as both the mating surfaces are covered with iron titanium oxides. However, the wear resistance is found to be dependent on the type of secondary carbide, as the support of the layers from the underlying oxide species is different. As the oxidation of tungsten leads to volatile WO3 formation at high temperature, the diffusion of the same through the iron titanium oxide to the other surface layer increases the severity of abrasion. The addition of TaC appears to enhance the resistance to plastic deformation. The stress-induced cracking of Nb2O5 increased the wear rate of NbC-containing cermets. The indirect effect of harder HfO2 enhanced the wear resistance of HfC-containing TiCN–Ni cermets, and this is possible partly because of the formation of a thick tribolayer on the countersurface.
REFERENCES ╇ 1â•… P. Ettmayer, H. Kolaska, W. Lengauer, and K. Dreyer. Ti(C,N) cermets-metallurgy and properties. Int. J. Refract. Met. Hard Mater. 13 (1995), 343–351. ╇ 2â•… S. Ahn and S. Kang. Formation of core/rim structures in Ti(C,N)-WC-Ni cermets via a dissolution and precipitation process. J. Am. Cer. Soc. 83(6) (2000), 1489–1494. ╇ 3â•… S. Ahn and S. Kang. Effect of various carbides on the dissolution behavior of Ti(C0.7 N0.3) in a Ti(C0.7 N0.3)-30 Ni system. Int. J. Refract. Met. Hard Mater. 19 (2001), 539–545. ╇ 4â•… W. T. Kwon, J. S. Park, S. W. Kim, and S. Kang. Effect of WC and group IV carbides on the cutting performance of Ti(C,N) cermet tools. Int. J. Mach. Tool Manufact. 44 (2004), 341–346. ╇ 5â•… S. Mun and S. Kang. Effect of HfC addition of microstructure of Ti(CN)-Ni system. Powder Metall. 42(3) (1999), 251–256. ╇ 6â•… U. Rolander, G. Winel, and M. Zwinkls. Effect of Ta structure on structure and mechanical properties of (Ti,Ta,W)-Co cermets. Int. J. Refract. Met. Hard Mater. 19 (2001), 325–328. ╇ 7â•… F. Qi and S. Kang. A study on microstructural changes in Ti(CN)–NbC–Ni cermets. Mater. Sci. Eng. A 251 (1998), 276–285. ╇ 8â•… H. Suzuki and H. Matsubara. Some properties of Ti(C, N)-WC-NI alloy. Jpn. Soc. Powder Metall. 33(4) (1986), 199–203. ╇ 9â•… I. M. Hutchings. Tribology: Friction and Wear of Engineering Materials. Butterworth-Heinemann Publications, Boston, 1992, 150–156. 10â•… K.-H. Zum Gahr. Microstructure and Wear of Materials. Elsevier, Amsterdam, 1987, 352–355. 11â•… T. Hisakado and N. Hashizume. Effect of normal loads on the friction and wear properties of metals and ceramic against cermet in vacuum. Wear 237 (2000), 98–106. 12â•… J. Pirso, M. Viljus, and S. Letunovitš. Sliding wear of TiC–NiMo cermets. Tribol. Int. 37 (2004), 817–824. 13â•… C. C. Degnan, P. H. Shipway, and J. V. Wood. Elevated temperature sliding wear behavior of TiC–reinforced steel matrix composites. Wear 251 (2001), 1444–1451. 14â•… M. Komac and S. Novak. Mechanical and wear behaviour of TiC-cemented carbides. Int. J. Refract. Met. Hard Mater. 4 (1985), 21–26.
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15â•… M. B. Peterson and R. E. Lee Jr. Sliding characteristics of the metal-ceramic couple. Wear 7(4) (1964), 334–343. 16â•… H. Engqvist, H. Högberg, G. A. Botton, S. Ederydand, and N. Axén. Tribofilm formation on cemented carbides in dry sliding conformal contact. Wear 239(2) (2000), 219–228. 17â•… A. Mukhopadhyay and B. Basu. Recent development of WC-based cermets and nanocomposites. J. Mater. Sci. 46 (2011), 571–589. 18â•… A. Blomberg, J. Lu, and S. Hogmark. An electron microscopy study of worn ceramic surfaces. Tribol. Int. 26 (1993), 369–381. 19â•… X. Zhao, J. Liu, B. Zhu, Z. Luo, and H. Miao. Effects of lubricants on friction and wear of Ti(CN)/1045 steel sliding pairs. Tribol. Int. 130(3) (1997), 177–182. 20â•… K. Kameo, K. Friedrich, J. F. Bartolome, M. Diaz, L.-E. Sonia, and J. S. Moya. Sliding wear of ceramics and cermets against steel. J. Eur. Cer. Soc. 23 (2003), 2867–2877. 21â•… E. T. Jeon, J. Joardar, and S. Kang. Microstructure and tribo-mechanical properties of ultrafine Ti(CN) cermets. Int. J. Refract. Met. Hard Mater. 20(3) (2002), 207–211. 22â•… D. Sarkar, B. V. Manoj Kumar, S. Ahn, S. Kang, and B. Basu. Fretting wear behavior of Ti(CN)based advanced cermets. Key Eng. Mater. 264–268 (2004), 1115–1118. 23â•… I. M. Hutchings. Tribology: Friction and Wear of Engineering Materials. Edward Arnold, London, 1992. 24â•… Y. R. Liu, J. J. Liu, B. L. Zhu, Z. B. Luo, and H. Z. Miao. The computer simulation of the temperature distribution on the surface of ceramic cutting tools. Wear 210 (1997), 39–44. 25â•… Z. Xingzhong, L. Jiajum, M. Hezhuo, and L. Zhenbi. Wear mechanisms of Ti(C, N) ceramic in sliding contact with stainless steel. J. Mater. Sci. 32 (1997), 2963–2968. 26â•… J. K. Lancaster. The influence of temperature on metallic wear. Proc. Phys. Soc. B 70 (1957), 112–118. 27â•… A. K. Vijh. The influence of metal-metal bond energies on adhesion, hardness and wear of metals. J. Mater. Sci. 10 (1978), 998–1004. 28â•… U. Persson, H. Chandrasekaran, and A. Merstallinger. Adhesion between some tool and work materials in fretting and relation to metal cutting. Wear 249 (2001), 293–301. 29â•… Z. Xingzhhong, L. Jiajun, Z. Baolinag, O. Jinlin, and X. Qunji. Tribological properties of TiC-based ceramic/high speed steel pairs at high temperatures. Ceram. Int. 24 (1998), 13–18. 30â•… J. Meng, J. Lua, J. Wang, and S. Yang. Tribological behavior of TiCN-based cermets at elevated temperatures. Mater. Sci. Eng. A 418 (2006), 68–76. 31â•… J. Joardar, S. W. Kim, and S. Kang. GI-XRD studies on surface structure of ultrafine Ti(C0.5N0.5)WC-Ni cermets at high temperature. Wear 261(3–4) (2006), 360–366. 32â•… F. Akhtar and S. J. Guo. Microstructure, mechanical and fretting wear properties of TiC-stainless steel composites. Mater. Character. 59(1) (2008), 84–90. 33â•… F. Monteverde and A. Bellosi. Oxidation behavior of titanium carbonitride based materials. Corros. Sci. 44 (2002), 1967–1982. 34â•… D. S. Park, C. Park, and Y. D. Lee. Oxidation of Ti(C,N)-based ceramics exposed at 1373â•›K in air. J. Am. Cer. Soc. 83 (2000), 672–674. 35â•… B. V. Manoj Kumar and B. Basu. Fretting wear properties of TiCN-Ni cermets: Influence of load and secondary carbide addition. Metall. Mater. Trans. A 39(3) (2008), 539–550. 36â•… B. V. Manoj Kumar, B. Basu, M. Kalin, and J. Vizintin. Tribochemistry in sliding wear of TiCN-Ni based cermets. J. Mater. Res. 23(5) (2008), 1214–1227. 37â•… B. V. Manoj Kumar, B. Basu, M. Kalin, and J. Vizintin. Load dependent transition in sliding wear properties of TiCN-WC-Ni cermets. J. Am. Cer. Soc. 90(5) (2007), 1534–1540. 38â•… B. V. Manoj Kumar and B. Basu. Mechanisms of material removal during high temperature fretting of TiCN-Ni based cermets. Int. J. Refract. Met. Hard Mater. 26 (2008), 504–513. 39â•… B. V. Manoj Kumar. Understanding the tribological properties of TiCN-Ni based cermets. Ph.D. thesis, Indian Institute of Technology Kanpur, India, November 2007. 40â•… D. K. Shetty, I. G. Wright, P. N. Mincer, and A. H. Clauer. Indentation fracture of WC-Co cermets. J. Mater. Sci. 20 (1985), 1873–1882. 41â•… I. Hussainova. Microstructure and erosive wear in ceramic-based composites. Wear 258 (2005), 357–365. 42â•… J. Bijwe and J. Indumathi. Influence of fibers and solid lubricants on low amplitude oscillating wear of polyetherimide composites. Wear 5–6 (2004), 562–572.
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43â•… E. Marui, M. Hashimoto, S. Kato, and W. Kojima. Dissipation of kinematic energy by slip at the interface of mating surfaces. Wear 159 (1992), 141–150. 44â•… M. Z. Huq and J.-P. Celis. Reproducibility of friction and wear results in ball-on-disc unidirectional sliding tests of TiCN-alumina parings. Wear 212 (1997), 151. 45â•… H. Czichos. Tribology. Elsevier, Amsterdam, 1978. 46â•… A. G. Evans and D. B. Marshall. In Fundamentals of Friction and Wear of Materials, D. A. Rigney (Ed.). Amer. Soc. Metals, Ohio, 1981, 439–452. 47â•… B. Bhushan. Principles and Applications of Tribology. John Wiley & Sons, Inc., New York, 1999, 507–509. 48â•… J. F. Archard. Contact and rubbing of flat surfaces. J. Appl. Phys. 24 (1953), 981–988. 49â•… S. E. Haggerty and D. Rumble III (Ed.). Oxide Minerals. Mineralogical Society of America, Washington, DC, 1981, Hg-103. 50â•… E. M. Levin, C. R. Robbins, and H. F. McMurdie. In Phase Diagrams for Ceramists, 4th ed. M. K. Reser (Ed.). American Ceramic Society, Columbus, OH, 1964, 62. 51â•… I. L. Singer, S. Fayeulle, and P. D. Ebnitt. Friction and wear behavior of TiN in air: The chemistry of transfer films and debris formation. Wear 149 (1991), 375–394. 52â•… J. E. Greene and J. L. Zilko. The nature of the transition region formed between dc-biased rf sputtered TiC films and steel substrates. Surf. Sci. 78 (1978), 109–124. 53â•… A. Roine. HSC Chemistry. Version 5.1 Outokumpu research Oy, Pori, Finland, 2002. 54â•… J. F. Archard. The temperature of rubbing surfaces. Wear 2 (1958–1959), 438–455. 55â•… J. F. Archard and R. A. Rowntree. Metallurgical phase transformations in the rubbing of steels. Proc. R. Soc. London A418 (1988), 405–424. 56â•… T. Polcar, T. Kunart, S. Novak, L. Kopecky, and P. Siroky. Comparison of tribological behavior of TiN, TiCN and CrN at elevated temperatures. Surf. Coat. Technol. 193 (2005), 192–199. 57â•… A. Ramalho and J.-P. Celis. High temperature fretting behavior of plasma vapor deposition TiN coatings. Surf. Coat. Technol. 155 (2002), 169–175. 58â•… P. Kofstad. High Temperature Oxidation of Metals. Wiley, New York, 1966. 59â•… L. S. Hsu, R. Rujkorakarn, J. R. Sites, and C. Y. She. Thermally induced crystallization of amorphous titania films. J. Appl. Phys. 59(10) (1986), 3475–3480. 60â•… R. B. Waterhouse. Fretting at high temperatures. Tribol. Int. 14 (1981), 203–207. 61â•… B. V. Manoj Kumar. Unpublished work, 2011.
CHAPTER
25
CASE STUDY: (W,Ti)C–CO CERMETS In an effort to develop cermets with better combinations of mechanical and tribological properties than one could achieve with conventional WC–Co cermets, (W,Ti) C–(20â•›wtâ•›%)Co cermets are being developed using a multistage pressureless sintering route. In this chapter, the wear-resistance data are discussed in the light of the difference in mechanical properties. Broadly, abrasive wear is the dominant wear mechanism and localized spalling of the tribochemical layer also partly contributes to the fretting damage of the mixed carbide cermets.
25.1 INTRODUCTION In the field of advanced ceramics, considerable research efforts are being invested to obtain better properties with transition-metal carbides such as WC, TiC, and TaC as a single phase or through a cermet approach using metallic binder, such as Co or Ni.1–6 In particular, WC-based hardmetals are widely used for wear-resistance applications. WC, because of its high melting point and high hardness (∼16–22â•›GPa), is one of the constituent phases in hardmetals (WC–Co). Almost 70% of the cutting tools in use today are made of cemented carbides, with WC and Co as major phases.7 Apart from the cutting tool applications, WC-based cermets also have established major dominance in various other applications, that is, in components of drilling and mining equipments,8 bearings,9 and face seals.10,11 Arenas and co-workers12 investigated the effect of TiC addition on wear behavior of WC/Co hardmetals against Cr-steel. They reported wear rates of 10−6â•›mm3/Nâ•›m for WC–TiC–Co cermets. The fretting wear behavior of binderless WC, fabricated by the plasma pressure compaction (P2C) process, was investigated by Staia et al.13 Pure WC, sintered using the P2C process, exhibited an extremely low wear rate on the order of 10−7â•›mm3/Nâ•›m. Larsen-Basse14 reported that sliding, under high loading and at low sliding speed, causes extrusion of Co binder, which leads to microfracturing and pullout of the carbide grains in WC–Co cermet. Trent15 observed that the diffusive wear is minimal in the case of WC–TiC–Co cermets. This is possibly due
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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CHAPTER 25â•… CASE STUDY: (W,Ti)C–CO CERMETS
to the formation of mixed carbides (WC/TiC), which in turn have lower solubility in the contacting steel counterbody. In another study, Jia and Fischer16 performed unlubricated fretting wear of WC–Co against Si3N4 (disk). The wear rate measured was proportional to the hardness of WC with average wear rate on the order of 10−6â•›mm3/Nâ•›m. Guilemany et al.17 performed fretting wear experiments on WC–Co coating against different mating materials, such as Al2O3, Si3N4, and martensitic steel. The excellent sliding properties of the hardmetal coating were reported to be due to the formation of oxides during the sliding process, which acted as a solid lubricant and, thereby, resulted in low friction coefficients. This chapter discusses some salient results to illustrate fretting wear of (W,Ti)C–(20%)Co cermets, processed from premixed WC/TiC powders (WCâ•›:â•›TiC╯=╯7:3) and more details can be found elsewhere.18
25.2 MATERIALS AND EXPERIMENTS Both (W,Ti)C solid solution powders and premixed WC/TiC powders were sintered in the temperature range 1400–1500°C to obtain mixed carbide cermets. The details of sample designation and sintering conditions as well as their mechanical properties are given in Table 25.1 and more details are reported in a research publication.18 The tribological results that are presented in this chapter were obtained using fretting wear tests at a constant frequency (8â•›Hz), a constant displacement stroke (50â•›µm), and for a large number of cycles (100,000), while keeping load (2–10â•›N) as the external variable of the tribosystem.
TABLE 25.1â•… Mechanical Properties of the Sintered (W,Ti)C–(20%)Co Cermets18
Sample composition (W,Ti)C–(20%)Co cermets (solid solution powder) (W,Ti)C–(20%)Co cermets (solid solution powder) (W,Ti)C–(20%)Co cermets (solid solution powder) W,TiC–(20%)Co cermet (premixed powder)
Sintering condition
Density (g/cm3)
Elastic modulus (GPa)
Hardness (Hv, GPa)
Fracture toughness (KIC, MPaâ•›m1/2)
1500°C, 1 hour
9.27
382
11.0╯±â•›â•¯0.2
14.8╯±â•¯1.1
1450°C, 1 hour and 1500°C, 1 hour 1400°C,1 hour; 1450°C,1 hour; and 1500°C,1 hour 1500°C, 1 hour
9.51
457
16.0╯±â•¯0.2
7.7╯±â•¯0.4
9.57
467
17.1╯±â•¯0.4
6.9╯±â•¯0.4
9.18
409
11.9╯±â•¯0.2
14.0╯±â•¯1.3
25.3 Microstructure and Mechanical Properties
â•… 409
25.3 MICROSTRUCTURE AND MECHANICAL PROPERTIES Various microstructural phases in Figure 25.1 include round-shaped grains (gray contrast), bright contrast faceted crystals, and darker contrast matrix phases. Compared with the microstructure prepared by single-stage sintering as shown in Figure 25.1a, a multistage sintering is effective, especially in increasing the number density and decreasing the average size of WC grains. It is also observed that in comparison with cermet processed from solid solution powder (S1 sample), the (W,Ti)C–Co cermet, processed from the premixed WC/TiC powder mixture and sintered at 1500°C/h contains finer faceted WC grains along with (W,Ti)C grains.
10 µm
(a)
10 µm
(b)
10 µm
(c)
10 µm
(d)
Figure 25.1â•… SEM image of polished surface microstructure of cermet sample processed from the following: (a) solid solution of (W,Ti)C powders at 1500°C/h; (b) solid solution of (W,Ti)C powders at 1450 and 1500°C/h; (c) solid solution of (W,Ti)C powders at 1400, 1450, and 1500°C/h; and (d) mixture of WC and TiC powders sintered at 1500°C/h.18
410â•…
CHAPTER 25â•… CASE STUDY: (W,Ti)C–CO CERMETS
Table 25.1 summarizes the density and mechanical properties data obtained with the developed cermets. A maximum density of 9.57â•›gm/cm3 is recorded with S3 cermet. Considering the density of WC (15.7â•›g/cm3), TiC (4.94â•›g/cm3), and Co (8.9â•›g/cm3), it has been found that the theoretical density of (W,Ti)C–Co cermet processed from premixed powders is 9.39â•›g/cm3. The elastic property data correlate well with the density data. The E-modulus varies in the range of 382–467â•›GPa and the highest modulus is measured with the densest cermet. In the case of cermet prepared with (W,Ti)C solid solution, the maximum hardness (∼17â•›GPa) is exhibited by the cermet processed by three-stage sintering (1400, 1450, and 1500°C/h). The cermet prepared with premixed WC and TiC powders exhibits a hardness of ∼12â•›GPa. The fracture toughness value varies between 7 and 15â•›MPaâ•›m1/2 for the cermet prepared from the solid solution powders; the cermet prepared from premixed WC/TiC powders has toughness ∼14â•›MPaâ•›m1/2. The crack-microstructure interaction is illustrated in Figure 25.2. It can be mentioned here that, using a high-frequency inductionheated combustion synthesis (HFIHCS) route, Kim and coworkers earlier developed WC–(5â•›wtâ•›%)Co cermets with fracture toughness of 7â•›MPa1/2 and hardness of 21â•›GPa.19 Kangwantrakool and Shinohara20 reported that WC–Co–TiC–Al2O3 composites derived from a spark plasma sintering process exhibit an improved hardness (18â•›GPa) after sintering at 1480°C for 5 minutes, when compared with the composite with similar composition (hardnessâ•›∼â•›15â•›GPa) hot pressed at 1480°C for 1 hour.
25.4 WEAR PROPERTIES Some interesting observations can be made from the plots of specific wear rate with hardness and abrasion parameter, as shown in Figures 25.3a and 25.3b, respectively. Figure 25.4 plots the maximum wear depth versus normal load for various cermets. A common observation is that the wear depth systematically increases with increase in load for all cermets. Among the cermets prepared from solid solution powders, the cermet with higher hardness exhibits the lowest wear depth (∼1–4â•›µm), while the cermet prepared from the premixed WC/TiC powders had a higher measured wear depth ranging from 3.6 to 5.4â•›µm for varying load of 2–10â•›N (Fig. 25.4). Summarizing the wear data reveals that the wear resistance of the developed cermet depends on the mechanical properties and processing conditions. Figure 25.5 reveals the topographical features of the fretted surface of the cermet prepared from (W,Ti)C solid solution powders. Broadly, abrasive wear dominates the material removal at all loads, and not much apparent variation in severity of abrasion was noted. The energy-dispersive x-ray spectrometry (EDS) compositional analysis of the wear debris and transfer layer reveals the presence of iron transferred from the counterbody (bearing steel), besides the stronger presence of W, Ti, and O peaks. The noticeable presence of oxygen indicates the possible formation of tribo-oxides. Figure 25.6 presents representative scanning electron microscopy (SEM) images of the wear scar on the (W,Ti)C–(20â•›wtâ•›%)Co cermet, presintered at 1450°C for 1 hour followed by final sintering at 1500°C for 1 hour. Mild abrasive grooves along with fine wear debris are observed over the fretted surface (Fig. 25.6a). In
25.4 Wear Properties
10 µm (a)
10 µm (b)
10 µm (c)
â•… 411
10 µm (d)
Figure 25.2â•… SEM image revealing the crack propagation behavior around the edges of Vickers indentation (Hv10) in S1 cermet revealing the crack blunting by hard WC particle (a), significant crack deflection contribution to the toughness of S2 (b) and S3 (c) cermet, and crack bridging and lateral crack generation in S4 cermet (d).18
Figure 25.6b, one can observe the cracking-induced spalling of the base material at the edge of the wear scar as well as the accumulation of wear debris around the cracked region. The EDS analysis shows the presence of base materials with a lesser amount of Fe. Figure 25.7 reveals the topographical features of the worn surfaces on the (W,Ti)C–(20â•›wtâ•›%)Co cermet, subjected to a three-stage sintering schedule of 1400, 1450, and 1500°C/h. The severity of abrasion-induced wear increases from intermediate to higher (from 5 to 10â•›N) load (Fig. 25.7d,f). The morphological features of the worn scar on the cermets prepared from the premixed WC and TiC powders are displayed in Figure 25.8. EDS analysis shows the presence of a stronger O peak and the presence of elements from the base material as well as iron from the counterbody, indicating the formation of WO3 as well as other oxides. At higher load (10â•›N), there
412â•…
CHAPTER 25â•… CASE STUDY: (W,Ti)C–CO CERMETS
Specific Wear Rate (× 10–7mm3/N m)
24
2N 5N 10 N
22 20 18 16 14 12 10 8 6 11
12
13
14
15
16
17
18
Hardness (Hv) (GPa)
Specific Wear Rate (× 10–7 mm3/N m)
(a) 22
S1 S2 S3 S4
20 18 16 14 12 10 8 6 0.060
0.065
0.070
0.075
0.080
0.085
0.090
Abrasion Parameter ([W 1/8/(Kc1/2 × H5/8)]) (b)
Figure 25.3â•… (a), Variation of measured specific wear rate with hardness (Hv) of the investigated (W,T)C–Co cermets, processed under varying sintering conditions. (b) Variation of specific wear rate measured for various cermets against abrasion parameter. Fretting conditions: duration, 100,000 cycles; frequency, 8â•›Hz; and stroke length, 100â•›µm; at various loads (2, 5, and 10â•›N).18
is formation of a nonprotective large tribolayer inside the wear scar along with fine wear debris with microcracks running over it (pointed arrow in Fig. 25.8c). Summarizing the analysis of the worn surface topography indicates that abrasion and mild oxidative wear are the dominant wear mechanisms for the investigated cermets. Occasionally, spalling on a small scale is observed for some cermets.
25.5 Correlation between Mechanical Properties and Wear Resistance
5.5 5.0
Wear Depth (mm)
4.5
â•… 413
s-1 s-2 s-3 s-4
4.0 3.5 3.0 2.5 2.0 1.5 1.0 2
4
6
8
10
Load (N)
Figure 25.4â•… Plot of the maximum wear depth on the various cermets (flat) versus normal load. The maximum depth of wear scar is measured using laser surface profilometer. Fretting conditions: frequency, 8â•›Hz; displacement, 100â•›µm; duration, 100,000 cycles; with varying load (2, 5, and 10â•›N).18
25.5 CORRELATION BETWEEN MECHANICAL PROPERTIES AND WEAR RESISTANCE One of the important aspects related to the development of new wear-resistant materials is establishing a relationship between mechanical and tribological properties. Some important issues need to be addressed, including the following: (1) How do the different processing parameters influence the mechanical properties? (2) How is the wear behavior of the newly developed cermets influenced by varying hardness and fracture toughness? (3) What are the dominant material removal processes resulting from the fretting wear? and (4) How are the tribological properties critically dependent on the microstructure of cermets prepared from (W,Ti)C solid solution and from the premixed WC and TiC powders? The wear rate linearly decreases with increase in hardness (Fig. 25.3) for all the cermets, irrespective of starting powder or sintering conditions. The inverse proportionality of the wear rate with hardness has been well established for two-body abrasion as well adhesive wear of the metallic materials.21 According to the classical Archard’s equation,22 the total volume of the materials displaced or worn away (V) is given by:
V =k
Wx , H
(25.1)
414â•…
CHAPTER 25â•… CASE STUDY: (W,Ti)C–CO CERMETS
W
W O Fe
Ti
Co Fe W
OCo
50 µm
2N
Ti Co Fe W
50 µm
5N
(a)
(b) W Fe O
Ti
Fe W Co
50 µm
10 N (c)
Figure 25.5â•… SEM images revealing the fretted surface of the (W,Ti)C–(20â•›wtâ•›%)Co prepared from starting powders of (W,Ti)C solid solution, sintered at 1500°C/h. Double-pointed arrow indicates the fretted direction. Fretting conditions: frequency, 8â•›Hz; displacement, 100â•›µm; duration, 100,000 cycles; with varying loads: 2â•›N (a), 5â•›N (b), and 10â•›N (c); and counterbody, bearing grade steel.18
where k is the nondimensional wear coefficient (which typically ranges from 10−6 to 10−1), x is the sliding distance, W is the applied load (N), and H stands for the hardness of the material (GPa). From this it can be stated that hardness is among the important parameters in determining the wear resistance of the (W,Ti)C–(20%)Co cermet materials. Such a strong dependence of hardness dependence on wear rate indicates that (W,Ti) C–(20%)Co cermets behave more like a metallic material at the tribocontact. Further, it can be recalled that the mechanism of material removal is dominated by abrasive wear, as mentioned earlier while reporting the worn surface topographical observation. However, Archard’s model is valid for metallic materials, which undergo plastic deformation at tribocontact. In the present case, the amount of metallic binder (matrix) is limited to 20â•›wtâ•›% in the cermets and hence the wear cannot be controlled
25.5 Correlation between Mechanical Properties and Wear Resistance
â•… 415
W Ti Co Fe W
Ti Co
50 µm
2N
50 µm
2N
(a)
(b) Ti
W
W Co
Ti Ti Co Co W
5N
25 µm
Fe
W Co
25 µm
5N
(c)
Ti
(d) W Ti Co
10 N
50 µm (e)
10 N
Ti
Fe CoW
10 µm (f )
Figure 25.6â•… Topographical features of the as-worn surface, as revealed by SEM, on (W,Ti)C–(20â•›wtâ•›%)Co cermet, pressureless sintered in two stages (1450 and 1500°C/h), fretted for 100,000 cycles at various loads: 2â•›N (a,b), 5â•›N (c,d), and 10â•›N (e,f). Double-pointed arrow indicates the fretting direction. EDS results of the noticeable tribological features are given as insets in the corresponding micrographs.18
416â•…
CHAPTER 25â•… CASE STUDY: (W,Ti)C–CO CERMETS
W
W Ti
50 µm
2N
O Ti
CoW
O Ti
Co Fe W
50 µm
2N
(a)
Ti
(b) W O Ti
Ti
Co Fe W
50 µm
5N
50 µm
5N
(c)
(d) W Ti O Ti
50 µm
10 N (e)
Fe W Co
50 µm
10 N (f )
Figure 25.7â•… Worn surface of the (W,Ti)C–(20â•›wtâ•›%)Co prepared with (W,Ti)C solid solution, pressureless sintered in three stages (1400, 1450, and 1500°C/h). Fretting conditions: frequency, 8â•›Hz; displacement, 100â•›µm; duration, 100,000 cycles; with varying load (2, 5, and 10â•›N); counterbody, bearing grade steel. Double-pointed arrow indicates the fretting direction.18
25.5 Correlation between Mechanical Properties and Wear Resistance
â•… 417
O W Ti Fe Fe CoW
W O Ti
Co Ti Fe W
20 µm
2N
50 µm
5N
(a)
(b) W O Ti Fe Co Ti W
50 µm
10 N (c)
Figure 25.8â•… SEM images revealing the details of the worn surface on the (W,Ti) C–(20â•›wtâ•›%)Co cermets, prepared from premixed WC and TiC powders and sintered at 1500°C/h. Fretting experiments were carried out at varying loads: (a) 2â•›N, (b) 5â•›N, and (c) 10â•›N. Fretting conditions: frequency, 8â•›Hz; displacement, 100â•›µm; duration, 100,000 fretting cycles; counterbody, bearing grade steel. Double-pointed arrow indicates the fretting direction.18
entirely by deformation-dominated abrasion. Following the assumption that the material removal for a brittle ceramic predominantly takes place by sharp-asperityinduced lateral crack formation, the dependence of wear rate (wear volume normalized with respect to normal load and sliding distance) V′ on material properties can be expressed as:23
V′ ∝
W 1/ 8 1/ 2 K IC H 5/8
(25.2)
where V′ is the wear volume (normalized with respect to normal load), W is the normal contact force, KIC is the fracture toughness, and H is the hardness. The term on the right-hand side of Equation 25.2 is known as the abrasion parameter, which will be used in further discussion.
418â•…
CHAPTER 25â•… CASE STUDY: (W,Ti)C–CO CERMETS
The data plotted in Figure 25.3b indicate that the wear rate is inversely proportional to the abrasion parameter for our cermets. If the cermets exhibited tribological response like a brittle material, the wear rate should have a direct proportionality relationship with abrasion parameter. From the preceding discussion, it is quite clear that wear behavior of the cermets does not follow the material removal behavior of classical brittle materials (brittle fracture and cracking). The high wear resistance of the cermets can also be explained from the wear model proposed by Roberts:24 3
P* =
54.47β K IC K IC πηθ H
(25.3)
where P* is the minimum load required to produce fracture from a point contact (N), β is the constant relating hardness to indentation diagonal (2.16 for Vickers indentation), η is a constant, θ is the geometrical constant (≈0.2), KIC is the fracture toughness of the material indented (MPa m½), and H is the hardness of the material indented (GPa). Incorporating the material property data for cermets from Table 25.1 provides an estimated load (P*) of more than 100â•›N for the investigated cermets. This indicates that the material removal in cermets is highly unlikely to take place by severe brittle fracture under the tribological regime. Based on the results discussed in this chapter, a summary of the fretting mechanisms for investigated cermets would include the following: (1) generation of lower amounts of debris accumulated with mild abrasive wear (at low load); (2) formation of a tribochemical layer, in addition to abrasion, enriched with W, Ti, Fe, Co as a result of agglomeration of the wear debris generated (at intermediate load); and (3) spalling of material as well as smearing of the tribolayer, in addition to abrasion (at higher load). As a concluding note, the experimental results demonstrate that the refinement of sintering conditions can lead to improved hardness and higher wear resistance in (W,Ti)C–(20â•›wtâ•›%)Co cermets.
25.6 CONCLUDING REMARKS The cermet processed from (W,Ti)C solid solution powders via three-stage sintering exhibits a combination of high hardness of 17â•›GPa and moderate fracture toughness of 7â•›MPaâ•›m1/2. A high toughness of more than 14â•›MPaâ•›m1/2 is recorded with (W,Ti) C–(20â•›wtâ•›%)Co cermets after one-stage sintering at 1500°C for 1 hour, independent of starting powders. Importantly, the cermet processed from solid solution powders, exhibits low wear depth (∼1–4â•›µm) and higher wear resistance ( wear rate ∼7╯×╯10−7– 18╯×╯10−7â•›mm3/Nâ•›m), compared with the reference cermet prepared from premixed WC/TiC powders (wear rateâ•›∼â•›14╯×╯10−7–22╯×╯10−7â•›mm3/Nâ•›m ). The analysis of wear data in the light of the phenomenological wear model study indicates that the measured wear rate cannot be explained entirely by classical fracture mechanics equations for brittle materials such as ceramics, pointing to the dominance of the metallic response.
â•… 419
REFERENCES
REFERENCES ╇ 1â•… A. Mukhopadhyay and B. Basu. Recent developments on WC-based bulk composites. J. Mater. Sci. 46 (2011), 571–589. ╇ 2â•… W. May. Phase decomposition and grain growth in (W,Ti)C-Co alloys. J. Mater. Sci. 6 (1971), 1209–1213. ╇ 3â•… L. E. Toth. Transition Metal Carbides and Nitrides. Academic Press, New York, 1971. ╇ 4â•… S. K. Bhaumik, G. S. Upadhyaya, and M. L. Vaidya. Microstructure and mechanical properties of WC-TiC-10Co and WC-TiN-Mo2C-Co(Ni) cemented carbides. Ceram. Int. 18 (1992), 327–330. ╇ 5â•… M. Tikkanen and H. Sipla. On the formation of WC-TiC solid solutions. Phy. Sint. 5(2) (1973), 67–77. ╇ 6â•… F. J. Arenas, A. Matos, M. Cabezas, C. Di Rauso, and C. Grigorescu. Densification, mechanical properties and wear behaviour of WC-VC-Co_Al hardmetals. Int. J. Refract. Met. Hard Mater. 19 (2001), 381–387. ╇ 7â•… R. Edwards. Cutting Tools. Institute of Materials, London, UK, 1993, 200. ╇ 8â•… G. Gille, B. Szesny, K. Dreyer, H. VandenBerg, J. Schmidt, T. Gestrich, and G. Leitner. Submicron and ultra fine grained hard metals for micro drills and metal cutting inserts. Int. J. Refract. Met. Hard Mater. 20 (2002), 3–22. ╇ 9â•… B. Montimer and J. K. Lancaster. Extending the life of aerospace dry bearings by the use of hard smooth counter faces. Wear 121 (1988), 289–305. 10â•… H. Tsujikawa, S. Maruoka, M. Koeda, S. Uryu, K. Funaba, K. Shbanuma, S. Katudate, N. Kanamori, E. Tada, and Y. Okhawa. Preliminary studies on large metal scaled gate valves for the international thermo nuclear experiment reactor. Vacuum 47 (1996), 639–646. 11â•… H. Engqvist, G. A. Botton, S. Ederyd, M. Phaneul, J. Fondelius, and N. Axen. Wear phenomena on WC-based face seal rings. Int. J. Refract. Met. Hard Mater. 18 (2000), 39–46. 12â•… F. Arenas, E. Ruiz, C. Di Rauso, and C. Grigorescu. Effect of TiC on the sintering and wear behavior of modified WC/Co hardmetals. Proceedings of the Tungsten and Refractory Metals International Conference 2000, M. S. Greenfield and J. J. Oakes (Eds.). MPIF, Annapolis, MA, 2000: 97–101. 13â•… M. H. Staia, I. J. Torres, C. Castillo, T. S. Sudarshan, J. Lesage, and D. Chicot. Tribological study of WC produced by plasma pressure compaction. Int. J. Refract. Met. Hard Mater. 24(1–2) (2006), 183–188. 14â•… J. Larsen-Basse. Binder extrusion in sliding wear of WC-Co alloys. Wear 105 (1998), 247–256. 15â•… E. M. Trent. Metal Cutting. Butterworth-Heinemann, Oxford, UK, 1991. 16â•… K. Jia and T. E. Fischer. Sliding wear of conventional and nanostructured cemented carbides. Wear 203–204 (1997), 310–313. 17â•… J. M. Guilemany, J. M. Miguel, S. Vizcaino, and F. Climent. Role of three-body abrasion wear in the sliding wear behaviour of WC-Co coatings obtained by thermal spraying. Surf. Coat. Technol. 140 (2001), 141–146. 18â•… S. Bodhak, B. Basu, T. Venkateshwaran, W. Jo, K.-H. Jung, and D.-Y. Kim. Mechanical and fretting wear behavior of novel (W,Ti)C–Co cermets. J. Am. Cer. Soc. 89(5) (2006), 1639–1651. 19â•… H. C. Kim, D. Y. Oh, J. Gujian, and I. J. Shon. Synthesis of WC and dense WC-5â•›wt% Co hard material by high-frequency induction heated combustion. Mater. Sci. Eng. A 368 (2004), 10–17. 20â•… S. Kangwantrakool and K. Shinohara. Sintering behaviour of mechanically coated WC-Co/TiCAl2O3 particles by high-speed rotational impact blending. Int. J. Refract. Met. Hard Mater. 21 (2003), 171–182. 21â•… A. Misra and I. Finnie. Some observations on two-body abrasive wear. Wear 68 (1981), 41–56. 22â•… J. F. Archard. Contact and rubbing of flat surfaces. J. Appl. Physics. 24 (1953), 981–988. 23â•… A. G. Evans and D. B. Marshall. Wear mechanisms in ceramics, in Fundamentals of Friction and Wear of Materials, D. A. Rigney (Ed.). Am. Soc. Met, Ohio, 1981, 439–452. 24â•… S. G. Roberts. Depth of cracks produced by abrasion of brittle materials. Scripta Mater. 40(1) (1999), 101–108.
SECTION
VI
FRICTION AND WEAR OF CERAMICS IN A CRYOGENIC ENVIRONMENT
CHAPTER
26
OVERVIEW: CRYOGENIC WEAR PROPERTIES OF MATERIALS It is widely recognized that tribological performance is a system-dependent property, primarily governed by material parameters (hardness, toughness, E-modulus of mating counterbodies), environmental conditions (ambient, lubrication), and operating variables (load, sliding speed, environment, etc.). Given such facts, it is important to note here that most of the studies evaluating the tribological properties of materials have concentrated on unlubricated experiments as well as experiments in water and other liquid media, including various oil lubricants and simulated body fluid, among others. In recent years, various traditional and advanced materials have received attention for use in tribological applications at cryogenic temperatures, and it is the purpose of this overview chapter to introduce this subject. This entire section of the book focuses on the wear behavior of bulk ceramic materials at cryogenic temperatures, especially in the presence of liquid nitrogen (LN2). In illustrating such issues, our recent research results on brittle ceramics (Al2O3, SiC, ZrO2) are presented. Hence, the primary goal of this section is to present the tribological data from the literature, summarize environment-specific wear mechanisms, and discuss the role of cryogenic fluids in causing the relative shift in wear mechanisms. This section would be of particular interest to scientists dealing with space applications who are involved in the selection and design of materials for tribological applications, for example high-speed cryogenic turbopumps. In this overview, a summary of research results obtained with metals (steel, high-purity Ti or Cu) are also presented and this will show how wear mechanisms of ductile metals can be characteristically different from those of brittle materials, as is discussed in other chapters of this section.
26.1 BACKGROUND Despite the fact that there was considerable development in cryogenic technology in last few decades, the interest of the materials science community in cryogenic Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
423
424â•…
CHAPTER 26â•… Overview: Cryogenic Wear Properties of Materials
technology has been restricted primarily to space technology. Apart from the existing engineering applications in cryogenic technology, the cryogenic environment has profound influence in technologically critical applications, such as the cryoturbopumps of the space shuttle main engine (SSME). The major driver of recent interest in understanding tribological properties of materials in a cryogenic environment stems from the reported failure of Stainless Steel 440C (martensitic grade) bearings in cryoturbopumps. Such failures result in frequent shutdown and operational difficulties of the SSME. In such background, research has been directed in two major directions: (1) assessment of some potential tribomaterials as well as understanding the underlying mechanisms of friction and material removal from sliding interfaces and (2) development of new bulk materials or surface coatings that can perform better under desired conditions. In cryogenic tribology, understanding the friction and wear mechanisms at subzero temperature has been the primary focus. This brings several elements of materials science into tribology, and the interest has been on evaluating materials for potential applications in liquid-fueled rocket engines. In the case of coating materials, the tribological properties are determined by surface composition and mechanical properties. As far as the applications of materials in cryogenic environments are concerned, one important factor that needs to be considered is the dramatic change in the mechanical properties of materials with decreasing temperature. From a fundamental point of view, materials for tribosystems in cryogenic applications have to sustain extreme low-temperature conditions and mechanical deformation at the surface. Materials for tribological applications span all material classes, which include metals, ceramics, and polymers and their composites.1 Importantly, these materials differ considerably in their physicomechanical properties.2 As mentioned earlier, the properties of both the mating materials influence the tribological response considerably. The properties that are relevant to the tribological applications can be reiterated here.3 For example, metallic materials are characterized by high tensile strength, fracture toughness, and thermal conductivity. Among all the material classes, ceramic materials show superior abrasion resistance and moderate tribochemical response. Outstanding properties of ceramics include high elastic moduli, compressive strength, and hardness. Less prominent properties of ceramics include higher modulus-dependent contact pressures, with the shift of the Hertzian shear stress maximum from bulk to the surface. Also, ceramics experience high frictioninduced temperature increase at the contact zone because of relatively low thermal conductivity (e.g., alumina, zirconia). More important, their major disadvantage is low fracture toughness. In contrast, because of their low interfacial adhesion energy, polymers such as polytetrafluoroethylene (PTFE) and polyethylene (PE) develop low friction conditions.4–6 A positive aspect of polymers is their low density, whereas the negative aspect is their poor thermal stability and low strength. Therefore, polymers can only be used under low Hertzian contact pressures (a few orders of megapascals) due to their premature failure by viscoelastic and plastic deformation modes at higher loads.7 Apart from the mechanical properties, the weight-to-density ratio of the material is an important criterion for use in moving components. Until now, SUS440Cgrade martensitic stainless steel has been the most widely used material for
26.2 Designing a High-Speed Cryogenic Wear Tester
â•… 425
ball-bearing applications, especially for liquid-fueled rocket engine turbopumps. However, such materials, if replaced with lower density substitutes, can evidently improve the bearing performance. It is to be noted that some of the structural ceramics, such as Al2O3, Y-TZP, SiC, and Si3N4, have moderate density (3–6â•›g/cm3) and high hardness (13–31â•›GPa), respectively. Clearly, high rotational speed as well as longer bearing life can be achieved with ceramic bearings.
26.2 DESIGNING A HIGH-SPEED CRYOGENIC WEAR TESTER The results discussed here were obtained with a designed high-speed ball-on-disk– type cryotribometer, which can operate in LN2 at high sliding speeds. Sliding occurs between a stationary ball and a high-speed rotating disk. The normal load and rotational speed both can be varied. The tangential frictional force along with normal load and speed can be simultaneously acquired dynamically during the test. Table 26.1 gives the overall technical specifications of the newly developed high-speed cryotribometer. Figure 26.1a shows the external view of the ball-on-disk cryotribometer and the details of the design drawing are shown in Figure 26.1b. The tribometer consists of the following major modules: spindle assembly unit; specimen housing with ball–pot assembly and bearing-mounting unit; mounting arrangement of test specimen and loading arrangement; drive motor and gear box; and data acquisition system and electronic controller. The spindle assembly consists of a flexible drive shaft that is secured into a stainless steel 440C shaft, which is held by a linear motion bearing (LMB) bushing and enclosed in an aluminum body. One end of the flexible shaft is connected to the test disk and the other end is connected to the drive motor. During the experiments, a constant level of LN2 is maintained inside the cryocan, wherein the disk–ball assembly is completely immersed.
TABLE 26.1â•… Specifications of the High-Speed Ball-on-Disk Cryotribometer11
Parameter Ball diameter Disk diameter Axial load Wear track diameter Test speed
Unit mm mm N mm
Minimum
Maximum
4 20 0 10
10 40 50 30
850
36,000
0
50
Frictional force
Revolutions per minute (rpm) N
Test environment
Normal ambient, LN2, liquid helium
Variables 4,6,8, and 10â•›mm 20–40â•›mm 1,2,5,10, and 20â•›N 10,15,20,25, and 30â•›mm 7 steps Any value in 0.1-N accuracy
426â•…
CHAPTER 26â•… Overview: Cryogenic Wear Properties of Materials
Flexible Shaft
Wear and Friction Monitor
Gear Box Motor Assembly
(a) Load Cell for Normal Loading Column Loaded Pan
Rope
Weights Motor Flexible Shaft
Detail-A Base Plate Emergency Switch Gear Box
Outer Cover Cryogenic Chamber Chamber Support
Structure
Motor
Anti-Vibration Pads Detail-A (b)
Figure 26.1â•… Overall external look at the cryotribometer (a) and the detailed design description of the ball-on-disk high-speed cryotribometer (b).11
The ball–pot assembly unit has a unique feature of holding a ball at four different radial positions on the disk specimen, and wear track radii of 7.5, 10, 12.5, and 15â•›mm can be obtained to facilitate different sliding speeds from 0.6 to 45â•›m/s. The linear sliding speed (v) can be computed as follows: v (m/s) = 2π × r × rotational speed (rpm ) / 60, where r is the track radius in meters.
26.3 Summary of Results Obtained with Ductile Metals
â•… 427
A loading pan is over-hung on a pulley with balancing dead weight, and a constant normal load can be maintained throughout the test. A maximum normal load up to 50â•›N can be applied. The drive motor used in the fabrication of the cryotribometer can operate at high speeds, up to 36,000â•›rpm, and the speed of the test spindle can be changed in 14 speed-steps in the range from 850 to 36,000â•›rpm. Also, the main motor cum gear unit can be easily isolated from the tribometer, which facilitates smooth operation with high sensitivity of frictional force measurement and related data acquisition. During sliding experiments, the frictional force is measured based on the high-speed torque measurement using a piezoelectric force transducer having high enough bandwidth and the data are continuously acquired using commercial software LABVIEW.
26.3 SUMMARY OF RESULTS OBTAINED WITH DUCTILE METALS Here, a summary of research results obtained with metals (steel, high-purity Ti or Cu) is presented. During dry unlubricated sliding of ductile metals, wear occurs either by mechanical damage of surface and subsurface because of severe localized plastic deformation or by oxidative wear due to the reaction of the mechanically damaged surface or products with the ambient environment.8 The latter process is significant in generating third-body material with a difference in chemistry compared with the base metal. Further, physicochemical behavior and, in particular, the reactivity of such third-body material influences further wear.9,10 However, it has been recognized that the total wear of a system is a result of simultaneous occurrence of the above processes, which makes the entire wear process a rather complex phenomenon. The results obtained with three different representative metallic materials illustrate some of the characteristic features of the sliding wear of ductile metals in a cryogenic environment.
26.3.1â•… Self-Mated Steel The sliding tests on self-mated SS304 stainless steel were carried out at a spindle speed of 17,100â•›rpm and at 850â•›rpm for self-mated 440C martensitic stainless steel. The results have been reported elsewhere.11 From Figure 26.2, it can be inferred that, at very high sliding speeds (22.4â•›m/s), the coefficient of friction (COF) of self-mated 304 stainless steel was found to be ∼0.1. Referring to Figure 26.3, for self-mated 440C martensitic stainless steel, COF increases from a very low value to around 0.15 in the run-in period (first 10 seconds); thereafter, the steady-state COF of 0.15 is maintained for the first 5 minutes of the test run. It can be noted that a COF of 0.5–0.6 is commonly reported in the literature for self-mated steel under unlubricated conditions.3 It should therefore be evident that the cryogenic environment has a significant influence on the reduction (∼50%) in COF of self-mated SS304 steel. Due to the severe plastic deformation, a highly uneven worn surface was observed on this wear track (Fig. 26.4). Also, the cracks originate from the central region of the wear track and propagate perpendicular to the sliding direction. The
428â•…
CHAPTER 26â•… Overview: Cryogenic Wear Properties of Materials
Coefficient of Friction
1.000
BALL: 10-mm dia.-AISI304 steel DISK: AISI-304 steel LOAD: 2 N (Peak Hertzian stress = 0.6 GPa) SLIDING SPEED: 22.4 m/s CRYOFLUID: LN2
0.800 0.600 0.400
µ = 0.1
0.200 0.000
0
25
50
75
100 125 150 175 200 225 250 275 Time (s)
Figure 26.2â•… Evolution of COF with sliding time of self-mated 304 stainless steel with sliding speed of 22.4â•›m/s.11
Coefficient of Friction
1.000
BALL: 8-mm dia.-SUS440C steel DISK: SUS-440C steel LOAD: 7 N (Peak Hertzian stress = 1.0 GPa) SLIDING SPEED: 0.89 m/s CRYOFLUID: LN2
0.800 0.600
µ = 0.15 to 0.278
0.400 0.200 0.000
0
50 100 150 200 250 300 350 400 450 500 550 600 Time (s)
Figure 26.3â•… Plot of COF versus sliding time of self-mated SUS440C steel at sliding speed of 0.89â•›m/s.11
occurrence of severe cracking-induced tribomechanical wear of the disk can be attributed to high sliding speed (22.4â•›m/s). The sliding test conditions as well as tribological data are summarized in Table 26.2. The wear of the ball was estimated by measuring the mass loss using a highprecision electronic balance having an accuracy of five digits. Looking at wear rate data for self-mated SS304 steel, the wear rate of the SS304 steel ball increases with increase in sliding speed, when all tests were conducted at identical load (2â•›N). For steels, abrasive wear is dominated by plastic deformation. However, at cryogenic temperature, the contribution of deformation would be much lower with dislocationactivated slip processes being limited at subzero temperature. Therefore, during high-speed abrasion, the fracture process (see Fig. 26.4) contributes largely to the high wear loss of material. Another observation is the presence of adhered iron oxide due to oxidative wear.
26.3 Summary of Results Obtained with Ductile Metals
â•… 429
Test conditions: self-mated 304 steel > in liquid nitrogen > hertz stress = 0.6 Gpa Sliding speed = 22.4 m/s > wear track width = 2.11 mm Wear rate of ball = 4.4 × 10–4 mm3/N m
50 µm
Figure 26.4â•… Optical micrograph showing the topographical features of the wear track on 304 steel disk worn against 304 steel ball.11 Test conditions: self-mated 304 steel in LN2; Hertzian stress, 0.6â•›GPa; sliding speed, 22.4â•›m/s; wear track width, 2.11â•›mm; and wear rate of ball, 4.4×10−4â•›mm3/Nâ•›m.
TABLE 26.2â•… Test Data of Experiments with High-Speed Ball-on-Disk Cryotribometer in LN2. The Sliding Tests on Self-Mated SS 304 Were Carried Out at 2-N Load, while That on Self-Mated 440C Steel at 7â•›N Load11
Tribosystem
Self-mated SS 304 steel Self-mated SS 304 steel Self-mated SS 304 steel Self-mated 440 steel
Hertz stress GPa (peak)
Sliding speed (m/s)
Sliding length, m (for ball)
COF (μ)
Wear rate of ball (mm3/Nâ•›m╯×╯10−4)
0.6
22.4
6160
0.1–0.12
4.4
0.6
17.9
5370
0.19–0.20
1.8
0.6
13.4
3819
0.19–0.21
1.25
498
0.15–0.28
0.2
1.0
0.89
In summary, with regard to self-mated SS304 steel, the frictional behavior clearly reflects that the high sliding speeds and cryogenic conditions have a strong influence on the steady-state value of μ. A very high wear rate of the ball of 4.4╯×╯10−4â•›mm3/Nâ•›m was measured at sliding speed of 22.4â•›m/s in LN2. The topographical investigation of the worn surface and the wear rate of the ball provide the evidence of severe plastic deformation and material transfer from the counterbody.
430â•…
CHAPTER 26â•… Overview: Cryogenic Wear Properties of Materials
26.3.2â•… Titanium/Steel Couple Among various metals, hexagonal closed packed (hcp) α-titanium (α-Ti) is known to have a favorable combination of physical and mechanical properties, such as lower density (∼4.5â•›g/cm3), high melting temperature (∼1668°C), high strength, good fatigue resistance, and superior corrosion resistance. The results of sliding wear experiments conducted on high-purity titanium (Ti) against bearing steel in a LN2 (boiling point, 77â•›K) environment are recently reported.12,13 High-purity Ti was thermomechanically processed and recrystallized to produce samples with grain size of 9, 17, and 37╛µm. These titanium samples were slid against a bearing-steel counterbody at 10-N load and at varying sliding speeds of 0.67, 1.11, and 4.19â•›m/s in an LN2 environment. A summary of important observations is given next. The steady-state COF for Ti/bearing-steel tribocouple varied in the narrow range of 0.28–0.38. The estimated wear rates for LN2-tested Ti samples were in the range of 10−3–10−4â•›mm3/Nâ•›m. The lowest wear rate (2.9╯×╯10−4â•›mm3/) was recorded in the fine-grained Ti (9-µm grain size) at the highest sliding speed (4.19â•›m/s). Scanning electron microscopy (SEM) results showed evidence of formation of adherent tribolayers that prevented further wearing of the metal (see Fig. 26.5). Substructure evaluation of worn surfaces revealed the formation of a dense array of deformation twins because of extensive plastic deformation, which often resulted in the subdivision of grains (see Fig. 26.5). Critical analysis of the worn surface topography reveals that the reduced wear rate was due to the formation of adherent and strain-hardened tribolayer. At a sliding speed of 4.19â•›m/s, the wear rate data appear to follow a Hall–Petch relationship versus grain size. However, this correlation was weaker with reduced sliding speed (see Fig. 26.6). To show various dominant wear mechanisms of Ti, a qualitative map was constructed in sliding speed–grain size space (see Fig. 26.7). Titanium is one of the reactive engineering materials and is prone to form titanium oxide layers even under moderate sliding conditions. In the case of testing in cryogenic liquid, there would be various influencing factors such as quenching of the friction surface temperature, blanketing of oxygen flow from the atmosphere to the friction surface, perhaps a degree of lubrication from the cryogenic liquid itself, and hardening of the test material and counterbody. The reduced test temperature also increases the hardness of the titanium test piece by 10 units of RB, as well as the counterbody. Against this backdrop, it is important to address the following questions: (1) Can LN2 influence frictional behavior to any noticeable extent? (2) Do the sliding tests at LN2 temperature influence the wear rate of titanium compared with the room-temperature (RT) sliding tests? and (3) Does deformation or fracture at subzero temperature contribute to the sliding wear properties of a Ti/steel tribosystem? In assessing these issues, two broad aspects, one based on flash temperature at the tribocontact and the other based on the stress-assisted deformation, are discussed in the following section. To better understand the wear mechanisms, a qualitative wear map for Ti/steel is provided in Figure 26.7. From a phenomenological point of view, high sliding speed causes an increase in the contact temperature (flash temperature), which often helps oxidation reactions to take place. A flash temperature of 340°C during ambient test conditions and 76°C
26.3 Summary of Results Obtained with Ductile Metals
50 µm
100 µm
(a)
20.00 kV
(b)
100 µm (c)
â•… 431
10.00 kV 800 ×
100 µm
(d)
Figure 26.5â•… (a) SEM image of worn intermediate grain size Ti-2 sample tested at 10-N load in LN2 at 4.19â•›m/s, showing severe delamination of the tribolayers and the wear debris. Representative images of coarse-grained Ti-3 sample tested at 4.19â•›m/s showing (b) as-worn surface after LN2 wear test (SEM image), (c) distribution of subsurface deformation twins in etched sample after LN2 wear test (optical microscopy image), and (d) twin boundaries as well as slip lines in etched sample after LN2 wear test (SEM image).12
during LN2 test has been predicted at a sliding speed of 1.34â•›m/s. Interestingly, the subzero flash temperature has only been predicted for LN2 test at a speed of 0.67â•›m/s. A maximum flash temperature of around 672°C was also estimated at the highest sliding speed (4.19â•›m/s) in an LN2 environment. Such a high flash temperature, however, is much lower than the temperature for phase transformation of α-Ti to body-centered cubic (bcc) β-Ti. Also, high flash temperatures of 369 and 480°C at sliding speeds of 2.51 and 3.35â•›m/s are estimated, respectively. At low sliding speeds (<1.5â•›m/s), lower flash temperature (<100°C) essentially lowers the chances of extensive oxide formation in tests in an LN2 environment. More important, the LN2 environment acts as a blanket and restricts oxygen flow to the friction surface. This would limit oxidative reactions at the friction surface.
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Wear resistance × 102 (mm3/N m)–1
40.0 35.0
0.67 m/s 1.11 m/s 4.19 m/s
30.0 25.0 20.0 15.0 10.0 5.0 0.0 0.16 0.18 0.20 0.22 0.24 0.26 0.28 0.30 0.32 0.34 Grain size–1/2 (mm–1/2)
Figure 26.6â•… Plot showing wear resistance of the titanium samples as a function of grain size tested against steel in an LN2 environment.12
cracking of tribolayer
smearing of tribolayer and delamination
abrasion and ploughing
smearing of tribolayer
grain size
sliding speed Figure 26.7â•… A qualitative wear mechanism map for titanium tested against steel counterbody at a load of 10â•›N and in the LN2 environment. Various dominant wear mechanisms are identified in the grain size–sliding speed space.12
The formation of tribolayers because of severe plastic deformation, as well as flow localization, in the form of elongated cavities in LN2, were also observed (Fig. 26.5). All such topographical development, in the light of the preceding discussion, can be attributed to large plastic shear strains that resulted from sliding wear tests. The contact pressure at the metal surface can be estimated by applying the concept of Hertzian contact stress field, that is, the mean (average) Hertzian contact pressure.3 Taking Poisson’s ratio equal to 0.3 for both titanium and the steel, the elastic modulus of the steel as 210â•›GPa, and that of titanium as 110â•›GPa, the mean contact stress and contact diameter at 10-N load have been computed to be 536.7â•›MPa and 77╛µm, respectively.
26.3 Summary of Results Obtained with Ductile Metals
â•… 433
Another experimental observation is the establishment of Ti/Ti contact, that is, the transfer of titanium from the test piece to the steel ball, leading to “galling,”14 which is a condition whereby excessive friction between high spots of mating parts results in localized welding with subsequent splitting and a further roughening of rubbing surfaces of one or both of the mating parts. The occurrence of galling is also reported in case of sliding of uncoated high-purity titanium against titanium alloys in unlubricated tests.15–17
26.3.3â•… Copper/Steel Sliding System Among various metallic materials, high-purity copper also has moderate mechanical strength and good fatigue resistance. In many engineering applications, tribological properties of copper dictate the performance of the component in service conditions. However, only a small number of studies have been conducted to evaluate the friction and wear behavior of high-purity copper. To demonstrate the tribological properties of ductile Cu in a cryogenic environment, results of the tribological study on work-hardened high-purity copper samples against bearing steel in LN2 environment at varying load (10, 15, and 20â•›N) and sliding speed (0.89, 1.11, and 1.34â•›m/s) conditions18 are summarized next. The friction and wear results are summarized in Figure 26.8. In cryogenic sliding experiments, a reduction in the steady-state COF (μF) was recorded with increasing load (10, 15, 20â•›N) at the highest sliding speed (1.34â•›m/s). High wear rate on the order of 10−4â•›mm3/Nâ•›m was measured, independent of the load or the sliding speed (see Fig. 26.8). On the basis of the experimental data and the characteristics of the worn surfaces, significant damage accumulation and plowing-induced material removal appear to have largely contributed to the wear losses. It is noteworthy that oxidative wear or mechanically mixed layer (transfer from steel counterbody) formation did not occur to any significant extent under the investigated sliding conditions. SEM analysis of the worn samples showed the dominance of plowing-induced severe mechanical damage (see Figs. 26.9 and 26.10). At lower load and sliding speed conditions, the copper surface was severely damaged by abrasion or plowing and delamination of the tribolayers. In contrast, at higher load and higher speed conditions, discontinuous delamination with large cavities, indicating severe deformation, were observed in the entire wear track. Also, no substantial evidence of oxidative wear or the formation of mechanically mixed tribolayer at the tribological interface was found. The characteristics of the sliding-induced severe deformation and damage accumulation can be discussed in the light of high contact stresses (more than 500â•›MPa) and thermal conditions (subzero flash temperature) at the tribological interface. Under the experimental conditions of copper sliding against steel in LN2 environment, the estimation on the basis of the Kong–Ashby model19 indicated that the flash temperatures at the mating interface can vary between −165°C and −180°C. Such low values of flash temperature are due to the high thermal conductivity of copper and also rapid heat transfer from the tribocontact to the LN2. Presumably, LN2 environment acts as a blanket and minimizes oxygen flow from the atmosphere to the friction surfaces and prevents the occurrence of tribochemical reactions.
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0.50
0.89 m/s 1.11 m/s 1.34 m/s
Average steady state (µf)
0.45 0.40 0.35 0.30 0.25 0.20 0.15 0.10 0.05 0.00
10
15 Load (N) (a)
8.0
0.89 m/s 1.11 m/s 1.34 m/s
7.0 Wear rate (10–4 mm3/N m)
20
6.0 5.0 4.0 3.0 2.0 1.0 0.0
10
15 Load (N) (b)
20
Figure 26.8â•… (a) Average steady-state μF of copper/steel sliding couple and (b) wear rate of copper, when slid against steel in LN2 environment.18
As far as the plasticity at the tribocontact of copper/steel is concerned, the load needed to initiate yield at the tribocontact is:3
WY = 21.17 R 2Y (Y / E*)2 ,
(26.1)
where Y is the yield strength of the test piece, E* is the effective elastic modulus, and R is the radius of the ball counterbody. The effective elastic modulus E* can be calculated by the following equation:
1/E* = (1 − ν12 ) /E1 + (1 − ν22 ) /E2
(26.2)
26.3 Summary of Results Obtained with Ductile Metals
15 N, 1.34 m/s
â•… 435
500.0 µm
(a)
15 N, 1.34 m/s
100.0 µm
(b)
Figure 26.9â•… Secondary electron (SE) images of worn surfaces of copper disk after sliding at 15-N load at different speeds against steel in LN2 environment: (a) damaged surface and local discontinuity in tribolayer and (b) appearance of large cavities throughout the wear track. Arrow indicates sliding direction.18
where E1 and E2 are the elastic modulus, and ν1 and ν2 are Poisson’s ratio of the mating solids. Taking Poisson’s ratios equal to 0.3 for both of the mating materials and the elastic modulus of the steel and copper as 210 and 130â•›GPa, respectively, the value of WY is around 2.8â•›N, implying that the plasticity is introduced to the copper samples at the tribocontact. For the case of plastic contacts with ductile metals as mating
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20 N, 1.11 m/s
200.0 µm
(a)
20 N, 1.34 m/s
200.0 µm
(b)
Figure 26.10â•… Worn surface morphology of copper disk that has undergone sliding at 20-N load at different speeds against steel in LN2 environment: (a) wider grooves and (b) the cavitation and delamination in tribolayer. Arrow indicates sliding direction.18
REFERENCES
â•… 437
solids, the maximum contact pressure (po) can be obtained by applying von Mises or Tresca’s condition:3
( po ) = 1.6Y ,
(26.3)
where Y is the yield strength in pure tension. Taking the yield strength of cold worked Cu as 345â•›MPa, the maximum contact pressure is estimated to be 552â•›MPa. Also, the maximum shear stress (τmax) for a circular contact is 0.31po for ν╯=╯0.3. For Cu/ steel sliding couple, the value of τmax at the subsurface region of worn Cu is estimated to be around 171.1â•›MPa. As a result, the observed plowing and smearing of the tribolayers are primarily the consequence of extensive deformation of the tribosurfaces.
26.4 SUMMARY In this chapter, the design of, as well as experimental results obtained with, a highspeed ball-on-disk cryotribometer is presented to demonstrate the challenging characterization of friction and wear behavior of different metallic materials immersed in cryogenic fluid (LN2) at higher sliding speed conditions. More description of the unique cryotribometer, as well as the use of this wear tester to evaluate the tribological properties of various coatings, has been reported elsewhere.20,21 The sliding wear results obtained with metallic materials demonstrate how the physical properties, such as thermal conductivity, are equally important parameters in the wear of metals at LN2 temperature (−196°C). In identical wear test conditions, high-purity Cu (wear rate ∼10−4â•›mm3/Nâ•›m) was found to be more wear resistant than the high-purity Ti (wear rate ∼10−3â•›mm3/Nâ•›m), though Cu and Ti have similar hardness. The reason for the high wear rate in Ti was found to be oxidative wear of Ti, which did not occur for Cu. Importantly, the research results demonstrate that materials with similar hardness subjected to identical LN2 wear test conditions can have significantly different wear rates and this is due to the difference in the flash temperatures, which depend on the thermal conductivity.
REFERENCES ╇ 1â•… D. K. Chaudhuri and R. Verma. Wear of liquid nitrogen–cooled 440C bearing steels in an oxygen environment. Engineered Materials for Advanced Friction and Wear Applications, Proceedings of an International Conference Gaithersburg, Maryland, USA, 1st–3rd March 1988, ASM International- 8801-006. ╇ 2â•… M. Nosaka, M. Kikuchi, M. Oke, and N. Kawai. Tribo-characteristics of cryogenic hybrid ceramic ball bearings for rocket turbopumps: Bearing wear and transfer film. Trib. Trans. 42(1) (1999), 106–115. ╇ 3â•… B. Bhushan. Principles and Applications of Tribology. Wiley-Interscience publication, John Wiley & Sons, Inc., New York, 1999, 447. ╇ 4â•… T. Gradt, T. Schneider, W. Hübner, and H. Börner. Friction and wear at low temperatures. Int. J. Hydrogen Energ. 23 (1998), 397–403. ╇ 5â•… I. V. Kragelsky, M. N. Dobychin, and V. S. Kombalow. Friction and Wear. Pergamon Press, Oxford, 1982, 224.
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╇ 6â•… Y. L. Ostrovskaya, T. P. Yukhno, G. D. Gamulya, Y. V. Vvedenskij, and V. I. Kuleba. Low temperature tribology at the B. Verkin Institute for Low Temperature Physics and Engineering (historical review). Trib. Int. 34 (2001), 265–276. ╇ 7â•… H. Czichos, D. Klaffke, E. Santner, and M. Woydt. Advances in tribology: The materials point of view. Wear 190 (1995), 155–161. ╇ 8â•… D. A. Rigney, L. H. Chen, and M. G. S. Naylor. Wear processes in sliding systems. Wear 100 (1984), 195. ╇ 9â•… D. A. Rigney. Transfer, mixing and associated chemical and mechanical processes during the sliding ductile material. Wear 245 (2000), 1–9. 10â•… M. Godet. Third bodies in tribology. Wear 136 (1990), 29–45. 11â•… B. Subramonian and B. Basu. Development of a high-speed cryogenic tribometer: Design concept and experimental results. Mater. Sci. Eng. A 415 (2006), 72–79. 12â•… A. Jain, J. Sarkar, B. V. Manoj Kumar. Harshavardhane, and B. Basu. Grain size–wear rate relationship for titanium in liquid nitrogen environment. Acta Mater. 58 (2010), 2313–2323. 13â•… B. Basu, J. Sarkar, and R. Mishra. Understanding friction and wear mechanisms of high-purity titanium against steel in liquid nitrogen temperature. Metall. Mater. Trans. 40A (2009), 472–480. 14â•… American Society for Testing and Materials. Standard G40-98b, Standard Terminology Relating to Wear and Erosion. ASTM, West Conshohocken, PA, 1999. 15â•… U. Wiklund and I. M. Hutchings. Investigation of surface treatments for galling protection of titanium alloys. Wear 251 (2001), 1034–1041. 16â•… M. Long and H. J. Rack. Friction and surface behavior of selected titanium alloys during reciprocatingsliding motion. Wear 249 (2001), 158–168. 17â•… L. Vitos, K. Larsson, B. Johansson, M. Hanson, and S. Hogmark. An atomistic approach to the initiation mechanism of galling. Comp. Mater. Sci. 37 (2006), 193–197. 18â•… B. Basu, B. V. Manoj Kumar, and P. Gilman. Sliding wear properties of high purity copper in cryogenic environment. J. Mater. Sci. 44 (2009), 2300–2309. 19â•… H. S. Kong and M. F. Ashby. Friction heating maps and their application. Mater. Res. Soc. Bull. 16(10) (1991), 41. 20â•… R. Khanna. Cryogenic wear of self-mated alumina and zirconia. MTech thesis, Indian Institute of Technology Kanpur, India, April 2006. 21â•… B. Subramonian, K. Kato, K. Adachi, and B. Basu. Experimental evaluation of friction and wear properties of solid lubricant coatings on SUS440C steel in liquid nitrogen. Tribol. Lett. 20(3–4) (2005), 263–272.
CHAPTER
27
CASE STUDY: SLIDING WEAR OF ALUMINA IN A CRYOGENIC ENVIRONMENT Structural ceramics are now recognized as potential candidate materials for hybrid bearings in rocket turbopumps. Friction- and wear-related failure is considered as one of the key factors in selecting the bearing materials for cryoturbopumps of the space shuttle main engine (SSME). To demonstrate fundamental understanding of the tribological properties of ceramics in cryogenic environment, this chapter reports the results of sliding wear study conducted on self-mated Al2O3, a model brittle ceramic material, in liquid nitrogen (LN2). The observed wear behavior is analyzed in terms of thermal heat dissipation and brittle fracture of alumina in LN2.
27.1 BACKGROUND There has been a growing interest in using advanced ceramic materials for spacerelated tribological applications, due to their unique combination of low density and high hardness. Hybrid ceramic bearings, consisting of ceramic balls enclosed in steel races, have exhibited the capability to extend bearing life under high-speed operation of liquid-fueled rocket engines.1,2 The lifetime of such rocket engines is usually limited to only a few minutes and suitable bearing material is expected to exhibit good tribological behavior (low friction and wear) within such a time frame. Nosaka et al.1 reported the self-lubricating performance of hybrid ceramic bearings (Si3N4), which are operated at speeds up to 50,000â•›rpm with a thrust load of 2840â•›N in LN2. In an LN2 environment, the hybrid ceramic bearings using Si3N4 balls are reported to have better wear resistance, provided by the thick transfer film consisting of FeF2/ iron oxide formed on the ceramic balls. Nosaka et al.2 also investigated the tribological performance of steel/ceramic (Si3N4) hybrid bearings in LN2 at speeds up to 120,000â•›rpm, and excellent performance in contrast to all-steel bearings was recorded. In cryogenic environment, tribosystems such as bearings, seals, and valves cannot be lubricated conventionally; therefore, the use of unlubricated bearing components is critical with respect to wear and frictional heat generation.3 It is known Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
439
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CHAPTER 27â•… Case Study: Sliding Wear of Alumina in A Cryogenic Environment
that space bearings typically operate in harsh conditions of high rotational speeds up to 50,000â•›rpm in pressurized cryogenic fluids, that is, LN2 (77â•›K), or liquid oxygen (LOX, 90â•›K), or liquid hydrogen (LH2, 20â•›K). The bearings also experience high transient axial loads with high contact stresses above the yield stress (∼1.8â•›GPa).4 Until now, SUS440C-grade martensitic stainless steel has been the most widely used material for the ball bearings used in the space shuttle, especially for the liquid-fueled rocket engine turbopumps. Study of the tribological behavior of ceramics in cryogenic environment, however, is limited (at least not widely reported in open literature), and the lack of information poses a barrier to their effective use as tribomaterials at cryogenic temperature. Among the structural ceramics, Al2O3 is considered for high-performance tribological applications, including extreme environments. For example, alumina has been widely used in seals and bearings due to its low density, rigid structure, and capability to retain smooth surfaces over a broad range of high stress loading conditions.5 Results of lubricated and unlubricated wear tests6–10 have shown the coefficient of friction (COF) to be in the range of 0.2–0.9, and wear rate was found to vary from 10−9 to 10−3â•›mm3/Nâ•›m for Al2O3. Such a large variation in the friction coefficient and wear rate is largely due to the strong influence of the variation in test conditions and test configurations. Xiong et al.11 found that the wear rate of self-mated Al2O3 decreases with an increase in grain size. Wang and Hsu12 investigated the wear transition of alumina and their study indicates that lubricants (paraffin oil and water) can effectively increase the wear transition load. It was postulated that the low friction of aluminabased ceramics at high humidity was predominantly due to the formation of aluminum hydroxides. Komvopoulos and Li13 reported the dominance of microplasticity, microfracture, and delamination, depending on the humidity and mating materials. Wang et al.14 reported wear mechanism maps for alumina at various operating conditions of dry sliding as well as lubricated sliding in nonreactive or reactive fluid. It was suggested that the cracks are due to the tensile stresses at the trailing end of the contact, which causes intergranular fracture and grain pullout in the severe wear regime.
27.2 MATERIALS AND EXPERIMENTS As described in Chapter 26, a high-speed ball-on-disk type of cryotribometer (DUCOM, India) was used to study friction and wear characteristics at sliding contacts in LN2 atmosphere at high sliding speeds. A detailed description of this equipment can be found elsewhere.15 Sliding occurs between a spherical alumina ball of 9.5-mm diameter and a hot-pressed alumina disk. Figure 27.1 shows the microstructure of thermally etched polycrystalline alumina ball and disk, revealing a typical bimodal grain-size distribution, most commonly observed for polycrystalline αalumina. The grain size of the α-alumina grains varies around 2.7 and 2.8â•›µm in alumina ball and disk sample, respectively. The sliding wear tests were conducted using normal loads of 2, 5, and 10â•›N with rotational speed of 2550â•›rpm for 10 minutes. Table 27.1 presents the selected
27.2 Materials and Experiments
â•… 441
10 µm (a)
10 µm (b)
Figure 27.1â•… SEM image of polished and thermally etched alumina ball (a) and alumina disk (b). Both of the microstructures exhibit a typical bimodal grain-size distribution.16
test parameters (load, sliding speed, and maximum Hertzian stress) for the cryogenic sliding wear tests conduced on self-mated alumina. It should be emphasized here that performing wear tests at high speeds (2550â•›rpm)—in LN2, at that—was a difficult task. This is particularly so when the tests were conducted on a large Al2O3 disk (35-mm diameter). Because of the inherent brittleness, no hole could be made
442â•…
CHAPTER 27â•… Case Study: Sliding Wear of Alumina in A Cryogenic Environment
TABLE 27.1â•… Summary of the Friction and Wear Data Obtained for Self-Mated Alumina under Varying Test Conditions16
Contact configuration
Operating environment
Ball-on-flat
LN2
Ball-on-flat
Water
Pin-on-disc
Distilled water, RT
Flat-on-flat
Dry air, 50% RH, RT
Ball-on-flat
Dry air, 40% RH, RT, 2 hours
Ball-on-flat
RT, 20–90% RH
Load (N)
Sliding speed (m/s)
Steady state COF
Wear rate (mm3/Nâ•›m)
2 5 10 49 24.5 9.8 49 24.5 9.8 49 24.5 9.8 10 40 10 20 20 40 40 40 80 80 80 10 10 30 50
2.67 3.33 3.67 0.18 0.18 0.18 0.54 0.54 0.54 1.18 1.18 1.18 0.1 0.1 4 2 4 1 2 4 1 2 4 0.1 0.9 0.2╯×╯10−2 0.2╯×╯10−2
0.17 0.13 0.15 0.28 0.28 0.28 0.29 0.24 0.28 0.28 0.25 0.28 0.2 0.3 0.65 0.67 0.98 1.12 1.18 0.90 1.10 0.90 0.64 0.48 0.48 0.75 0.45
1.57╯×╯10−5 4.94╯×╯10−5 7.76╯×╯10−5 0.2╯×╯10−7 0.3╯×╯10−7 1.2╯×╯10−7 1.8╯×╯10−8 4.5╯×╯10−8 6.1╯×╯10−8 1╯×╯10−8 0.3╯×╯10−7 0.3╯×╯10−7 0.5╯×╯10−7 0.3╯×╯10−7 0.4╯×╯10−5 0.5╯×╯10−5 1.9╯×╯10−5 1.4╯×╯10−5 1.8╯×╯10−5 1.9╯×╯10−5 1.4╯×╯10−3 1.3╯×╯10−3 4.4╯×╯10−3 10−8–10−6 ∼10−3 ∼10−6
through the disk and, instead, a specially designed disk holder was used. Overall, the cryogenic wear experiments were a difficult task—in particular, to carry out reproducible experiments. A more detailed description of the results presented in this chapter can be found elsewhere16 and, in the following sections, some important results as well as observations are summarized.
27.3 TRIBOLOGICAL PROPERTIES OF SELF-MATED ALUMINA As illustrated in Figure 27.2, frictional behavior of self-mated alumina in the cryogenic environment reveals severe initial fluctuation, with steady-state COF varying
â•… 443
27.3 Tribological Properties of Self-Mated Alumina
Coefficient of Friction (COF)
0.5
Load: 2–10 N, rotational speed: 2550 rpm, max. Hertzian stress: 0.57–0.99 GPa, LN2 environment
0.4
0.3
5N 2N
0.2
10 N 0.1
0.0 0
100
200
300
400
500
600
Sliding Time (s)
8
Maximum wear depth (mm)
8
6 6
4 4
Specific wear rate Maximum wear depth
2 2
4
6
8
2
Specific wear rate (× 10–5 mm3/N m)
Figure 27.2â•… Frictional behavior of self-mated alumina under loads of 2–10â•›N, at 2550â•›rpm rotational speed for 600 seconds in LN2 (maximum Hertzian stress, 0.57–0.99â•›GPa).16
10
Load (N)
Figure 27.3â•… Specific wear rate and maximum wear depth of disk after sliding in LN2 at loads of 2–10â•›N and rotational speed of 2550â•›rpm for 600 seconds.16
from 0.12 to 0.17. In the running-in period, a sudden increase in COF values can be attributed to sharp asperity interactions under maximum Hertzian contact stresses and, with continued sliding, COF starts decreasing as the real area of contact further increases due to flattening of asperities.7 From the variation of wear rate and maximum wear depth (Fig. 27.3), the wear rate of Al2O3 disk is found to be 1.57╯×╯10−5â•›mm3/Nâ•›m at 2-N load, followed by a continuous increase from 4.9╯×╯10−5 to 7.8╯×╯10−5â•›mm3/Nâ•›m at 5 and 10â•›N, respectively.
Wear depth (µm)
444â•… 0 –4 –8 –12 –16 –20 –24 –28 –32 –36 –40
CHAPTER 27â•… Case Study: Sliding Wear of Alumina in A Cryogenic Environment
10 N
5N
2N
Figure 27.4â•… Surface profiles taken across the wear track on the disk specimen after sliding for 600 seconds at loads of 2, 5, and 10â•›N with rotational speed of 2550â•›rpm in LN2.16
The wear rate of the ball exhibited almost a linear increase with the value going from 0.8╯×╯10−6â•›mm3/Nâ•›m (2-N load, 2â•›m/s sliding speed) to 1.8╯×╯10−5â•›mm3/Nâ•›m (5-N load, 2.67â•›m/s sliding speed) and finally to 2.4╯×╯10−5â•›mm3/Nâ•›m (10-N load and 3.3â•›m/s sliding speed). Two-dimensional surface profiles are useful to quantify wear depth. In Figure 27.4, the maximum wear depth at 2-N load is observed to be ∼3â•›µm, and follows a linear increase from ∼5 to ∼8.5â•›µm is recorded with an increase in load from 5 to 10â•›N, respectively. Both the wear rate and maximum wear depth data thus clearly indicate an increase in severity of wear damage with increasing load. Scanning electron microscopy (SEM) images (Fig. 27.5) reveal the extremely rough surfaces with indication of large abrasive grooves of 15–20â•›µm width and grain pullouts, after sliding a distance of 1200â•›m at 2-N load, 2â•›m/s sliding speed in LN2. Figure 27.6 shows the occurrence of intergranular cracking and various deep grooves of varied depth, which is a signature of differential wear. This differential wear depends on the preferential cleavage of the material, which can be due to orientation difference between grains. Figure 27.7 illustrates the smearing of the surfaces with parallel steps, which suggests mixed-mode fracture (both transgranular and intergranular). Broadly, the cleavage steps and intergranular cracking are the most characteristic features (Figs. 27.6–27.8). As far as the sliding wear damage of the Al2O3 ball is concerned, Figure 27.8 reveals the characteristic features of cleavage steps on various cleavage facets joining two parallel cleavage fractures, which are suggestive of transgranular fracture. The tortuous cracks are also observed along the grain boundaries, indicating intergranular fracture (see Fig. 27.8a). The tribological properties of self-mated alumina are compared in Table 27.1. Takadoum17 reported COF of 0.75–0.45 and wear rate of 10−3–10−6â•›mm3/Nâ•›m at 20% and at 90% relative humidity (RH), respectively, for self-mated Al2O3 at 30-N load. It was suggested that humidity favors the formation of a more protective layer of Al(OH)3, causing reduction in friction and wear of alumina. Ravikiran7 reported the unlubricated ball-on-disk wear data for self-mated alumina, when slid at a 10-N load.
â•… 445
27.3 Tribological Properties of Self-Mated Alumina
G
F
G
IG
10 µm (a)
LAG
1 µm (b)
Figure 27.5â•… (a) SEM images revealing the worn surfaces of alumina disk after sliding at 2â•›N, 2â•›m/s for 600 seconds (sliding distance of 1200â•›m) in LN2. Various characteristic wear features (G, grooves; IG, intergranular cracking; F, faceted and cleaved grains) are indicated in (a). The presence of LAG, that is, large abrasive grooves (15–20â•›µm wide) with debris entrapped on grooves, can be seen in (b), indicative of the pinning action of hard asperities.16
446â•…
CHAPTER 27â•… Case Study: Sliding Wear of Alumina in A Cryogenic Environment
G
1 µm (a)
IG G
G
IG
10 µm (b)
Figure 27.6â•… SEM images revealing the topographical features of worn surface on alumina disk after sliding at 5-N load, 2.67â•›m/s sliding speed for 600 seconds (sliding distance of 1602â•›m) in LN2: (a) grain boundary cracking (IG), grooving (G), and (b) differential grooving debris particle entrapment inside the deep abrasive grooves.16
â•… 447
27.3 Tribological Properties of Self-Mated Alumina
F
C F
IG
G
G F 10 µm (a)
SM
1 µm (b)
Figure 27.7â•… SEM images revealing the worn surfaces of alumina disk after sliding at 10-N load, 3.3â•›m/s sliding speed for 600 seconds (sliding distance of 1980â•›m): (a) intergranular cracking (IG), faceted and cleaved grains (F), and grooves (G), indicating grain pullout and mixed mode of fracture; (b) smeared (SM) surface with intergranular cracking. These features indicate extensive brittle fracture due to combined effect of cryogenic environment and high Hertzian contact stresses (0.99â•›GPa).16
448â•…
CHAPTER 27â•… Case Study: Sliding Wear of Alumina in A Cryogenic Environment
I
D
10 µm (a)
F
IG
CS
1 µm (b)
Figure 27.8â•… SEM images revealing the worn surfaces of alumina ball after sliding for 600 seconds at 10-N load with a sliding speed of 3.3â•›m/s (sliding distance of 1980â•›m). Various characteristic features (D, fine debris; F, faceted and cleaved grains; IG, intergranular cracking; CS, cleavage steps) are also indicated. Arrows marked along IG show the intergranular fracture. Cleavage steps and facets within the grains indicate the features of transgranular fracture.16
â•… 449
27.4 Genesis of Tribological Behavior in A Cryogenic Environment
Steady-state friction coefficient is reported to be 0.48, whereas wear rate varied in the range of 10−8–10−6â•›mm3/Nâ•›m. Andersson18 performed a number of water-lubricated pin-on-disk tests on sintered alumina (α-Al2O3) with loads of 10 and 40â•›N at a sliding speed of 0.1â•›m/s. While the COF showed a marginal increase from 0.2 to 0.3, the wear rate varied in the range of (0.5–0.7)╯×╯10−7â•›mm3/Nâ•›m with the increase in load from 10 to 40â•›N, respectively. It was reconfirmed that aluminum hydroxide layers can act as a boundary layer, which reduces the COF. Blomberg et al.19 reported the wear mechanisms of Al2O3 of varying loads (10–80â•›N) and high sliding speeds (1–4â•›m/s) under ambient conditions. Two predominant wear regimes of alumina were identified: “mild wear,” with extremely low specific wear rates of 10−8–10−7â•›mm3/Nâ•›m; and “severe wear,” with wear rates of the order of ∼10−3â•›mm3/Nâ•›m. From the wear mechanism point of view, microfracture at the edges of pores and microabrasive grooves dominated in the mild wear regime, whereas extensive surface fracture as well as plowing dominated the severe wear regime. Consequently, a large increase in friction coefficient up to 0.9 is measured at 80-N load. From the preceding discussion, it is evident that the wear rate, measured in ambient atmosphere and under similar sliding conditions19 (speed, 4â•›m/s; load, 10â•›N), is comparable with that measured in cryogenic environment (see Table 27.1). The dual effect of cryogenic environment appears, first, to result in the considerable reduction in friction coefficient (∼0.15) and, second, to cause very high wear (∼10−5â•›mm3/Nâ•›m).
27.4 GENESIS OF TRIBOLOGICAL BEHAVIOR IN A CRYOGENIC ENVIRONMENT 27.4.1â•…Friction of Self-Mated Alumina in LN2 It must be emphasized at this juncture that the steady-state friction coefficient in LN2 (∼0.13–0.18) is lower than that under the lubricated and unlubricated environments (0.2–0.9).6–10,13,17–19 The cryogenic fluid does not offer any lubrication effect, and hydrodynamic lubrication is ruled out due to low viscosity of all cryogenic fluids, with the exception of extremely turbulent LN2. The role of cryogenic fluids appears to be to serve as a coolant to dissipate the frictional heating and causing a reduction of the COF. The major reason for lower COF can be attributed to the presence of a considerable amount of finer wear debris particles (∼1â•›µm), which get trapped at the contacting surface between ball and disk. Due to the three-body abrasion phenomenon, the finer wear debris particles can act as roller elements, causing the reduction in COF. Another possible factor is the increased thermal conductivity of material at sliding contacts. The conductivity at high temperatures is limited by the mutual scattering of phonons, but at very low temperatures, phonon–phonon interactions can increase the thermal conductivity (k╯∼╯exp [−θ/αT]).20 The scattering of phonons at grain boundaries results in maximum thermal conductivity at very low temperatures, as reported by Kingery and coworkers.20 In fact, thermal conductivity of
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CHAPTER 27â•… Case Study: Sliding Wear of Alumina in A Cryogenic Environment
alumina follows a 10-fold increase with decrease in temperature to 77â•›K.20,21 Therefore, high thermal conductivity of alumina at LN2 temperature and continuous flushing by LN2 leads to more intense heat dissipation and consequently to reduction in friction coefficient.
27.4.2â•… Brittle Fracture and Wear of Self-Mated Alumina in LN2 Wear rate data (10−5â•›mm3/Nâ•›m) in LN2 environment indicate the severe wear regime (see Fig. 27.3). Also, large abrasive grooves suggest that the material removal occurs by abrasion due to hard and sharp asperity interactions in the cryogenic atmosphere. The wear behavior of ceramics is less understood in cryogenic environments. Unlike metals, most ceramics are inherently brittle; ceramics, however, possess high compressive strength, but low tensile strength. In a tribological scenario, tensile stresses cause the onset of wear in a contact, especially in the mild wear regime. The tensile stresses cause microfracture on the surface as well as in the subsurface region.22 It can be noted that, for single-crystal alumina, fracture energy at 298â•›K is less than that at LN2 temperature (77â•›K). Strong evidence of increased fracture stress of alumina at low temperatures has been proposed by Heuer.23 Heuer observed the rhombohedral twinning on the fracture surface of sapphire at −196°C, using the back-reflection Laue technique. Congleton and Petch24 reported the observation of thin flakes on fracture surfaces of polycrystalline alumina, mechanically fractured at 77â•›K. They also found the presence of rhombohedral twinning and enhanced dislocation density. Therefore, dislocations can glide in the preferred grain orientations under high Hertzian contact stresses in LN2 (see Table 27.1). As the crack propagates along the grain boundaries, a critical stage is reached and fractureinduced wear processes14 are accelerated. When fracture occurs in a single crystal, two new surfaces are formed with the supply of additional surface energy: 2γs (γs is the surface free energy), that is, the energy required for transgranular fracture. However, combining two fracture surfaces of different crystallographic orientation results in a grain boundary. Therefore, under ideal condition, the intergranular fracture energy should be equal to (2γs–γgb), so that grain boundaries (γgb is the grain boundary energy) are more easily fractured than the grains.25 To explain the initiation of fracture processes, Figure 27.9 presents a schematic representation of mixed-mode fracture (both intergranular and transgranular type), which simulates the fracture surface topography of as-worn surfaces on alumina in LN2. In polycrystalline ceramics, the grains have a range of orientations, with some grains being more favorably oriented for cleavage than others. When a crack passes from a grain of one orientation into a grain of a different orientation, the crack must reinitiate and propagate through the second grain to form the cleavage steps (Fig. 27.9). This explains the presence of extremely cleaved surfaces after sliding in an LN2 environment. In a cryogenic environment, brittle material such as alumina becomes very stiff as well as hard, which suggests that two hard bodies now abrade over each other more vigorously, leading to enhanced wear. With temperature of 77â•›K, the
â•… 451
27.4 Genesis of Tribological Behavior in A Cryogenic Environment
Intergranular Fracture 1
Transgranular Fracture 2
Cleavage Steps
Figure 27.9â•… Schematic of fracture processes revealing the intergranular fracture, transgranular fracture and cleavage steps. It can be noted that cleavage steps impinge on a grain boundary and reinitiate in the new crystal at positions 1 and 2 at the grain boundary.16
severity of contact between the mating alumina bodies is increased due to the pronounced effect of frictional surface hardening. Furthermore, Kenichi26 reported the increase in shear modulus of alumina at low temperatures (<200â•›K). One can therefore predict that alumina can only accommodate limited shear strain under the existing stress conditions. Also, higher elastic modulus and low Poisson’s ratio would result in a high contact stress. From the preceding discussion, it can be suggested that deleterious effect of both increase in load and LN2 environment can result in an increase in wear rate of alumina. For the ball-on-flat contact configuration, the frictional energy dissipation can be estimated using the following formula:27
q = µ Pa V (A a /A r ),
(27.1)
where: q ╯=╯heat generated per unit area (J/m2), μ ╯=╯COF, Pa╯=╯load per unit apparent area (P/Aa) (N/m2), P ╯=╯load (N), Aa ╯=╯apparent area of contact (m2), Ar ╯=╯real area of contact (m2), and V ╯=╯sliding speed (m/s). Based on this calculation, frictional heat dissipated for self-mated Al2O3 is around 96╯×╯1010â•›J/m2 at 10-N load and at sliding speed of 3.3â•›m/s in LN2. The fracture surface energies of Al2O3 are reported to vary from the range 6–24â•›J/m2 at 298–293â•›K to the range 16–32â•›J/m2 at 77â•›K, depending on the crystallographic orientations.24 This observation thus indicates that the brittle fracture of alumina at cryogenic temperature is easily facilitated.
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As a concluding note, the wear of self-mated alumina is an interplay between plastic flow and microfracture depending on the wear transition in water or oil medium9,13,22,25 at room temperature, whereas wear takes place predominantly by a brittle fracture process rather than plastic flow in LN2 environment.
27.5 CONCLUDING REMARKS One of the highlights of the results discussed in this chapter is that extremely low steady-state COF of 0.13–0.17 can be obtained for self-mated alumina during highspeed sliding (2550â•›rpm) in LN2. The high thermal conductivity of alumina at LN2 temperature and continuous flushing by LN2 result in intense heat dissipation from the contact zone, leading to reduction in the COF. Nevertheless, the wear rate is quite high and varies within the order of 10−5â•›mm3/Nâ•›m. Overall, the combined effect of increased surface hardening in LN2 and high Hertzian contact stresses (0.57– 0.99â•›GPa) appears to trigger the onset of brittle fracture in alumina. Material removal occurs by both transgranular and intergranular modes of brittle fracture. The effect of low temperature (LN2) is pronounced; first, in lowering the friction coefficient and second, in triggering the brittle fracture of alumina.
REFERENCES ╇ 1â•… M. Nosaka, M. Kikuchi, M. Oke, and N. Kawai. Tribo-characteristics of cryogenic hybrid ceramic ball bearings for rocket turbopumps: bearing wear and transfer film. Trib. Trans. 42(1) (1999), 106–115. ╇ 2â•… M. Nosaka, S. Takada, M. Kikuchi, T. Sudo, and M. Yoshida. Ultra-high speed performance of ball bearings and annular seals in liquid hydrogen at up to 3 million DN (120,000â•›rpm). Trib. Trans. 47 (2004), 43–53. ╇ 3â•… T. Gradt, T. Schneider, W. Hübner, and H. Börner. Friction and wear at low temperatures. Int. J. Hydrogen Energ. 23 (1998), 397–403. ╇ 4â•… I. V. Kragelsky, M. N. Dobychin, and V. S. Kombalow. Friction and Wear. Pergamon Press, Oxford, 1982, 224. ╇ 5â•… P. Andersson. A study on dry and poorly lubricated alumina journal bearings. Finnish J. Tribol. 1 (1993), 97. ╇ 6â•… D. Amutha Rani, Y. Yoshizawa, H. Hyuga, K. Hirao, and Y. Yamauchi. Tribological behavior of ceramic materials (Si3N4, SiC and Al2O3) in aqueous medium. J. Eur. Ceram. Soc. 24 (2004), 3279–3284. ╇ 7â•… A. Ravikiran. Influence of apparent pressure on wear behavior of self-mated alumina. J. Am. Cer. Soc. 83(5) (2000), 1302–1304. ╇ 8â•… J. D. O. Barceinas-Sánchez and W. M. Rainforth. On the role of plastic deformation during the mild wear of alumina. Acta Mater. 46(8) (1998), 6475–6483. ╇ 9â•… M.-C. Jeng and L.-Y. Yan. Environmental effects on wear behaviour of alumina. Wear 161 (1993), 111–119. 10â•… S. M. Hsu and M. C. Shen. Ceramic wear maps. Wear 200 (1996), 154–175. 11â•… F. Xiong, R. R. Manory, L. Ward, M. Terheci, and S. Lathabai. Effect of grain size and test configuration on the wear behavior of alumina. J. Am. Cer. Soc. 80(5) (1997), 1310–1312. 12â•… Y. Wang and S. M. Hsu. The effects of operating parameters and environment on the wear and wear transition of alumina. Wear 195 (1996), 90–99.
REFERENCES
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13â•… K. Komvopoulos and H. Li. The effect of tribofilm formation and humidity on the friction and wear properties of ceramic materials. J. Tribol. 114 (1992), 131–140. 14â•… Y. S. Wang, S. M. Hsu, and R. G. Munro. Ceramics wear maps: Alumina. Lubr. Eng. 47(1) (1991), 63–69. 15â•… B. Subramonian and B. Basu. Development of a high-speed cryogenic tribometer: Design concept and experimental results. Mater. Sci. Eng. A 415 (2006), 72–79. 16â•… R. Khanna and B. Basu. Low friction and severe wear of alumina in cryogenic environment: A first report. J. Mater. Res. 21(4) (2006), 832–843. 17â•… J. Takadoum. Tribological behavior of alumina sliding on several kinds of materials. Wear 170 (1993), 285–289. 18â•… P. Andersson. Water lubricated pin-on-disc tests with ceramics. Wear 154 (1992), 37–47. 19â•… A. Blomberg, M. Olsson, and S. Hogmark. Wear mechanisms and tribo mapping of Al2O3 and SiC in dry sliding. Wear 171 (1994), 77–89. 20â•… W. D. Kingery, H. K. Bowen, and D. R. Uhlmann. Introduction to Ceramics, 2nd ed. A WileyInterscience Publication, John Wiley & Sons, Inc., New York, 2004, Ch. 12, 615, 616, 619, 797. 21â•… D. A. Wigley. Materials for low temperature use. Eng. Des. Guides 26 (1978), 32–33. 22â•… S. M. Hsu and M. Shen. Wear prediction of ceramics. Wear 256 (2004), 867–878. 23â•… A. H. Heuer. Deformation twinning in corundum. Philo. Mag. 13 (1966), 379–393. 24â•… J. Congleton and N. J. Petch. Crack branching. Philo. Mag. 142 (1967), 749–760. 25â•… R. W. Davidge. Mechanical Behavior of Ceramics. Cambridge University Press, Cambridge, UK, 1979, 12. 26â•… O. Ken´ichi. Elastic modulus and the measurement of structural ceramics at cryogenic temperatures. Cryogenics 35 (1995), 735–737. 27â•… B. Bhushan. Principles and Applications of Tribology. Wiley-Interscience Publication, John Wiley & Sons, Inc., New York, 1999, 447.
CH A P T E R
28
CASE STUDY: SLIDING WEAR OF SELF-MATED TETRAGONAL ZIRCONIA CERAMICS IN LIQUID NITROGEN In this chapter, the friction and wear of self-mated ZrO2 ceramics in cryogenic environment are described. Using a specially designed high-speed cryotribometer, the obtained results are analyzed to answer some important issues: (1) can sliding in liquid nitrogen (LN2) reduce the coefficient of friction (COF) of self-mated ZrO2? and (2) does the tetragonal zirconia (t-ZrO2) transformation occur in a cryogenic environment and, if it occurs, how does it affect the fracture behavior?
28.1 INTRODUCTION In the case of metallic alloys (SUS440C and 340 AISI stainless steel), engineering polymers (polytetrafluoroethylene [PTFE], PA6), and coatings (self-lubricating coatings, solid lubricant, diamond-like carbon), a number of tribological studies at cryogenic temperatures have been reported.1–4 However, such studies on bulk ceramics at cryogenic temperature are rather limited.1 Among various structural ceramics, ZrO2 is one of the most important ceramic materials, because of its high fracture toughness and strength, as also discussed earlier in this book. Almost three decades back, CSIRO (Australia) scientists first reported that zirconia-based ceramics can be toughened considerably by the martensitic tetragonal-to-monoclinic (t→m) transformation in localized transformation zones around cracks.5 Phenomenologically, transformation toughening is accompanied by volume dilation (4–5%) during the t╯→╯m ZrO2 phase transition, which leads to crack tip shielding.6 The relationship between wear and toughness is an important issue in the field of ceramic tribology. Fischer and coworkers7 examined the wear behavior of
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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28.1 Introduction
â•… 455
TABLE 28.1â•… Summary of Some Available Tribological Data Obtained with Self-Mated Y-TZP Material under Different Environments16
Material
Configuration
Environment
3Y–TZP
Ball-on-disk
LN2 (77â•›K)
3Y–TZP, (0.3-µm grains)
Unidirectional rotation (pinon-disk)
TZP–TZP
Air (50%RH) Water Hexadecane Air
TZP–TZP
Air
TZP–TZP
Distilled water
TZP–TZP
Ethanol
3Y–TZP
Pin-on-plate
3Y–TZP
Pin on disc
Dry air RH: 50╯±â•¯5% Dry air
Tribological conditions
COF
Wear rate (mm3/Nâ•›m)
5â•›N, 1.1â•›m/s 10â•›N, 1.1â•›m/s 15â•›N, 1.1â•›m/s 0.001â•›m/s, 9.8â•›N
2.0â•›N, 0.1â•›m/s 3.4â•›N, 0.1â•›m/s 5.0â•›N, 0.1â•›m/s 2.0â•›N, 0.1â•›m/s 3.4â•›N, 0.1â•›m/s 5.0â•›N, 0.1â•›m/s 3.4â•›N, 0.1â•›m/s 5.0â•›N, 0.1â•›m/s 3.4â•›N, 0.1â•›m/s 5.0â•›N, 0.1â•›m/s 0.08â•›m/s, 3.7â•›N
0.55–0.75 0.45–0.50 0.30–0.35 0.35 0.6 0.11 0.58–0.63 0.51–0.59 0.50–0.56 0.58–0.63 0.56–0.62 0.50–0.56 0.29–0.37 0.35–0.40 0.10–0.15 0.10–0.14 0.62–0.69
3.8╯×╯10−4 5╯×╯10−5 7╯×╯10−6 0.5╯×╯10−6 0.15╯×╯10−6 0.3╯×╯10−7 2.6╯×╯10−5 2.1╯×╯10−5 1.9╯×╯10−5 3.2╯×╯10−5 2.1╯×╯10−5 3.5╯×╯10−5 5.8╯×╯10−7 1.4╯×╯10−7 3.2╯×╯10−7 2.8╯×╯10−7 1.6╯×╯10−5
19â•›N, 0.24â•›m/s 10â•›N, 0.016â•›m/s
0.12–0.44 0.18–0.39
8.9╯×╯10−4 2.4╯×╯10−7
self-mated zirconia, having varied fracture toughness. After conducting tribological experiments on stabilized ZrO2 with a broad spectrum of toughness variation (2.5– 11.6â•›MPaâ•›m1/2), it was observed that the wear rate decreased with increasing fracture toughness.7 The friction and wear of self-mated yttria-stabilized tetragonal zirconia polycrystalline (Y-TZP) ceramics under different triboenvironments (air, water, paraffin, hexadecane, oil, relative humidity [RH]) have been reported in a number of studies.8–15 A summary of related literature data is provided in Table 28.1. Lubricated and unlubricated wear tests,8–12 conducted on Y-TZP at room temperature (RT) revealed the friction coefficient to be in range of 0.1–0.9 and wear rate varied from 10−7 to 10−1â•›mm3/Nâ•›m. Such a large variation in the friction coefficient and wear rates is due to the strong influence of operating parameters, tribological environment, and test configurations on the wear mechanisms as well as wear transitions in the case of ZrO2. Under water-lubricated conditions, the influence of sliding speed on wear is more significantly pronounced than for dry sliding.12–14 Such an effect was initially thought to be due to frictional-heating-induced phase transformation during the wear of zirconia. Later, it was argued that, for Y-stabilized zirconia, the leaching of yttria can lead to destabilization of tetragonal zirconia in an aqueous environment. Such
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an effect causes more wear in water compared with unlubricated sliding in ambient environment.12–14 In a different work, Fischer et al.7 investigated the wear behavior of Y-TZP in various environments (air, water, hexadecane). The worn surface topographical features indicated that plastic deformation, accompanied by rolling of wear debris, was the predominant wear mode for Y-TZP at ambient conditions (air, 50% RH). Scott11 studied the unlubricated wear behavior of Y-TZP as a function of temperature. It was observed that under the critical sliding conditions, the stresses and temperatures caused the Y-TZP material to undergo phase transformation from tetragonal to monoclinic phase. This chapter summarizes the friction and wear behavior of a high-toughness ceramic, Y2O3-stabilized ZrO2, which is reported elsewhere.16
28.2 MATERIALS AND EXPERIMENTS A high-speed ball-on-disk–type cryotribometer (DUCOM, India) is used to study the tribological behavior in different environments (RT and LN2) over a range of high rotational speeds (850–36,000â•›rpm). A detailed description of this equipment is provided in the overview given in Chapter 26. The tribological tests were conducted using normal loads of 5, 10, and 15â•›N (corresponding Hertzian stresses, 0.49, 0.62, and 0.71â•›GPa, respectively) with rotational speed of 850â•›rpm (sliding speed, 1.1â•›m/s) for 300 seconds. Spherical Y-TZP balls of 10-mm diameter and hot-pressed zirconia disk samples were used; the average surface roughness (Ra) of a polished disk is around 0.05â•›µm. X-ray diffraction (XRD) analysis of polished surfaces indicated the predominant presence of t-ZrO2 in both ball and disk samples. Figures 28.1 and 28.2 show scanning electron microscopy (SEM) images, which reveal equiaxed grain morphology; the grain size (measured by line intercept method) of the tetragonal zirconia grains varies around 0.1–1.2â•›µm and 0.2–1.2â•›µm in zirconia ball and disk samples, respectively.
28.3 FRICTION OF SELF-MATED Y-TZP MATERIAL IN LN2 To illustrate frictional response, Figure 28.3 plots the COF of self-mated Y-TZP under varying load (5–15â•›N) with a sliding speed of 1.1â•›m/s in LN2. For the 5-N load test, COF fluctuates between 0.55 and 0.75, during the testing time of 300 seconds, and COF fluctuates between 0.45 and 0.50 for 10-N load. For 15-N load, COF varies between 0.28 and 0.38 for the major part of the test duration. Broadly, the frictional data indicate stable frictional behavior as well as lower steady-state COF at higher load (15â•›N) in LN2 environment for self-mated ZrO2. A number of researchers measured COF to be in the range of 0.1–0.9 with self-mated ZrO2 under lubricated and unlubricated environments (ambient, variable RH, water, paraffin oil, hexadecane).8–15 Recognizing the fact that the wide variation in sliding parameters would have direct influence on differences in the severity of
28.3 Friction of Self-Mated Y-TZP Material in LN2
â•… 457
Figure 28.1â•… SEM image of polished and thermally etched Y-TZP disk, revealing the equiaxed grain morphology of tetragonal zirconia.16
Figure 28.2â•… SEM image of polished and thermally etched Y-TZP ball, showing the equiaxed tetragonal zirconia grains.16
tribological response, it is, however, safe to state that LN2 does not have any beneficial influence on the reduction in COF of self-mated Y-TZP. In the case of self-mated ZrO2, the cryocooling effect has marginal effect in influencing COF, and this can be explained in the light of the following: ZrO2 does not exhibit any noticeable change in thermal conductivity with decrease in
458â•…
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SLIDING WEAR OF SELF-MATED TETRAGONAL ZIRCONIA CERAMICS
Load: 5–15N, sliding speed: 1.1 m/s, time: 300 s,
1.0
max. Hertzian stress: 0.49–0.71 GPa, LN2
Friction Coefficient
0.8 5N 0.6 10 N
0.4
15 N
0.2
0.0 0
50
100
150 Time (s)
200
250
300
Figure 28.3â•… Frictional behavior of self-mated Y-TZP during sliding in LN2 at load 5–15â•›N and sliding speed 1.1â•›m/s for 300 seconds; maximum Hertzian stress, 0.49–0.71â•›GPa.16
TABLE 28.2â•… The Maximum Flash Temperature as Well as Peclet Number (L) for the Investigated Tribological Conditions.16 The Details of the Procedure for Analytically Estimating the Flash Temperature Follows the Original Model, Proposed by Archard18
Load, W (N) 5 10 15
Sliding speed, V (m/s)
COF
Peclet no., L
Maximum flash temperature, θm (°C)
1.1
0.55–0.75 0.45–0.50 0.30–0.35
36.6 46.6 52.7
1023–1123 921–1053 717–883
temperature and the thermal conductivity of ZrO2 remains low (∼2â•›W/m/K), in a large temperature range down to LN2 temperature.17 The primary role of cryogenic environment therefore appears to be to keep the contacting surface cool, but the poor thermal conductivity of zirconia at 77â•›K does not allow faster heat dissipation from the contacting surface. Therefore, the cryocooling effect is nullified by the slower heat dissipation from the contacting surface and thereby does not cause reduction in COF. According to Archard,18 the velocity dependent thermal behavior at sliding contact can be analyzed in terms of the Peclet number (L), which can be uniquely defined by
L = V a/2κ,
(28.1)
where V is the sliding speed, a is the contact radius, and κ is the thermal diffusivity. For the set of experimental conditions, the value of L is calculated to be around 36–52 (see Table 28.2), depending on load variation. Following Archard’s model for
â•… 459
28.4 Cryogenic Wear of Zirconia
moderately high speeds (5╯<╯L╯<╯100), maximum flash temperature can be calculated using the following expression:
θm = 0.435 γ N L1 / 2 ,
(28.2)
where N and γ can be computed by the following mathematical equations: N = (µg/J ) (πpm /ρC p ) and γ = 1/ (1 + 0.87 L−1 / 2 ). In this set of expressions, μ is the COF, g is the acceleration due to gravity, J is the mechanical equivalent of heat, Q is the rate of heat supply from the contact area, W is the normal load, pm is the mean Hertizan contact pressure, ρ is density, and Cp is specific heat. Based on analytical calculations, the maximum flash temperature values are predicted to vary from 1023–1123°C at 5-N load to 717–883°C at 15-N load, at constant sliding speed of 1.1â•›m/s in cryogenic test conditions (see Table 28.2). It can therefore be expected that the boiling off of LN2 can easily occur as a result of extreme heat generation at the interface. The preceding observation is in contrast to the experimental results, obtained with self-mated Al2O3 under similar cryogenic sliding conditions (reported in the preceding chapter). It was observed that high thermal conductivity of alumina at cryogenic temperatures causes faster heat dissipation and reduced thermal effect lowering in COF.1 It has been experimentally observed that wear debris, when generated during sliding, were removed from the track during the sliding wear tests of Y-TZP tribocouple in LN2 (see Fig. 28.8). Therefore the possibility of third-body abrasion leading to lowering of COF can be neglected in the case of self-mated ZrO2.
28.4 CRYOGENIC WEAR OF ZIRCONIA Figure 28.4 shows the variation of wear rate and maximum wear depth with normal load. The wear rate of the disk was found to be 3.8╯×╯10−4â•›mm3/Nâ•›m 5-N load; with load increasing to 15â•›N, the wear rate decreased to 7╯×╯10−6â•›mm3/Nâ•›m. The wear rate of the ball increased from 2.3╯×╯10−4â•›mm3/Nâ•›m (5-N load) to 3.2╯×╯10−4â•›mm3/Nâ•›m (10-N load) and with a marginal increase to 3.5╯×╯10−4â•›mm3/Nâ•›m (15-N load). Clearly, the zirconia ball undergoes more severe wear than the disk specimen. At low temperatures (77â•›K), it is expected that tribocontact becomes severe due to increase in contact stresses (higher elastic modulus and Poisson’s ratio in LN2) and surface hardening in cryogenic environment. Therefore, the surface of zirconia specimens becomes very stiff as well as hard, which implies that two hard bodies can abrade over each other more vigorously leading to enhanced wear. Figure 28.5 illustrates the wear track on the zirconia disk after the tribological testing at the load of 10â•›N, sliding speed of 1.1â•›m/s run for a sliding time of 1200 seconds in cryogenic and ambient test environments, respectively. Overall, the wear track maintains a smooth circular nature under both testing conditions. Such observations also indicate that the concentricity of the track is maintained during the entire sliding test. Clearly, the track width after testing in RT is much larger than that after cryotesting, indicating much higher wear damage at ambient temperature. Some of
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SLIDING WEAR OF SELF-MATED TETRAGONAL ZIRCONIA CERAMICS
Maximum wear depth (mm)
10
Maximum wear depth Specific wear rate
50
8
40
6
30
4
20
2
10
0
Specific wear rate (× 10–5 mm3/Nm)
460â•…
0 4
6
8
10
12
14
16
Load (N) Figure 28.4â•… Specific wear rate and maximum wear depth of disk after sliding in LN2 at loads of 5–15â•›N and rotational speed of 850â•›rpm for 300 seconds.16
the characteristic features of the worn surfaces of zirconia ball and flat after cryogenic/ ambient sliding are shown in Figures 28.6–28.8.
28.5 CRYOGENIC SLIDING-INDUCED ZIRCONIA PHASE TRANSFORMATION An important observation is related to the presence of orthorhombic zirconia (oZrO2) after sliding in LN2 (Fig. 28.9c). All the characteristic peaks of o-ZrO2, in the present case, are located at 2θ╯=╯31.1°, 35.4°, 36.0° and 51.3°, 60.4°, 60.9°, 63.6°. Importantly, no trace of any major characteristic peaks of monoclinic zirconia (mZrO2) could be detected near, 2θ╯=╯28.2° (100% relative intensity) and 2θ╯=╯31.5° (68% relative intensity), irrespective of the sliding conditions. It can be commented that the combination of tensile stress field at the tribocontact and large temperature increase is not sufficient to initiate t-ZrO2 phase transformation in the present case. However, such a phase transformation was earlier reported during sliding under various operating conditions12–14 and was reported elsewhere as well.19,20 It is possible that the finer tetragonal grains of the investigated zirconia (see Fig. 28.1) are extremely resistant to t╯→╯m ZrO2 phase transformation, even under extreme sliding conditions. Instead, the transformation to o-ZrO2 is favored during cryogenic sliding conditions. It can be recalled here that, besides three widely reported ZrO2 polymorphs (cubic, tetragonal, and monoclinic), there have been reports21–25 of an additional metastable orthorhombic phase (o-ZrO2) forming within ZrO2 particles dispersed in partially stabilized zirconia (PSZ) samples, subjected to thermomechanical treatments. In the light of the earlier observations, the o-ZrO2 formation
â•… 461
28.5 Cryogenic Sliding-Induced Zirconia Phase Transformation
500 mm
(a)
500 mm
(b)
Figure 28.5â•… SEM images revealing the shape of the circular track formed on the disk after sliding in LN2 (a) and at RT (b). The sliding conditions include a load of 10â•›N and sliding speed of 1.1â•›m/s.16
during sliding (under the presence of high temperature and large Hertzian stress field), which can be visualized as multiple asperity interaction under a given thermomechanical stress field, is potentially facilitated. Based on earlier reports, it can be said additionally that the orthorhombic phase is the result of a stress-activated transformation from the tetragonal phase and forms as an intermediate transformation product, before its subsequent transformation to monoclinic symmetry.26 In view of this, the formation of the orthorhombic phase implies that it could be seen as an intermediate step in the t╯→╯m transformation and probably during longer-term sliding under similar sliding conditions to the present case, subsequent transformation of o╯→╯m-ZrO2 could be expected to take place.
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5N
Microcracks
Sliding direction
10 µm
(a) 10 N LC
10µm
C
Sp Sliding direction
25 µm
(b) 15 N
Sliding direction
25 µm
(c)
Figure 28.6â•… SEM images revealing the worn surfaces of zirconia ball after sliding at loads of 5–15â•›N with a sliding speed of 1.1â•›m/s for 300 seconds in LN2. Various topographical features include the cracks perpendicular to the sliding direction; C, chipping; Sp, spalling; intergranular cracking; LC, large cracks and fish-scale pattern.16
â•… 463
28.5 Cryogenic Sliding-Induced Zirconia Phase Transformation
5N
Sliding direction
10 µm
(a) 10 N
Sliding direction
Abrasive scratches
5 µm
(b) 15 N
Fish-scale pattern
Microcracks
Sliding direction
(c)
5 µm
Figure 28.7â•… SEM images revealing the worn surfaces of Y-TZP disk after sliding at loads of 5–15â•›N, with 1.1â•›m/s sliding speed for 300 seconds in LN2. Various characteristic wear features include chipping, large cracks, and spalling, as well as the fish-scale pattern indicating the microcracks.16
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~35 µm width
Sliding direction 25 µm (a)
Sliding direction
10 µm (b)
Figure 28.8â•… SEM images revealing the worn surfaces of Y-TZP disk after sliding at 10-N load with a sliding speed of 1.1â•›m/s for 1200 seconds at RT. The cracks on the worn surface are indicated by dotted arrows. While the signature of extensive deformation on rather deeper grooves can be noted in (a), the adherence of tribolayer is observed in (b).16
28.6 WEAR MECHANISMS OF ZIRCONIA IN LN2 Figure 28.6 reveals a worn surface on the zirconia ball after cryogenic sliding in LN2 at loads of 5–15â•›N and 1.1â•›m/s sliding speed for 300 seconds. The topographical features clearly reveal rough surfaces, characterized by mixed mode of fracture
â•… 465
o
t
t
t
t
(c)
(b)
t
t
t
t
t
Intensity (arbitrary units)
o
o
o
o
28.6 Wear Mechanisms of Zirconia in LN2
(a)
20
30
40
50
60
70
80
2q (degree) Figure 28.9â•… XRD spectra recorded with thermally etched and unworn Y-TZP disk specimen (a), worn surface of Y-TZP ball after sliding at 10-N load at RT (b), and in LN2 (77â•›K) (c).16 The symbols “t” and “o” indicate tetragonal (t-ZrO2) and orthorhombic (o-ZrO2) zirconia, respectively.
(intergranular and transgranular types). In Figure 28.6a, the cracking of thick tribolayer is the major observation, which suggests the typical mechanism of cleavage induced brittle fracture. Observing the characteristic microstructural features in Figure 28.6a and b, it is clear that the ball undergoes severe wear under the operating conditions. The worn surface topographical features indicate very rough surface with the extensive intergranular and transgranular cracking. In general, the damaged tribolayer, like the work hardened layer commonly observed in case of metal (after sliding tests), has been observed. Figure 28.7 reveals that the wear of Y-TZP material (flat) takes place predominantly by brittle fracture, as observed by chipping due to fracture, exhibiting fishscale-like pattern. Similar cracking pattern on worn surfaces of self-mated zirconia has been earlier reported with microcracks exhibiting geometric and periodic patterns. A critical observation of Figure 28.7 indicates that large cracks can form potentially due to coalescence of finer microcracks during fatigue wear process. This phenomenon is revealed in fish scale patterns, which exhibit a steplike pattern along
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the sliding direction. Figure 28.8 demonstrates the topographical features of the worn surface on Y-TZP disk after RT sliding at the most severe conditions (15-N load, 1.1â•›m/s sliding speed, 0.71â•›GPa maximum Hertzian stress). The worn track appears to be largely deformed with deep and wide lateral grooves (∼35-µm width). Broadly, the shear flow of worn material is observed at various locations of wear track. Besides deformation-induced wear damage, the brittle fragmented tribolayer with localized but limited cracking was also noted. The genesis of spalling in cryogenic environment is explained below. Spalling is considered to be a brittle mode of fracture that originates from nucleation of microcracks and, on repeated cycles of stresses (fatigue mechanism), the cracks grow further and finally join to generate wear debris particles. The formation of small cracks has been the result of poor thermal shock resistance. It is well known that the thermal shock resistance of a material is directly proportional to thermal conductivity and inversely proportional to elastic modulus. Due to high elastic modulus and poor thermal conductivity of zirconia at cryogenic temperature (77â•›K),27 its thermal shock resistance is even poorer. This explains our observation that the material tends to spall off to a greater extent in cryogenic sliding conditions (see Figs. 28.7 and 28.8). The worn surface topographical features reveal the extensive shear or plastic flow of the tetragonal phase after RT sliding (refer Fig. 28.8). Furthermore, high contact temperature appears to favor deformation of ZrO2 surface during RT sliding and, interestingly, the combination of high temperature and Hertzian contact stress does not promote the transformation from t-ZrO2 to m-ZrO2. This is presumably the cause for larger track width, as measured after RT sliding tests (see Fig. 28.6). As a concluding note, cryogenic wear experiments on self-mated ZrO2 at higher sliding speeds should be carried out in future to see whether the decrease in wear rate with increase in load can be noticed under high-speed sliding conditions. Also, the potential application at RT under high-speed sliding conditions can be investigated in future by carefully planning the experiments at varying sliding speeds (up to 5–6â•›m/s). Such planned experiments will establish the transition conditions from deformation-dominated wear mechanism to fracture-triggered wear. Another aspect of future research is the evaluation of the influence of toughness on wear of Y-TZPs in cryogenic conditions. Such influence on fretting wear of monolithic Y-TZP 28 and TZP(Y)–ZrB2 composites29 is reported elsewhere.
28.7 CONCLUDING REMARKS The observation of high COF and wear rate in cryogenic sliding conditions is attributed to poor thermal properties of zirconia as well as to large flash temperature increase. Under the selected operating conditions, tetragonal ZrO2 transforms to orthorhombic ZrO2 during sliding in LN2 environment. Spalling and microcrackinduced damage involving fatigue wear are observed to be the major wear mechanisms under the cryogenic test conditions, whereas severe deformation takes place during RT sliding. A characteristic fish-scale pattern of microcracks, oriented preferentially along the sliding directions, has been critically observed after cryogenic
â•… 467
REFERENCES
sliding conditions. Some deeper abrasive grooves with limited cracking of tribolayer are, however, noticed in experiments at ambient sliding conditions.
REFERENCES ╇ 1â•… T. Gradt, T. Schneider, W. Hübner, and H. Börner. Friction and wear at low temperatures. Int. J. Hydrogen Energ. 23 (1998), 397–403. ╇ 2â•… Y. L. Ostrovskaya, T. P. Yukhno, G. D. Gamulya, Y. V. Vvedenskij, and V. I. Kuleba. Low temperature tribology at the B. Verkin Institute for Low Temperature Physics and Engineering (historical review). Trib. Int. 34 (2001), 265–276. ╇ 3â•… W. Hübner. Phase transformations in austenitic stainless steels during low temperature tribological stressing. Trib. Int. 34 (2001), 231–236. ╇ 4â•… W. Hübner, T. Gradt, T. Schneider, and H. Börner. Tribological behaviour of materials at cryogenic temperatures. Wear 216 (1998), 150–159. ╇ 5â•… R. C. Garvie, R. H. Hannick, and R. T. Pascoe. Ceramic steel. Nature 258(25) (1975), 703. ╇ 6â•… B. Basu. Toughening of Y-stabilized tetragonal zirconia ceramics. Int. Mater. Rev. 50(4) (2005), 239–256. ╇ 7â•… T. E. Fischer, M. P. Anderson, and S. Jahanmir. Influence of fracture toughness on the wear resistance of yttria-doped zirconium oxide. J. Am. Cer. Soc. 72(2) (1989), 252–257. ╇ 8â•… Y. He, L. Winnubst, A. J. Burggraaf, and H. Verweij. Influence of porosity on friction and wear of tetragonal zirconia polycrystal. J. Am. Cer. Soc. 80(2) (1997), 377–380. ╇ 9â•… Y. He, L. Winnubst, A. J. Burggraaf, and H. Verweij. Grain-size dependence of sliding wear in tetragonal zirconia polycrystals. J. Am. Cer. Soc. 79(12) (1996), 3090–3099. 10â•… M. Woydt, J. Kadoori, K.-H. Habig, and H. Hausner. Unlubricated sliding behaviour of various zirconia-based ceramics. J. Eur. Ceram. Soc. 7(3) (1991), 135–145. 11â•… H. G. Scott. Friction and wear of zirconia at very low sliding speed, in Wear of Materials, K. C. Ludema (Ed.). American Society of Mechanical Engineers, New York, 1985, 8–12. 12â•… J. D. O. Boncenias-Sanchez and W. M. Rainforth. Transmission electron microscopy study of 3Y-TZP worn under dry and water- lubricated sliding conditions. J. Am. Cer. Soc. 82(6) (1999), 1483–1491. 13â•… W. M. Rainforth. The sliding wear of ceramics. Ceram. Int. 22 (1996), 365–372. 14â•… T. E. Fischer, M. P. Anderson, S. Jahanmir, and R. Salher. Friction and wear of tough and brittle zirconia in nitrogen, air, water, hexadecane, and hexadecane containing stearic acid. Wear 124 (1988), 133. 15â•… S. W. Lee, S. M. Hsu, and M. C. Shen. Ceramic wear maps: Zirconia. J. Am. Cer. Soc. 76(8) (1993), 1937–1947. 16â•… R. Khanna and B. Basu. Sliding wear properties of self-mated yttria-stabilised tetragonal Zirconia ceramics in cryogenic environment. J. Am. Cer. Soc. 90(8) (2007), 2525–2534. 17â•… W. D. Kingery, H. K. Bowen, and D. R. Uhlmann. Introduction to Ceramics, 2nd ed. A WileyInterscience Publication, John Wiley & Sons, Inc., New York, 2004, Ch. 12, 615. 18â•… J. F. Archard. The temperature of rubbing surfaces. Wear 2 (1958), 438–455. 19â•… S. Hori. Strength-toughness relations in sintered and iso-statically hot-pressed ZrO2-toughned Al2O3. J. Am. Cer. Soc. 69(3) (1986), 169–172. 20â•… M. Ruhle, N. Claussen, and A. H. Heuer. Transformation and microcrack toughening as complementary processes in ZrO2-toughened Al2O3. J. Am. Cer. Soc. 69(3) (1986), 195–197. 21â•… L. K. Lenz and A. H. Heuer. Stress-induced transformation during subcritical crack growth in partially stabilized zirconia. J. Am. Cer. Soc. 65(11) (1982), C-192–C-194. 22â•… L. H. Schoenlein. Microstructural studies of an Mg-PSZ. PhD thesis, Case Western Reserve University, Cleveland, Ohio, 1982. 23â•… L. H. Schoenlein and A. H. Heuer. Transformation zones in MgO-PSZ, in Fracture Mechanics of Ceramics, Vol. 6. R. C. Bradt, A. G. Evans, D. P. H. Hasselman, and F. F. Lange (Eds.). Plenum, New York, 1983, 309–325.
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24â•… A. H. Heuer, L. H. Schoenlein, and S. Farmer. New microstructural features in magnesia partiallystabilized zirconia (Mg-PSZ), in Science of Ceramics, Vol. 12. P. Vincenzini (Ed.). Ceramurgica s.r.l., Faenza, Italy, 1984, 257–266. 25â•… T. A. Bielicki, U. Dahmen, G. Thomas, and K. H. Westmacott. Transmission electron microscopy investigation of orthorhombic-type phases in stabilized zirconias, advances in ceramics, in Science and Technology of Zirconia III, Vol. 24. A. H. Huer and L. W. Gibbs (Eds.). American Ceramic Society, Columbus, OH, 1981, 485–491. 26â•… S. Veitch, M. Marmach, and M. V. Swain. Strength and toughness of Mg-PSZ and Y-TZP materials at cryogenic temperatures. Advanced Structural Ceramics, MRS Symposium Proceedings, edited by P. F. Becher, and S. Scmiya. Mater. Res. Symp. Proc. 78(97) (1987), 97–106. 27â•… Y. Ma and E. H. Kisi. In situ neutron diffraction study of a liquid nitrogen-quenched Mg-PSZ under load: A microcrack-dominated system. J. Am. Cer. Soc. 88(9) (2005), 2510–2514. 28â•… B. Basu, J. Vleugels, and O. Van Der Biest. Microstructure-toughness-wear relationship of tetragonal zirconia ceramics. J. Eur. Cer. Soc. 24(7) (2004), 2031–2040. 29â•… S. Das Bakshi, B. Basu, and S. K. Mishra. Fretting wear properties of sinter-HIPed ZrO2-ZrB2 composites. Composites A 37 (2006), 1652–1659.
CHAPTER
29
CASE STUDY: SLIDING WEAR OF SILICON CARBIDE IN A CRYOGENIC ENVIRONMENT This chapter discusses the results of experiments on the sliding wear properties of self-mated SiC in liquid nitrogen (LN2) over a broad spectrum of operating conditions with variation in either load or sliding speed. Limited tribochemical wear as well as grain boundary microfracture-induced damage mechanisms contribute to wear of self-mated SiC. A comparison with results obtained with self-mated Al2O3 or ZrO2 confirms the good tribological potential of self-mated SiC in LN2, in terms of exhibiting a better combination of coefficient of friction (COF) and wear rate.
29.1 INTRODUCTION In the context of tribological applications, silicon carbide (SiC) could be a useful bearing material because of a combination of properties, including low density (∼3.2â•›g/cm3), high hardness (>30â•›GPa), and good thermal conductivity (∼100â•›W/mâ•›K). SiC has two polymorphs: α-SiC, with a hexagonal structure; and β-SiC, with a cubic structure. Like other ceramics, it is relatively stable in corrosive environments.1 Apart from applications as bearings, SiC ceramics are considered to be potential candidates for important engineering applications, such as mechanical seals for water pumps, valves, nozzles, and cutting tools.1–3 The main disadvantage of SiC, however, is low fracture toughness (2.5–4â•›MPaâ•›m1/2).2,3 Another disadvantage is that SiC is difficult to sinter and requires high sintering temperature and use of a sinteraid.4–6 Typically, hot pressing at 1700°C or higher temperatures is necessary to obtain dense SiC ceramic. Nevertheless, SiC is attractive for tribological applications. Reviewing the existing literature,7–21 it has been noted that the majority of the tribological studies on SiC was carried out under dry lubricated conditions as well as in an aqueous environment.
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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Many nonoxide ceramics, such as SiC and Si3N4-based materials, are considered to be a good choice for hybrid-bearing material in cryoturbopumps of the space shuttle. Another important aspect related to tribology of nonoxide ceramics is the occurrence of tribochemical wear. The results of this chapter illustrate whether SiC ceramics are prone to oxidation at sliding contact in LN2; a more detailed description of the results are available elsewhere.22
29.2 MATERIALS AND EXPERIMENTS All the sliding wear tests were carried out on self-mated α-SiC, 8-mm diameter balls and 35-mm-diameter SiC disks. The disks were hot pressed at 1700°C in argon atmosphere with 0.5â•›wtâ•›% boron as sinter-additive. The density of both the ball and disk specimens was 3.15â•›g/cm3 (fully dense). X-ray diffraction (XRD) analysis of polished surfaces confirmed the predominant presence of α-SiC in both ball and disk samples. While the hardness of SiC both at room temperature (RT) and in LN2 is around 33.5â•›GPa, RT indentation toughness values were found to lie between 2.5 and 3â•›MPaâ•›m1/2. In the first set of sliding tests, the linear speed was varied up to 1.1â•›m/s at a constant load of 5â•›N (corresponding mean Hertzian stress, 0.82â•›GPa) and the test duration was 10â•›minutes, in both LN2 environment and ambient conditions (see Table 29.1). The second set of sliding tests were performed at a much higher linear speed of 3.3â•›m/s with varying load of 5, 10, and 15â•›N, and each test was conducted for 15â•›minutes in LN2. Table 29.2 summarizes the selected test parameters (load, sliding speed and maximum Hertzian stress) for the sliding wear tests, conduced on self-mated SiC material in LN2.
29.3 FRICTION AND WEAR PROPERTIES To illustrate frictional behavior, the continuously recorded frictional data are plotted for both the set of sliding tests in Figure 29.1. The steady-state COF data are summarized along with other results in Tables 29.1 and 29.2. For the first set of tests at
TABLE 29.1â•… The Experimental Parameters Used for the First Set of Sliding Tests with Varying Sliding Speeds at the Same Load Along with Important Results Obtained with Self-Mated SiC in LN2 and Ambient Conditions22
Load (N)
5
Sliding speed (m/s)
COF (LN2 test)
COF (RT test)
Maximum Hertzian stress (GPa)
Mean/ average Hertzian stress (GPa)
Flash temperature in LN2 tests (°C)
Flash temperature in RT tests (°C)
0.67 0.89 1.1
0.40 0.36 0.33
0.36 0.38 0.35
1.2
0.8
−13 47 105
202 261 316
29.3 Friction and Wear Properties
â•… 471
TABLE 29.2â•… The Experimental Parameters Used for the Second Set of Sliding Tests with Varying Load at the Same Sliding Speed (3.3â•›m/s) Along with Some Important Results Obtained with Self-Mated SiC in LN222
Load (N)
5 10 15
Maximum Hertzian contact stress (GPa)
Mean Hertzian contact stress (GPa)
COF (LN2 tests)
Wear rate (mm3/Nâ•›m)
Wear depth (µm)
Flash temperature (°C)
1.2 1.45 1.66
0.8 0.96 1.1
0.35 0.30 0.28
1.6×10−6 6×10−5 8.3×10−5
1.0 2.2 4.8
625 841 1009
5-N load, considerable fluctuation in frictional force is observed within the first 200â•›seconds during the running-in period in LN2 environment (see Fig. 29.1a). For ambient sliding tests, however, the steady-state COF varies in a narrow window of 0.35–0.38, depending on sliding speed (see Table 29.1). Overall, COF data reveals that the sliding speed does not have any significant influence on the frictional behavior of self-mated SiC at low load (5â•›N). A load-dependent frictional response is recorded in Figure 29.1b and a close observation of Figure 29.1b reveals that steady-state COF is reached only after 400â•›seconds. From Figure 29.1b, it is apparent that the time required to reach steadystate COF depends on the normal load and also corresponds to the time required to aggregate the wear debris into a protective layer, which is consistent with an earlier report.17 During the running-in period, a rather large fluctuation in COF in the window 0.2–0.5 is also recorded, depending on load. From the differences in frictional response, the frictional behavior of self-mated SiC in LN2 environment is found to be dependent more on load than on sliding speed. Also, a lower steady-state COF can be achieved at a higher sliding speed (3.3â•›m/s). As far as wear resistance is concerned, both the wear rate and wear depth appear to decrease with sliding speed in LN2 environment (Fig. 29.2a). The maximum wear depth is less than 0.5â•›µm in LN2 environment and it varies between 0.1 and 0.4â•›µm in ambient conditions. Despite high-speed sliding conditions, the variation of wear rate remains on the order of 10−6â•›mm3/Nâ•›m. However, wear depth varies in the range 1–5â•›µm. From Figure 29.2b, the severity of wear damage appears to increase with load at sliding speed 3.3â•›m/s. Two-dimensional (2D) profile measurements after sliding tests at 3.3â•›m/s reveal greater severity of wear compared with that at lower speeds of 1.1â•›m/s or less (see Fig. 29.3). Also, the smoothness of the 2D profiles can be attributed to attrition wear, which is defined as material removal at a submicroscopic level at higher loads and speeds.20 From the wear damage of the ball, the wear rate has been estimated to be on the order of 10−6â•›mm3/Nâ•›m; the total wear of self-mated SiC indicates good potential for tribological applications of SiC in subzero temperature, particularly with reference to the current use of stainless steel bearings or other competitive materials.23
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CHAPTER 29â•… SLIDING WEAR OF SILICON CARBIDE IN A CRYOGENIC ENVIRONMENT
0.8 Self-mated SiC LN2 environment Load: 5 N
Coefficient of friction (COF)
0.7
0.67 m/s 0.89 m/s 1.1 m/s
0.6 0.5 0.4 0.3 0.2 0.1 0.0
0
100
200
300 400 Sliding time (s) (a)
500
600
0.8 Self-mated SiC LN2 environment Sliding speed: 3.3 m/s
Coefficient of friction (COF)
0.7
5N 10 N 15 N
0.6 0.5 0.4 0.3 0.2 0.1 0.0 0
200
400 600 Sliding time (s) (b)
800
Figure 29.1â•… Frictional behavior of self-mated SiC for the first set of experiments with varying sliding speed under a load of 5â•›N load in LN2 (a), and at higher speed (3.3â•›m/s) in LN2 with different load (b).22
â•… 473
29.4 Thermal Aspect and Limited Tribochemical Wear
20
1.0 Self-mated SiC LN2 environment Load: 5 N
14
0.8 0.7
12
0.6
10
0.5
8
0.4
6
0.3
4
0.2
2
0.1
0
0.6
0.7
0.8 0.9 1.0 Sliding speed (m/s) (a)
1.1
10
20 SiC flat LN2 environment
18 Wear rate 10–6 mm3/N m
0.0
Wear rate Wear depth
9
16
8
14
7
12
6
10
5
8
4
6
3
4
2
2
1
0
4
5
6
7
8
9 10 11 12 13 Normal load (N) (b)
14
15 16
Wear depth (µm)
Wear rate 10–6 mm3/N m
16
0.9
Wear depth (µm)
18
Wear rate Wear depth
0
Figure 29.2â•… (a) Specific wear rate and maximum wear depth of SiC disk after sliding against SiC ball versus sliding speed under 5-N load for 600 seconds in LN2 and (b) versus normal load in LN2 at sliding speed of 3.3â•›m/s for 15 minutes.22
Some representative worn surface topographical features are presented in Figures 29.4–29.8.
29.4 THERMAL ASPECT AND LIMITED TRIBOCHEMICAL WEAR The severity of sliding damage significantly increased during sliding for longer duration (15â•›minutes) in LN2, at a sliding speed of 3.3â•›m/s. At a load of 5â•›N, a
Depth (µm)
0.50
0.67 m/s
0.00 –0.50 –1.0 0.0
0.21
Distance (mm) 0.42
0.62
Depth (µm)
1.00 0.50
0.89 m/s
0.00 –0.50 –1.00 0.00
0.27
Distance (mm) 0.55
0.83
Depth (µm)
1.00
1.1 m/s
0.50 0.00 –0.50 –1.00 0.00
0.25
Distance (mm) 0.50 (a)
0.75
Depth (µm)
2.0
5N
1.0 0.0 –1.0 –2.0 0.0
0.20
0.41 Distance (mm)
Depth (µm)
5.0
10 N
2.5 0.0 –2.5 –5.0 0.0 5.0
Depth (µm)
0.83
0.62
0.26
0.53
0.79 Distance (mm)
1.05
1.32
1.58 15 N
2.5 0.0 –2.5 –5.0 0.0
0.28
0.57
0.86 1.14 Distance (mm) (b)
1.42
1.71
2.0
Figure 29.3â•… 2D surface profiles, as acquired using laser surface profilometer, taken across the wear track on SiC disk specimen after sliding with varying linear speeds for 600 seconds at load of 5â•›N (a) as well as with varying load for 900 seconds at sliding speed 3.3â•›m/s (b), all after testing in LN2.22
474
â•… 475
29.4 Thermal Aspect and Limited Tribochemical Wear
(a)
(b)
Figure 29.4â•… Some selected SEM images demonstrating the wear surface characteristics of SiC disk (flat sample) after sliding against itself at 5-N load for 600 seconds in LN2 with different sliding speeds: 0.67â•›m/s (a) and 0.89â•›m/s (b). Single-pointed bold arrow indicates the sliding direction; dashed arrow indicates the microfracture-induced spalling region. Limited formation of tribochemical layer is shown by dash-dotted arrow.22
significant area of wear track is found to be covered by a rather thick layer of SiO2rich tribochemical film (see Fig. 29.7a) and a closer look at Figure 29.7a reveals parallel cracking. In Figure 29.7b, a representative scanning electron microscopy (SEM) image is presented to illustrate the sliding damage on the SiC disk after sliding at 10-N load for 15â•›minutes. A number of topographical features, indicated by dashed arrows, show evidence of severe grain boundary microfracture and grain removal, as well as signs of transgranular fracture (Fig. 29.7a,b). In Figure 29.8a, the finer details of an entrapped SiO2 layer as well as evidence of transgranular cracking on the worn surface (15-N load, with sliding speed 3.3â•›m/s) are presented. A typical illustration of crack deflection around an elongated SiC grain due to
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(a)
(b)
Figure 29.5â•… Representative SEM images illustrating the wear surface characteristics of SiC disk after sliding against itself at 5-N load with a sliding speed of 1.1â•›m/s for 600 seconds in LN2 (a) and worn surface on SiC disk (flat sample) after sliding against itself at 5-N load for 600 seconds in ambient conditions with a sliding speed of 0.89â•›m/s (b). Single-pointed bold arrow indicates the sliding direction; dashed arrow indicates the microfracture-induced spalling region. Limited formation of tribochemical layer is shown by dash-dotted arrow.22
interaction with the grain boundary (15-N load, sliding speed 3.3â•›m/s) is provided in Figure 29.8b. The predominant presence of SiC with characteristically sharp Raman bands appearing at around 767, 789 and 969â•›cm−1 is recorded from both the polished and worn ceramic surfaces (see Fig. 29.9).24 In addition, the broad Raman band with peak located around 1086â•›cm−1 acquired from the worn surface is due to the stretching vibration of Si–O bond.25 Additionally, such a broad peak reflects the amorphous nature of silica.
â•… 477
29.4 Thermal Aspect and Limited Tribochemical Wear
(a)
(b)
Figure 29.6â•… Selected SEM images revealing the worn surface features of SiC ball after sliding against itself at 5-N load for 600 seconds with a sliding speed of 0.67â•›m/s in two different environments: (a) ambient condition and (b) cryogenic condition. Single-pointed bold arrow indicates the sliding direction.22
One of the key findings of these results is the observation of limited tribochemical reaction and the formation of a discontinuous tribochemical layer, depending on operating parameters (load, sliding speed, environment). The thermal severity at sliding contact can explain this. As summarized in Table 29.1, a flash temperature of 316°C during tests under ambient conditions, and 105°C during LN2 tests, can be predicted at a sliding speed of 1.1â•›m/s. Interestingly, the subzero flash temperature has only been predicted for LN2 test at the lowest speed (0.67â•›m/s). From the flash temperature estimation, it can be said that, at low sliding speeds (<1.5â•›m/s), lower flash temperatures (<400°C) reduce the chances of oxide formation in LN2 environment. More important, the LN2 blanket restricts oxygen flow to the
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CHAPTER 29â•… SLIDING WEAR OF SILICON CARBIDE IN A CRYOGENIC ENVIRONMENT
(a)
(b)
Figure 29.7â•… Selected SEM images demonstrating the wear surface characteristics of SiC flat after sliding against itself for 900 seconds in LN2 environment with a sliding speed of 3.3â•›m/s at 5-N load (a) and at 10-N load (b). Single-pointed bold arrow indicates the sliding direction; dashed arrow indicates the microfracture-induced spalling region. Limited formation of tribochemical layer is shown by dashed arrow.22
friction surface and, consequently, limits occurrence of tribochemical reactions at the friction surface. Since the COF of self-mated SiC is much higher than 0.1, the possibility of any lubricating effect from LN2 can be ruled out. Rough worn surfaces of samples tested in LN2, as shown in Figures 29.4–29.8, correlate well with the absence of lubricating effect from LN2. In Table 29.2, high flash temperature (up to 1000°C) is predicted at sliding speed 3.3â•›m/s, and this explains the observation of extensive silica-rich layer formation on the wear track.
29.5 Tribomechanical Stress-Assisted Deformation and Damage
â•… 479
200 nm (a)
200 nm (b)
Figure 29.8â•… Representative SEM images revealing the worn surfaces of SiC flat after sliding against itself at 15-N load for 900 seconds in LN2 environment with a sliding speed of 3.3â•›m/s. Single-pointed bold arrow indicates the sliding direction; dashed arrow indicates the crack path, mostly around the grain boundary.22
29.5 TRIBOMECHANICAL STRESS-ASSISTED DEFORMATION AND DAMAGE Figure 29.4a illustrates the sliding damage on SiC disk after a test with sliding speed 0.67â•›m/s for 300â•›seconds at 5-N load in LN2. Overall, mild abrasive scratches are observed without any sign of tribochemical layer formation. Some of the abrasive grooves are found to be 10â•›µm wide (see inset of Fig. 29.4a). In contrast, a different wear behavior is observed after tests at a sliding speed of 0.89â•›m/s in an LN2
480â•…
CHAPTER 29â•… SLIDING WEAR OF SILICON CARBIDE IN A CRYOGENIC ENVIRONMENT
Figure 29.9â•… Raman spectra obtained with polished/unworn SiC disk as well as worn SiC surface after sliding at 10-N load with sliding speed of 3.3â•›m/s in LN2 (77â•›K).22
environment (see Fig. 29.4b). A closer observation of Figure 29.4b reveals limited tribochemical layer formation (appearing in dark contrast and indicated by dashdotted arrow) as well as microcracking and spalling. Such microcracked regions are commonly observed in several regions along the wear track. A SEM image of overall wear track (Fig. 29.5a) as well as the inset of Figure 29.5a reveals the characteristic presence of a sliding-induced damaged region. Other characteristic features include severe microfracture and grain pullout in a rough wear-groove, whose width is around 20–25â•›µm. Energy-dispersive x-ray spectrometry (EDS) analysis could not detect an O peak (not shown), thereby indicating the dominance of only tribomechanical wear. Similar observation has also been made after sliding tests with 0.89â•›m/s speed in ambient condition, as shown in Figure 29.5b. EDS analysis of dark contrast region, however, confirms the layer to be rich in silica (not shown). The occurrence of cracking and spalling is also observed. Some representative SEM images of the worn surface of the SiC ball are shown in Figure 29.6. The SiC ball undergoes wear damage to a similar extent in both ambient and LN2 environments. In both cases, the phenomena of microcracking and localized spalling are observed at various locations on the ball’s worn surface. The formation of a SiO2 layer is not observed to any noticeable extent. Taking Poisson’s ratio equal to 0.14 for SiC and the elastic modulus of the SiC as 310â•›GPa, the mean Hertzian contact stress and contact diameter, at 5-N load, can be estimated to be 0.77â•›GPa and 45.6â•›µm, respectively. The maximum Hertzian contact stress is 1.15â•›GPa. Considering the typical toughness of SiC as 2–4â•›MPaâ•›m1/2, such high stress level is expected to cause microcracking. Also, the wear behavior dominated by cracking-induced spalling is typical of nontransforming ceramics, such as SiC.
â•… 481
29.6 Comparison with Sliding Wear Properties of Oxide Ceramics
29.6 COMPARISON WITH SLIDING WEAR PROPERTIES OF OXIDE CERAMICS Next, the sliding results presented in the two preceding chapters are compared with those of self-mated SiC. Looking carefully at the data in Table 29.3, it is clear that, although all three sliding couples experience a similar level of maximum Hertzian contact stress, a difference in wear rate and maximum wear depth is measured. The difference in friction and wear rate needs to be explained in the light of the differences in hardness, E-modulus, Hertzian contact pressure, and thermal conductivity. As far as the hardness is concerned, SiC has the highest hardness, while ZrO2 has the lowest. Since plastic-deformation-induced damage does not dominate in the case of brittle ceramics, the difference in hardness cannot explain any difference in wear rate. Concerning the elastic property, SiC has the highest elastic modulus, while ZrO2 has the lowest modulus. Following Hertzian contact mechanics theory, one can realize that the higher the E-modulus, the lower is the contact damage dimension. This explains why SiC exhibits higher wear resistance than Al2O3 or ZrO2. The COF (μ) due to adhesion at elastic contacts can be expressed as26 µ = ατ a /E*(σ/β*),
(29.1)
where τa represents the average shear strength of the dry contact, E* is the composite or effective elastic modulus, σ is the composite standard deviation of surface heights, and β* is the composite correlation length. Since all three sliding couples are smoothly polished to similar roughness values (Ra╯∼╯0.1â•›µm), it is clear from this expression that self-mated SiC should have higher COF than self-mated Al2O3. Despite the same, we also notice higher COF of self-mated ZrO2 than self-mated SiC. Therefore, the influence of other parameters, for example thermal conductivity,
TABLE 29.3â•… A Comparison of the Sliding Wear Properties of Self-Mated SiC with Research Results Discussed in Earlier Chapters and Obtained with Other Self-Mated Structural Ceramics in LN2. All the Sliding Experiments Were Conducted on the Same Cryogenic Tribometer22
Tribosystem
Vickers hardness, GPa
E* (GPa)
Load(N)
Sliding speed (m/s)
Maximum Hertzian contact stress (GPa)
COF
Wear rate (mm3/Nâ•›m)
Self-mated Al2O3 Self-mated ZrO2 Self-mated SiC
20.1 (Hv500) 16.2 (Hv500) 33.5 (Hv500)
214.3
10
3.3
0.99
0.15
115.4
5
1.1
0.49
0.55–0.75
4╯×╯10−4
170.3
5 10
1.1 3.3
1.20 1.45
0.33 0.30
7╯×╯10−7 6╯×╯10−6
8.3╯×╯10−5
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should be considered. SiC has the highest thermal conductivity (λ), while ZrO2 has the lowest λ value. Following the discussion in an earlier chapter, it is likely that thermally induced brittle fracture should lead to enhanced wear in the case of selfmated ZrO2. This should be followed by Al2O3 and SiC. Although such a quantitative trend is observed in wear rate data, no such trend is followed in frictional properties. Relatively lower COF in the case of self-mated Al2O3 can be attributed to increased λ (10 times the RT value) in LN2 and, therefore, much lower flash temperature as well as easier dissipation of heat from contacting surfaces to heat sink. As far as the Hertzian contact stress (max.) is concerned, the wear rate of self-mated SiC, Al2O3, and ZrO2 are 6╯×╯10−6â•›mm3/Nâ•›m at 1.45â•›GPa pressure; 8.3╯×╯10−5â•›mm3/Nâ•›m at 0.99â•›GPa pressure, and 4╯×╯10−4â•›mm3/Nâ•›m at 0.49â•›GPa contact pressure, respectively. Such an inverse relationship in the case under discussion here confirms that Hertzian contact pressure is an important parameter that should be considered when selecting materials for given tribological applications. On the basis of the combination of COF and wear rate under similar operating conditions, the better tribological properties of self-mated SiC, compared with those of Al2O3 or ZrO2, can therefore be established.
29.7 CONCLUDING REMARKS Considerable fluctuation in frictional behavior during the running-in-period is recorded for self-mated SiC over a wide spectrum of sliding conditions. The steadystate COF is found to be dependent more on load than sliding speed in LN2. Although wear rate varies around 10−7–10−6â•›mm3/Nâ•›m, better wear resistance is measured under LN2 sliding conditions, independent of sliding speed, at 5-N load. A comparison with the earlier published results, obtained with self-mated alumina or zirconia using the same cryotribometer, establishes the better tribological potential of SiC in cryogenic sliding conditions. For self-mated SiC, grain boundary microfracture, as well as spalling, is identified as a dominant wear mechanism. Limited
Limited tribochemcial wear for metals/nonoxide ceramics (e.g., SiC)
Cryocooling effect and faster heat dissipation
Friction and wear in cryogenic environment
Brittle fracture enhancement ((e.g., Al2O3) and phase transformation (e.g. ZrO2) as well as unique cracking characteristics
Changes in mechanical property and surface hardening
Figure 29.10â•… Summary of various aspects of wear mechanisms and different associated factors observed in cryogenic wear study of various materials, as reported in this section.
â•… 483
REFERENCES
contribution from tribochemical wear is explained in terms of the estimated flash temperature at tribocontacts. Raman spectroscopy analysis confirms the formation of an amorphous silica-rich tribolayer, after sliding in LN2 environment. As a summary, Figure 29.10 illustrates the various wear mechanisms identified with three different structural ceramics (Al2O3, SiC, ZrO2). Although brittle fracture and characteristic cracking patterns are observed on worn oxide ceramics, interestingly, limited tribochemical wear is recorded with SiC.
REFERENCES ╇ 1â•… W. D. G. Böcker. Covalent high-performance ceramics. Adv. Mater. 4(3) (1992), 169–178. ╇ 2â•… N. P. Padture. In situ toughened silicon carbide. J. Am. Cer. Soc. 77(2) (1994), 519–523. ╇ 3â•… N. P. Padture and B. R. Lawn. Toughness properties of a silicon carbide with an in situ induced heterogeneous grain structure. J. Am. Cer. Soc. 77(10) (1994), 2518–2522. ╇ 4â•… C. Greskovich and J. H. Rosolowski. Sintering of covalent solids. J. Am. Cer. Soc. 59(7–8) (1976), 336–343. ╇ 5â•… F. F. Lange. Hot-pressing behavior of silicon carbide powders with additions of aluminium oxide. J. Mater. Sci. 10 (1975), 314–320. ╇ 6â•… K. Negita. Effective sintering aids for silicon carbide ceramics: Reactivities of silicon carbide with various additives. J. Am. Cer. Soc. 69(12) (1986), C-308–C-310. ╇ 7â•… K. H. Zum Ghar, R. Blattner, D.-H. Wang, and K. Pöhlmann. Micro- and macro tribological properties of SiC ceramics in sliding contact. Wear 250 (2001), 299–310. ╇ 8â•… T. E. Fischer, Z. Zhu, H. Kim, and D. S. Shin. Genesis and role of wear debris in sliding wear of ceramics. Wear 245 (2000), 53–60. ╇ 9â•… R. Wäsche and D. Klaffke. Wear of multiphase SiC based ceramic composites containing free carbon. Wear 249 (2001), 220–228. 10â•… S. M. Hsu and M. Shen. Wear prediction of ceramics. Wear 256 (2004), 867–878. 11â•… V. S. R. Murthy, H. Kobayashi, N. Tamari, S. Tsurekawa, T. Watanabe, and K. Kato. Effect of doping elements on the friction and wear properties of SiC in unlubricated sliding condition. Wear 257 (2004), 89–96. 12â•… J. F. Li, J. Q. Huang, S. H. Tan, Z. M. Cheng, and C. X. Ding. Tribological properties of silicon carbide under water-lubricated sliding. Wear 218 (1998), 167–171. 13â•… X. Dong, S. Jahanmir, and L. K. Ives. Wear transition diagram for silicon carbide. Tribol. Int. 28(8) (1995), 559–572. 14â•… S. M. Hsu and M. C. Shen. Ceramic wear maps. Wear 200 (1996), 154–175. 15â•… P. Andersson and P. Lintula. Load-carrying capability of water lubricated ceramic journal bearings. Tribol. Int. 27 (1994), 315–321. 16â•… P. Anderson and A. Blomberg. Instability in the tribochemical wear of silicon carbide in unlubricated sliding contacts. Wear 174 (1994), 1–7. 17â•… J. Takadoum, Z. Zsiga, and C. Roques-Carmes. Wear mechanism of silicon carbide: New observations. Wear 174 (1994), 239–242. 18â•… L. C. Erickson, A. Blomberg, S. Hogmark, and J. Bratthall. Tribological characterization of alumina and silicon carbide under lubricated sliding. Tribol. Int. 26(2) (1993), 83–92. 19â•… S. Sasaki. The effects of the surrounding atmosphere on the friction and wear of alumina, zirconia, silicon carbide and silicon nitride. Wear 134 (1989), 185–200. 20â•… R. Komanduri and M. C. Shaw. Wear of silicon carbide in high speed sliding. Wear 36 (1976), 363–371. 21â•… K. Sang and Z. Jin. Unlubricated friction of reaction-sintered silicon carbide and its composite with nickel. Wear 246 (2000), 34–39. 22â•… T. Kumar Guha and B. Basu. Microfracture and limited tribochemical wear of Silicon Carbide during high speed sliding in Cryogenic Environment. J. Am. Cer. Soc. 93(6) (2010), 1764–1773.
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23â•… B. Subramonian, K. Kato, K. Adachi, and B. Basu. Experimental evaluation of friction and wear properties of solid lubricant coatings on SUS440C steel in liquid nitrogen. Tribol. Lett. 20(3–4) (2005), 263–272. 24â•… A. Killer, K. G. Nickel, and Y. Gogotsi. Raman microscopy of nanocrystalline and amorphous phases in hardness indentations. J. Raman Spectrosc. 30 (1999), 939–946. 25â•… Y. Sasaki, Y. Nishina, M. Sato, and K. Okamura. Raman study of SiC fibers made from polycarbosilane. J. Mater. Sci. 22 (1987), 443–448. 26â•… B. Bhushan. Principles and Applications of Tribology. A Wiley-Interscience Publication, John Wiley & Sons, New York, 1999.
SECTION
VII
WATER-LUBRICATED WEAR OF CERAMICS
CHAPTER
30
FRICTION AND WEAR OF OXIDE CERAMICS IN AN AQUEOUS ENVIRONMENT 30.1 BACKGROUND The significant amount of research conducted on the “water lubrication” of ceramic materials has led to its use in various mechanical systems. In particular, silicon nitride and silicon carbide can provide very low wear and low friction under specific conditions.1–4 Due to the low viscosity of the water, the lubricating films of water are very thin, which means the typical lubrication regime for water-lubricated systems can be described as boundary lubrication. Despite this, silicon nitride and silicon carbide can experience, under specific conditions, very low wear. In fact a super-low friction of 0.002 and 0.0035 was measured for silicon nitride1,2 when lubricated with water. Several key mechanisms have been proposed for such excellent properties. Some researchers have suggested that the properties of the boundary film, due to tribochemical reactions, are responsible for such excellent properties;2,5,6 others have argued that the hydrodynamic lubrication promoted by the extremely smooth surfaces can explain such a low level of friction.3,7 The importance of a combination of tribochemical reactions, the presence of hard particles, and the formation of smooth surfaces was emphasized in another study.8 Although these mechanisms vary in terms of several details, it is obvious that the beneficial effect of silicon nitride and silicon carbide is due to some strong tribochemical effects in the water, and that the tribological performance is always very positive. Alumina can be a suitable material for use in water-lubricated bearings due to its low wear properties; however, the friction in this case is relatively high to moderate.9–11 The other important oxide ceramic, that is, zirconia, often suffers from high wear and friction in water-lubricated applications because of the hydrothermal tetragonal-to-monoclinic transformation.12,13 Therefore, the behavior of oxide ceramics in aqueous solutions tends to be less favorable. Nevertheless, the wear behavior of oxide ceramics in aqueous environments is less well understood and the results are often contradictory. As a result, in this chapter we look primarily at some of
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
487
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CHAPTER 30â•… Friction and Wear of Oxide Ceramics in an Aqueous Environment
these contradictions associated with oxide ceramics and describe some of the fundamental reasons for their behavior in different aqueous solutions. Despite the fact that, in the many years of research in this and related fields, it was mainly “water” that was used as the lubricating medium. Here we would like to introduce a broader concept, including aqueous solutions instead of just water in the contact, and discussing some of the fundamental behavior that depends on these conditions. It seems obvious that the results, especially when contradictions are observed, will depend on the “water” parameters, such as pH and surface charge, and their interdependence with the different chemical compositions of the ceramic materials. In the past, some extremely important effects were not recognized as critical parameters for the wear and friction properties during boundary lubrication in water. The pH of the solution, which is associated with the surface charge and the zetapotential (ZP), and which, in turn, affects the electrochemical and other properties of surfaces (both bulk and particles) in aqueous solutions, was not investigated systematically in order to analyze its effect on the wear mechanisms and tribological behavior in various regimes. However, we have found that the pH, which influences the surface charge in aqueous solutions, affects the wear of alumina and zirconia by an order of magnitude, and the friction by 2–3 times,13–16 and that this all occurs within the boundary-lubrication regime.
30.2 TRIBOLOGICAL BEHAVIOR OF ALUMINA IN AN AQUEOUS SOLUTION The materials used were alumina hemispherical pins with diameters of 16.6 mm and 5-mm-thick disks with diameters of 25â•›mm. The as-polished surfaces had an Ra roughness value of 0.05╯±â•¯0.014â•›µm for the pin surfaces and 0.03╯±â•¯0.005â•›µm for the disk surfaces. The measured microhardness of the pins and disks was about 21╯±â•¯0.6â•›GPa. The experiments were performed in a reciprocating sliding device with a 7-mm stroke. The frequency of the oscillation was 1â•›Hz. The tests were performed for 2 hours, giving a total sliding distance of 100â•›m, and a load of 50â•›N was applied via a stationary loading system to each contact. The flat specimens were fixed in a corrosion-resistant holder in the form of a tub, in which the aqueous solution with a selected pH value was subsequently poured to provide the waterlubrication bath. The lubricating media with controlled compositions were prepared by using distilled water and HCl, or NH4OH, to form either acid or alkaline aqueous solutions with various pH values (0.85, 4, 6, 8.5, 9.5, and 13). The worn disk surfaces were quantified using a stylus-tip profilometer and the wear mechanisms were studied using an electron microscope. Figure 30.1a shows the wear-volume results for the pins in the various aqueous solutions. The high wear-volume region was obtained for the most acidic and the most alkaline solutions, that is, at pH values of 0.85 and 13, respectively, and for the intermediate pH value of 8.5. In contrast, the wear volumes in the solutions with pH 4, 6, and 9.5 were much lower and their extents were approximately the same. The dimensionless wear coefficients for the selected conditions are presented in Table 30.1. The values of the wear coefficients for pH 4, 6, and 9.5 were in the range
30.2 Tribological Behavior of Alumina in an Aqueous Solution
â•… 489
Wear volume × 10–3 (mm3 )
7.0 6.0 5.0 4.0 3.0 2.0 1.0 0.0
0.85
4
6
pH
8.5
9.5
13
8.4
9.5
13
8
10
13
(a) 1 0.9 Coefficient of friction
0.8 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0
0.85
4
6
0.9
3.9
6
ph (b)
50
3
Wear loss × 10 (mm )
60
–3
40 30 20 10 0
pH
(c)
Coefficient of friction
1 0.8 0.6 0.4 0.2 0
0.9
3.9
6
pH
(d)
8
10
13
Figure 30.1â•… (a) Wear volume of alumina pins after 100â•›m of sliding at 50â•›N as a function of pH. (b) Steady-state coefficient of friction in alumina–alumina contacts as a function of pH. (c) Wear volume of zirconia pins after 100â•›m of sliding at 50â•›N as a function of pH. (d) Steady-state coefficient of friction as a function of pH.14,15
490â•…
CHAPTER 30â•… Friction and Wear of Oxide Ceramics in an Aqueous Environment
TABLE 30.1â•… Dimensionless Wear Coefficient of Alumina Pins at Different pH Values14
pH
Wear factor (/)
0.85 4 6 8.5 9.5 13
2.4╯×╯10−5╯±â•¯2.4╯×╯10−6 1.0╯×╯10−6╯±â•¯1.1╯×╯10−7 2.1╯×╯10−6╯±â•¯1.3╯×╯10−7 1.6╯×╯10−5╯±â•¯6.9╯×╯10−7 2.9╯×╯10−6╯±â•¯7.5╯×╯10−8 1.2╯×╯10−5╯±â•¯1.0╯×╯10−6
of 10−6, suggesting a mild wear regime, while for pH 0.85, 8.5 and 13 the wear coefficients were an order of magnitude higher, that is, in the range of 10−5, which is the intermediate region between severe and mild wear. The coefficient of friction (Fig. 30.1b) shows, in a similar way to the wear volume, two regions with comparable values. The low values (about 0.22) of the coefficient of friction were obtained in the acidic and alkaline conditions at pH 0.85 and 13, while a much higher friction (about 0.55) was measured for all the other conditions, that is, pH 4, 6, 8.5, and 9.5. The variation of the coefficient of friction in the low-friction region was extremely low; in the high-friction region, the variation was much higher, in particular for the solution with a pH of 8.5. It is interesting to note that the friction behavior is almost opposite to that of the wear behavior: in the region where the wear volume was high, the coefficient of friction was low and vice versa. The only exception to this is at pH 8.5, where the wear volume is high, while the coefficient of friction varies between values below 0.5 and above 0.6. Moreover, the scatter of values in the coefficient of friction is rather high, and so distinguishing between the different pH conditions in the high-friction region might be difficult. To determine the wear mechanisms, scanning electron microscopic (SEM) analysis was performed for the various pH conditions (Fig. 30.2). Two distinct features can be observed on the worn surfaces: the absence of any tribolayer at the surface and the presence of a tribolayer formed in the contact. The worn surfaces generated in the most acidic (pH 0.85) and the most alkaline (pH 13) environments look very smooth and are similar to the unworn surfaces. In contrast to the unworn surfaces, several pits can be observed on the worn surfaces (Fig. 30.2a,d). The size of the pits and their form (with sharp edges) suggest that these pits were formed by fracture and the delamination of grains from the surface. No signs of other types of mechanical wear or the formation of any tribochemical layer can be found on these surfaces. Figure 30.2b also shows worn surfaces that are completely covered with a distinctive tribochemical layer. The tribolayers can be clearly distinguished by several cracks that are formed during the drying of the hydrated layer after the waterlubricated experiments, which was carefully studied in one of our earlier investigations.17,18 An analysis of the edge of the wear scar revealed that the tribolayer formed
30.2 Tribological Behavior of Alumina in an Aqueous Solution
(a)
(b)
(c)
(d)
â•… 491
Figure 30.2â•… SEM images of worn alumina pin surfaces after 100â•›m of sliding at 50â•›N as a function of pH: (a) pH 0.85, (b) pH 4, (c) pH 8.5, and (d) pH 13.14
at pH 8.5 (Fig. 30.2d) was probably thicker than the others. This is consistent with this sample having a high wear volume (Fig. 30.1a), since the removal of the thicker layer most probably results in a higher wear loss. Moreover, the appearance of the tribolayers at pH 4, 6, and 9.5 also suggests that they are rather thin. In addition to their thickness, the appearance and morphology of these three tribolayers look very similar and different from the one at pH 8.5. This is consistent with the wear-volume results, where a much lower wear (and with a similar extent for each sample) compared with the wear at pH 8.5 was measured at pH 4, 6, and 9.5.
30.2.1â•… Electrochemical Properties and Wear Characterization of Self-Mated Alumina Our results show that, when the alumina ceramic is worn in aqueous solutions, the wear and friction behavior significantly depend on the pH value. The wear volume varied by as much as an order of magnitude (Fig. 30.1a), while the coefficient of friction varied by a factor of 3, that is, from approximately 0.2 to 0.6 (Fig. 30.2b). From our results it appears that the reason for such variations is the formation of a tribolayer and its properties. SEM analyses showed that in the most acidic or alkaline aqueous solutions, pH 0.85 and 13, tribolayers were not formed on the worn surface
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(Fig. 30.2a,f). In these cases, the amount of wear depends primarily on the solubility of the alumina in the water.19 Since the wear is mainly caused by the dissolution of alumina, no other externally induced mechanical wear was observed on the surfaces. The observed pits on the worn surface were most probably formed by the delamination of grains, caused by the loss of the surrounding bulk material that originally supported the grains. After the preferential dissolution of individual grains, which results in the formation of a step-like surface, particular grains that are no longer adequately supported by this partially dissolved neighboring material tend to delaminate and form the pits on the surface. In these two extreme regions the wear loss was very high. However, although the wear was too high to be described as mild, due to a lack of significant mechanical wear, the conditions were not severe either (Table 30.1). In the intermediate pH region, pH 4, 6, 8.5, and 9.5, tribolayers were always formed on the worn surfaces. However, the tribolayer at pH 8.5 (Fig. 30.2d) was found to be thicker than the layer formed at pH 4. Obviously, more wear debris was generated during the wear process, which later compacted into a coherent and thick tribolayer, as seen in Figure 30.2d. Due to its high thickness more material was removed from the surface than for the other pH values, where the tribolayers formed were much thinner (Fig. 30.2b). In addition, the wear volume at pH 8.5 was the second highest observed in our study. The reason for the different wear behavior of the tribolayer at pH 8.5 compared with the other investigated pH values most probably lies in the electrochemical properties of the alumina. The net surface charge on the alumina for a pH of approximately 8 is about 0 (isoelectric point [IEP]), which causes flocculation and, hence, increases the wear volume.13 At pH values lower and higher than around 8, the net surface charge (the ZP) is highly positive or negative, respectively, and this changes the tribolayer formation and the wear behavior. Indeed, from the SEM analyses, it is clear that the appearance of the tribolayer formed at pH 8.5, where the surface charge is about 0, is quite different from the other tribolayers, where the surface charge is positive (pH 4 and 6) or negative (pH 9.5). Depending on the different tribolayers formed in the contact, the obtained wear regime could be either mild or intermediate.
30.2.2â•… Surface Roughness and Frictional Behavior The friction behavior was quite different, and to some extent it contradicted the results of the wear-behavior tests. SEM images (Fig. 30.2) reveal the properties for the tribolayers formed at pH 4, 6, and 9.5, and similar properties for the wear surfaces formed at pH 0.85 and 13. Since the coefficient of friction was very different for these two groups of worn surfaces (see Fig. 30.1b), it seems reasonable to suggest that surface roughness could affect the frictional behavior. To investigate this possibility we measured the roughness after the experiments on the worn surfaces for all the conditions. As anticipated, the surface roughness was significantly lower for pH 0.85 and 13, while for all the conditions where the tribolayer formed the worn surface’s roughness was 5 to 10 times higher. Accordingly, we can confirm the assumption that the formation of the tribolayer results in an increased coefficient of friction in our experiments. Although the variation in the roughness between the
â•… 493
30.3 Tribological Behavior of Self-Mated Zirconia in an Aqueous Environment
different tribolayers is relatively high, and the coefficient of friction could be associated mainly with the formation of the tribolayer (or not), it can also be concluded that the coefficient of friction also follows the worn-surface roughness pattern in its simplistic model: that is, if the surface roughness increases, the coefficient of friction increases and vice-versa. This mechanism can even be found when observing only the friction-roughness behavior in the high-friction (intermediate pH) region. Nevertheless, it should be clearly pointed out that the variations in the coefficient of friction between the conditions where the tribolayers form (high-friction region) are within the scatter and could therefore be questionable, while the effect of the formation of the tribolayers itself is much more pronounced. However, a more detailed inspection of the results shows that the tribolayers at pH 8.5 were the smoothest and the coefficient of friction was the least reproducible, resulting in a relatively high scatter compared with the experiments at pH 4, 6 and 9.5. In some of our experiments this friction was the highest, while in others it was the lowest, and so within the limits of this investigation we cannot be certain of the reason for such behavior. Nevertheless, this is consistent with all the previous results, which suggested that the tribolayer formed at pH 8.5 had different properties than the other tribolayers. From these results, we can conclude that for water-lubricated alumina, there exist regions of pH that are favorable for material-removal applications (the wear volume is high) and for low-wear applications (the wear volume is low). In analogy to this, the coefficient of friction also varies significantly with pH, but the low and high coefficients of friction are obtained for the pH values, where the wear volume is low and high, respectively. The only exception to these rules is at pH 8.5, where the coefficient of friction and wear are both high. In conclusion, this means that the wear and friction behavior of water-lubricated alumina can be tailored for particular applications by using different pH values.
30.3 TRIBOLOGICAL BEHAVIOR OF SELF-MATED ZIRCONIA IN AN AQUEOUS ENVIRONMENT As was mentioned in Section 30.2, the pin-on-disk tests were conducted on selfmated 3â•›mol% yttria-stabilized ZrO2. The hardness of the disk and pin varied in the range 11–12â•›GPa, while the surface roughness was in the range 0.03–0.05â•›µm. The aqueous solutions used in the experiments were prepared with distilled water and HCl or NH4OH to obtain solutions with a pH of 0.9, 3.9, 6, 8, 10, and 13. The wear experiments were performed with a TE 77 reciprocating-sliding test machine at a constant frequency of 1â•›Hz and a stroke length of 6.8â•›mm. The disks were stationary in the test machine, while the counterbody balls were sliding with a reciprocating motion. The total sliding distance in each test was approximately 100â•›m, corresponding to 14,400 loading cycles. A load of 50â•›N was used; this resulted in an initial Hertzian contact pressure of 0.9â•›GPa. After the experiments, the ball and disk samples were also analyzed with a stylus-tip profilometer to assess the topographical characteristics of the worn surfaces. Selected samples from each test condition were sputter coated with gold and examined in the SEM.
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After each experiment the wear debris was collected by passing the water solution through a 1-µm filter. The wear debris on the filters was examined in the SEM after coating with a layer of gold. For selected conditions, the wear debris collected after the experiments from the water solutions was dried and then analyzed with x-ray powder diffraction using Cu Kα radiation. Figure 30.1c shows the wear volume of the ZrO2 pin samples as a function of pH. At pH 0.9, the wear volume is much lower—about an order of magnitude—than for all the other conditions. The corresponding dimensionless wear coefficient was about 5╯×╯10−5. In the region of pH between 3.9 and 13, the wear was significantly higher and varied by about 30–40% for the different aqueous solutions. The corresponding dimensionless wear coefficient was of the order of 10−4. In this region the highest wear was measured at pH 6 and the lowest at pH 10. The wear data suggest that in the most acidic environment the wear corresponds to an intermediate regime, close to mild wear, whereas for all the other pH values, the wear was severe. The coefficient of friction exhibited very similar behavior to the wear; that is, it had a lower value in the most acidic environment at pH 0.9, but at pH 3.9 it quickly reached a plateau with an almost constant value, where it remained for all the other pH values (see Fig. 30.1d). However, both regions had rather high coefficients of friction: about 0.45 for pH 0.9 and 0.75 for the higher pH values. Figure 30.3 presents SEM images of the tested worn surfaces for various pH values. In agreement with the wear and profilometric results, an obvious distinction between the surface at pH 0.9 and all the other surfaces can be seen. The surface at pH 0.9 was very smooth and had no specific features (see Fig. 30.3a). No wear debris or any other signs of any kind of macroscopic damage can be observed. A completely different appearance was found with all the other surfaces (see Fig. 30.3b,d). The surfaces are covered with layers, obviously consisting of a lot of wear debris, as can be seen at the fractured areas inside the wear-debris layer. However, this debris seems to be bonded together with an amorphous-like phase. Moreover, from the outermost surface of the debris layer in particular, which is at least locally very smooth, it appears that this debris was compacted and substantially smoothed at the interface, as seen from the top, thin tribochemical film with an apparently amorphous appearance (see Fig. 30.3b,d). These debris layers can be estimated (from the wornsurface profilometrical measurements) to have a thickness of about 2â•›µm, and are extensively delaminated and spalled. There is much fracture and deformation found on these layers. The severity of the deformation and spalling is consistent with the very high wear (Fig. 30.1c) measured in this pH region, that is, between 3.9 and 13. Moreover, a more obvious difference between the surfaces in the high-wear region can be found at lower magnification (Fig. 30.4a). It is clear that the surface at pH 6 (Fig. 30.4a) is much rougher than the surface at pH 4 and exhibits continuous spalling, which agrees also with the wear results (see Fig. 30.1c). Similar to the results presented previously, two different types of wear debris can be distinguished (see Fig. 30.5). In the very acidic aqueous solutions at pH 0.9, a very small amount of wear debris remained on the filter (see Fig. 30.5a). The debris particles were mostly separated from each other and were rather small, about 1–3â•›µm. However, at higher magnification some debris at this pH appears to have an amorphous-like form (see Fig. 30.5), as observed many times in experiments
â•… 495
30.3 Tribological Behavior of Self-Mated Zirconia in an Aqueous Environment
(a)
(b)
(c)
(d)
Figure 30.3â•… SEM images of zirconia worn surfaces after 100â•›m of sliding at 50â•›N as a function of pH: (a) pH 0.9, (b) pH 3.9, (c) pH 8, and (d) pH 13. Arrows indicate sliding direction.15
(a)
(b)
Figure 30.4â•… SEM images of worn zirconia surfaces after 100â•›m of sliding at 50â•›N at lower magnification showing, macroscopic topographical differences at (a) pH 6 and (b) pH 10.15
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(a)
(b)
Figure 30.5â•… SEM images of various distinctive types of zirconia wear debris collected after 100â•›m of sliding at 50â•›N shown as a function of different conditions: (a) pH 0.9 and (b) pH 6.15
involving water-lubricated ceramics.18,20,21 In addition, much of the wear debris was smaller than 1â•›µm (see Fig. 30.5), which is about the same or less than the original grain size of the zirconia used in these experiments, that is, about 0.3â•›µm. Completely different wear debris was found in the aqueous solutions of all the other pH values. Figure 30.5b presents a micrograph of the wear debris collected at pH 6; for all other conditions in the high-wear region, similar features can be observed. There was always a huge amount of wear debris in the collected water samples, and the pieces had many sizes and shapes: from less than 1â•›µm to more than 30â•›µm. Some of the wear debris appears as fractured pieces of bulk zirconia ceramic (see Fig. 30.5), but much of the collected debris is delaminated pieces of previously compacted wear-debris tribochemical layers, like that found during the SEM analyses of the worn surfaces (see Fig. 30.3). Evidence for the mechanism of formation of the compacted layers from fractured and/or tribochemically transformed wear debris through the applied mechanical loads and tribochemically assisted processes during sliding can be seen from the layer presented in Figure 30.5. The observed wear-debris layer is broken into smaller pieces of debris, many of them as small as the zirconia grains (around 0.3â•›µm) or smaller. The smoothness and the morphology of this layer suggest that at the sliding interface this debris layer was compacted into a dense and coherent surface top-layer and smoothed via a tribochemical wear mechanism. The process of debris-layer de-cohesion seen in Figure 30.5 is probably enhanced by the drying process during sample preparation17 for the weakly bonded pieces of debris with smaller tribochemically enhanced cohesive forces between them. A dedicated transmission electron microscopy (TEM) study15 revealed a quite similar size and shape for the finest wear debris generated under different conditions, even in both wear regions, that is, at pH 0.9 and pH 6. Their sizes varied from 50 to 500â•›nm, although these pieces of wear debris were agglomerates. The primary particles were relatively spherical, with sizes of around 3–5â•›nm. Moreover, x-ray diffraction (XRD) was used to further characterize the collected wear debris (Fig. 30.6). The results confirmed that in addition to the
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30.3 Tribological Behavior of Self-Mated Zirconia in an Aqueous Environment
200 pH6
Intensity (Count)
T M M
150
T M T
100
M
T
50 0 20
30
40
50
60
70
2 Theta
Figure 30.6â•… Wear debris collected in the aqueous solution after experiments at pH 6.0. The T-peaks correspond to the tetragonal phase; the M-peaks correspond to the monoclinic phase.15
tetragonal peaks (denoted as T), which were the only peaks present before the wear experiments, several new peaks were recorded (Fig. 30.6). These peaks correspond to a monoclinic zirconia phase (denoted as M), suggesting that part of the zirconia ceramic subjected to the tribological action transformed from the tetragonal to the monoclinic phase, as reported previously for tribological experiments involving zirconia in water22 or humid air.23 This transformation was reported as a hydrothermal phase transformation, associated with an increase in volume, causing large internal stresses and, as a consequence, stress-corrosion cracking and fracturing.5 The low intensity of the peaks in the XRD spectrum, shown in Fig. 30.6, suggest that part of the collected wear debris was also transformed in the amorphous phase (most probably Zr–OH compounds23), which could correspond to the observed amorphous-like appearance of the worn surfaces and layers of wear debris (see Figs. 30.3 and 30.5) and TEM results.15
30.3.1â•… Zirconia Transformation and Wear The results of the work presented here suggest that the wear of zirconia ceramics can be differentiated into two major wear regions, depending on the pH. The coefficient-of-friction behavior exactly coincides with these two regions. The first region is at a very low pH (i.e., 0.9), and the second at all other higher pH values (i.e., from 3.9 to 13). The surfaces at the low pH were very smooth and the wear debris was rarely found, because of its small size. Moreover, some of the wear debris appears to be amorphous. These characteristics are clear from the SEM images of the worn surfaces and the wear debris (see Fig. 30.3a), but also from the separate TEM analyses.15 The lowest coefficient of friction and the lowest wear (a wear factor in the range of 10−5) under the present conditions suggest that this kind of wear behavior corresponds to tribochemical wear. The observed tribochemical wear process can be explained by the high solubility of the zirconia at very low pH values.
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In contrast, the high-wear region from pH 3.9 to 13 is characteristic for distinctive thick tribochemical layers (based on profilometric analyses, assumed to be up to 2â•›µm), which consist of agglomerated wear debris that is compacted together under high stress at the interface due to repeated sliding and tribochemical reaction products. The suggested wear process therefore consists of several phases. During the initial sliding, the local spot-to-spot contact temperature24 at the surface asperities increases enough to promote the hydrothermal transformation of zirconia from the tetragonal to the monoclinic phase, as was determined with the XRD analyses (see Fig. 30.6). This transformation is associated with a volume increase,22 which causes stress-corrosion cracking and fracturing of the bulk material, resulting in the formation of wear debris and gradually leading to a wear-debris layer. Due to the locally high stresses at the “highest” wear-debris asperities, the debris is further fractured and/or deformed. At the same time, the tribochemical reaction products are formed by the dissolution and subsequent precipitation of the very small subgrain, 3- to 5-nm-sized particles.15 These newly formed particles, together with the larger bulk debris, were covered with a tribochemically transformed amorphous material (most probably Zr–OH compounds23), as observed with the SEM (see Figs. 30.3 and 30.5), and supported by the TEM and XRD results. This soft amorphous phase assists in the bonding together of the various debris (from a few nanometers to a few tens of micrometers), in particular at the “every-time” sliding surface, and forming a compacted debris layer (see Figs. 30.3 and 30.5) as a result of the mechanical stresses at the interface and the tribochemically assisted mechanism. The interface was, therefore, rather smooth, where it was not fractured, which is typical for a tribochemically dominated processes.5,17,25 However, due to the substantial influence of the high macroroughness, the low friction typical of the smooth tribochemical layers cannot be obtained. This layer is thus seen to form a quasi-static “transition” layer in which the material removed from the layer, either by gradual attrition or by sudden detachment, is continually replenished by freshly formed wear debris, causing a high wear factor, of the order of 10−4. Such a high debris-removal rate and the rough surfaces with predominant mechanical wear also caused a high coefficient of friction, despite the locally smooth tribochemical layer at the interface.
30.3.2â•… Electrochemical Aspect of Wear In addition to the previously discussed processes, another process needs to be introduced to fully explain the wear behavior in the high-wear region from pH 3.9 to 13. Zirconia ceramic has its IEP where the net surface charge is close to zero, that is, at pH 6.0.15 This condition exactly coincides with the pH where the highest wear was observed in our investigation. The increase or decrease in the ZP, that is, in the number of positively or negatively charged ions at the native and/or debris surface, leads to a change in the wear rate. The same wear mechanism was also observed earlier for alumina ceramics.19 In this case, the mechanism was confirmed by using anionic polyelectrolytes, which could shift the ZP without a change in the pH value. In these experiments, the highest wear consistently followed the shift in the IEP. The effect of the electrochemical potential on the tribological performance of iron and iron oxides in aqueous systems has been previously described26 in terms of
30.4 Concluding Remarks
â•… 499
modifying the normal force via electrostatic repulsion and the surface chemistry, and consequently, the shear strength of the surface films. However, the high normal load in our experiments, compared with the electrostatic forces, does not support the suggestion that the direct changes in the external load are due to a repulsive action. However, the surface chemistry is modified, and the resulting repulsive and/ or attractive forces change the nature of the surface layers, primarily affecting the strength of the debris agglomeration and compaction and thus controlling the coherency, the smoothness, and the thickness of the surface tribolayers. With water lubrication of alumina, these changes in the surface layers are more pronounced because of the mild wear,14 and thus the “wear sensitivity” is higher. On the other hand, due to the predominant effect of severe mechanical wear of the zirconia (fracture and debris detachment) caused by the hydrothermal transformation in the pH region from 3.9 to 13, the differences in the nature of the debris films due to electrochemical forces are smaller and relatively less pronounced than in the case of alumina ceramics, where an order-of-magnitude change was observed. Nevertheless, the absolute difference in the specific wear rate for zirconia, depending on the ZP, is still very high, about 3.5╯×╯10−6â•›mm3/Nm. Moreover, clear differences in terms of the wear, the coherency, and the appearance of the surface tribochemical layers were found to depend on the surface-charge (ZP) variation also in these experiments involving the water lubrication of zirconia, as seen in Figures 30.3 and 30.4.
30.4 CONCLUDING REMARKS The wear volume in water-lubricated alumina can vary by one order of magnitude, depending on the pH. High wear is obtained in very acidic and very alkaline aqueous environments, due to the solubility effect, and also around the IEP, due to a (net) zero-surface-charge effect. In very acidic and very alkaline aqueous solutions tribolayers do not form, and wear is primarily a result of alumina dissolution. In the intermediate pH region tribolayers always form; however, near the IEP (around pH 8) their properties are different. For water-lubricated alumina, pH regions were found where the wear loss and the coefficient of friction could be significantly reduced or increased. These findings suggest an increased potential for achieving the desired wear and friction behavior if the proper water-lubrication conditions are applied. In the case of water-lubricated zirconia, we have observed two wear and friction regions, which were dependent on the pH and varied by about an order of magnitude for the wear and by about a factor of 2 for the friction. In the low-wear region (pH 0.9) the dissolution of zirconia predominates and leads to smooth surfaces, a small amount of wear-debris generation, relatively low wear and a low coefficient of friction. In the high-wear region (pH 3.9 to 13) distinctive tribolayers are formed as a result of the hydrothermal transformation of zirconia, leading to fracture of the bulk material and the formation of a debris layer. At the same time, tribochemical reaction products are formed by dissolution and precipitation. The debris is compacted and bonded together by the tribochemical reaction products via mechanical stresses and tribochemically assisted mechanisms. The continuous process of debris-layer
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removal and freshly formed wear debris caused a very high wear rate, and the wear factor was on the order of 10−4. The very rough surfaces under these conditions also cause a high coefficient of friction. Despite the clear and important influence of electrochemical effects on the formation of the debris layer and the resulting wear mechanisms, the severe mechanical wear of zirconia (fracture and debris detachment due to hydrothermal transformation) was the main prevailing wear mechanism in the region of pH from 3.9 to 13. Therefore, the differences in the nature of the debris films due to the electrochemical forces in this region are smaller and relatively less pronounced in comparison with the differences between the solubility-dominated tribochemical wear in the low pH (0.9) region and mechanical wear mechanisms, which dominate in the whole high pH (3.9–13) region.
REFERENCES ╇ 1â•… L. Jordi, C. Iliev, and T. E. Fischer. Lubrication of silicon nitride and silicon carbide by water: Running in, wear and operation of sliding bearings. Tribol. Lett. 17(3) (2004), 367–376. ╇ 2â•… M. Chen, K. Kato, and K. Adachi. The comparisons of sliding speed and normal load effect on friction coefficients of self-mated Si3N4 and SiC under water lubrication. Tribol. Int. 35 (2002), 129–135. ╇ 3â•… H. Tomizawa and T. E. Fischer. Friction and wear of silicon nitride and silicon carbide in water. ASLE Trans. 30 (1987), 41–46. ╇ 4â•… S. Jahanmir and T. Fisher. Friction and wear of slilicon nitride lubricated by humid air, water, hexadecane and hexadecane + 0.5 percent stearic acid. Tribol. Trans. 32 (1989), 32–43. ╇ 5â•… J. Xu and K. Kato. Formation of tribochemical layer of ceramics sliding in water and its role for low friction. Wear 245 (2000), 61–75. ╇ 6â•… R. S. Gates and S. M. Hsu. Effect of selected chemical compounds on the lubrication of silicon nitride. Tribol. Trans. 34(3) (1991), 417–425. ╇ 7â•… V. A. Muratov, T. Luangvaranunt, and T. E. Fischer. The tribochemistry of silicon nitride: Effects of friction, temperature and sliding velocity. Tribol. Int. 31 (1998), 601–611. ╇ 8â•… R. S. Gates and S. M. Hsu. Tribochemistry between water and Si3N4 and SiC: Induction time analysis. Tribol. Lett. 17(3) (2004), 399–407. ╇ 9â•… M. G. Gee. The formation of aluminum hydroxide in the sliding wear of alumina. Wear 153 (1992), 201–207. 10â•… R. S. Gates, S. M. Hsu, and E. E. Klaus. Tribochemical mechanism of alumina in water. Tribol. Trans. 32 (1989), 357–363. 11â•… P. Andersson. Water-lubricated pin-on-disc tests with ceramics. Wear 154 (1992), 37–47. 12â•… T. E. Fischer, M. P. Anderson, S. Jahanmir, and R. Salher. Friction and wear of tough and brittle zirconia in nitrogen, air, water, and hexadecane containing stearic acid. Wear 124 (1988), 133–148. 13â•… S. Novak, G. Dražicˇ, and M. Kalin. Structural changes in ZrO2 ceramics during sliding under various environments. Wear 259 (2005), 562–568. 14â•… M. Kalin, S. Novak, and J. Vižintin. Wear and friction behaviour of alumina ceramics in aqueous solutions with different pH. Wear 254 (2003), 1141–1146. 15â•… M. Kalin, G. Dražicˇ, S. Novak, and J. Vižintin. Wear mechanisms associated with the lubrication of zirconia ceramics in various aqueous solutions. J. Eur. Ceram. Soc. 26 (2006), 223–232. 16â•… S. Novak and M. Kalin. The effect of pH on water-lubricated alumina and zirconia ceramics. Tribol. Lett. 17(4) (2004), 727–732. 17â•… M. Kalin, B. Hockey, and S. Jahanmir. Wear of hydroxyapatite sliding against glass-infiltrated alumina. J. Mater. Res. 18(1) (2003), 27–36.
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18â•… M. Kalin, S. Jahanmir, and G. Dražicˇ. Wear mechanisms of slip-cast glass infiltrated alumina sliding against reference pure alumina in water. J. Am. Cer. Soc. 88(2) (2005), 346–352. 19â•… S. Novak, M. Kalin, and T. Kosmacˇ. Chemical aspects of wear of alumina ceramics. Wear 250(1/12) (2001), 318–321. 20â•… J. Takadoum. Tribological behaviour of alumina sliding on several kinds of materials. Wear 170 (1993), 285–290. 21â•… M. G. Gee and N. M. Jennett. High resolution characterisation of tribochemical films on alumina. Wear 193 (1996), 133–145. 22â•… R. H. J. Hannink, M. J. Murray, and H. G. Scott. Friction and wear of partially stabilized zirconia: Basic science and practical applications. Wear 100 (1984), 355–366. 23â•… B. Basu, R. G. Vitchev, J. Vleugels, J. P. Celis, and O. Van Der Biest. Influence of humidity on the fretting wear of self-mated tetragonal zirconia ceramics. Acta Mater. 48 (2000), 2461–2471. 24â•… M. Kalin. Influence of flash temperatures on the tribological behaviour in low-speed sliding: A review. Mater. Sci. Eng. A 374(1) (2004), 390–397. 25â•… M. Kalin and S. Jahanmir. Influence of roughness on wear transition in glass-infiltrated alumina. Wear 255(1/6) (2003), 669–676. 26â•… Y. Y. Zhu, G. H. Kelsall, and H. A. Spikes. The influence of electrochemical potentials on the friction and wear of iron and iron oxides in aqueous systems. Tribol. Trans. 37(4) (1994), 811–819.
SECTION
VIII
CLOSURE
CHAPTER
31
PERSPECTIVE FOR DESIGNING MATERIALS FOR TRIBOLOGICAL APPLICATIONS In this concluding chapter, the need for appropriate designing of ceramic materials for tribological applications is discussed, followed by a discussion of the persistent need for the development of ceramics and ceramic composites with better fracture toughness and strength properties. Also, designing new materials with the ability to generate self-lubricating film in situ along with good resistance against material removal is stressed. The need for selection of appropriate testing conditions and environments is mentioned. An improved understanding of wear micromechanisms requires the use of some of the new characterization tools and this is also highlighted, along with the need to characterize the subsurface damage behavior. Finally, the need for modeling aspects coupled with control experiments to validate the models used to predict material properties and/or dependence of wear resistance on operating parameters is highlighted. Each material has some specific properties defining its potential use in practice. Not all materials can be used in every possible application, no matter how new, advanced, and extraordinary they may be in some respects. Therefore, materials need to be tailored for a particular use. Accordingly, before the best use can be determined, the materials need to be characterized by their relevant properties (mechanical, chemical, physical, etc.) and used in conditions that give them comparable advantages over other materials. Tribological properties are one of those characteristics that need to be studied and verified under the conditions that are already presupposed for a given application. Thereafter, the potential material is selected. So they are one step further toward the actual conditions and the performance under these conditions. In this book, a significant effort has been invested to discuss the wear micromechanisms of ceramics and composites. Ceramics have outstanding properties such as high elastic modulus, compressive strength, and hardness. Many ceramics have low density and they have the ability to retain hardness and strength at high Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
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temperature. In view of these properties, these materials can be used in various tribological applications; however, care must always be taken for some of the specific properties. In the past, ceramics have found some niche applications where they are almost irreplaceable. For example, silicon nitride had been successfully used in high-speed precision ball bearings due to its low density compared with steel and to its noncatastrophic failure mode. They are widely used in spindle hybrid or allceramic bearings. Silicon carbide has the strong advantage of being very corrosion resistant, as well as highly wear resistant, especially for a broad range of sliding applications. Therefore, one key application for silicon carbide is as mechanical axial face seals used for sealing technological media with aggressive properties for steels and rubbers. The most widely used ceramic is alumina, known for its high hardness and wear resistance. It is used both under dry and especially under water-lubricated conditions. Alumina is largely used in water-tap mechanisms. Biomedical applications must also be mentioned, with great potential for alumina that is used in joint prostheses. Apart from these positive examples, major disadvantages of ceramics are related to their poor fracture toughness, poor mechanical and thermal shock resistance, and cost of machining for precision engineering applications. These properties characterize all ceramics, some more than others; however, we cannot neglect them. The use of ceramics would be tremendously higher if the negative properties could be overcome! Many times, engineers look on ceramic materials as a problem-solving material, because of some great advantages. Namely, in many applications that require high loads and high precision, but high wear, ceramics may be the best suitable material. However, at the same time as they are introducing ceramics, the engineers are looking to replace it with a new material, because of the fear of breakdowns and fracture. It is difficult to predict fracture, and defects causing stress concentrations and premature failure are difficult to trace in advance. These materials need to improve primarily in these respects, but their manufacturing methods and mechanics should also improve. Finally, but probably most importantly, the component design must be suitable for ceramics, with all their positive and negative properties. If this is not considered, it is impossible to expect excellent results from any material, but especially ceramics. So engineers should understand these properties better and consider them in designs for ceramic material! By reducing the stress concentrations, tensile stresses, thermal shocks, thermal expansion coefficient mismatch, and so on, ceramics can become the perfect design material. We hope this book will bring forward some of this information to a broader audience. Nevertheless, it should be stressed that ceramics have improved significantly over the last decade as well. One who considers using ceramic material should know that there is no single type of a certain class of ceramics, but several of them. The major properties remain the same, but some can be significantly improved. For example, the production of silicon carbide is slowly but steadily increasing, coming very close to alumina, thanks to many specific improvements in certain properties— but not all in the same material! There are several microstructures, several grain
â•… 507
CHAPTER 31â•… Perspective for Designing Materials for Tribological Applications
boundary properties, and several sintering techniques that give new properties to these materials. They could improve toughness, or improve corrosion resistance, or reduce cost, and so on. Therefore, care should be taken in selecting the correct type of material. Last but far from least, there have been significant improvements and complementary properties in ceramic composites. The ever-increasing demand for developing ceramics or ceramic-based composites with better toughness and strength properties requires the careful selection of material combinations (second phase, dopant) as well as processing tools and parameters. The advent of fast sintering techniques, such as spark plasma sintering and hot isostatic pressing (HIP), offers such an advantage. The defect size in ceramics or brittle materials typically scales with grain size and, therefore, the availability of ultrafine-grained and nanostructured ceramics and composites is clearly a step ahead in realizing better wear-resistance applications. On the modeling aspect, some recent models for tribomechanical and tribochemical wear have been presented. The validation of such models requires control experiments with materials having properties measured over a wider range of operating parameters (load, sliding velocity). To verify the functional dependence of wear volume on material parameters, control experiments need to be performed on different materials with identical operating conditions. In the case of the tribochemical wear model, an absolute comparison of the experimental values for wear volume could not be made with the model-predicted values in the work presented here, due to unavailability of the experimental values of some material parameters. Therefore, work is required for the absolute validation of the tribochemical wear model. A transition from mild wear regime to tribochemical wear regime still cannot be identified using any existing model. To this end, the modeling efforts are required to define the conditions for such a transition. Another point of concern is the estimation of contact temperature. In the absence of any widely acceptable model, researchers use different models, and often the estimated temperatures vary over a considerable range for the same tribocouple. Since temperature is an important parameter, particularly for tribochemical wear models, a model’s predictions would therefore depend on the use of an appropriate thermal model. Among the various materials for biomedical applications, this book covers a range of materials, starting from hydroxyapatite-based composites to stabilized zirconia and then to machinable glass-ceramics. It has been recognized that the machinable mica-based glass-ceramics are worthy of huge technological interest because of their unique property of being machinable with the same machining tools used for metals. However, tribological testing on biomaterials needs to be carried out with a carefully chosen environment. For hard tissue replacement applications, the test medium should contain protein, such as bovine serum albumin, along with carefully controlled pH and temperature. Similarly for dental restoration applications, simulated tribological studies under the operational conditions of the oral cavity (load, frequency, duration, presence of saliva, etc.) need to be performed, with natural teeth as mating material. Also, results should be compared with baseline tests on natural bone and teeth. In various sections of this book, research results obtained using various counterbodies are presented. Since tribology is a system-dependent property, any claim
508â•…
CHAPTER 31â•… Perspective for Designing Materials for Tribological Applications
of industrial applicability of a new material requires testing under closely simulated conditions as well as with different counterbody materials over a range of speeds and/or test durations. Such research results will enhance the understanding of novel materials in wear applications, involving even situations with high contact stress. Often the potential role of humidity variation during real applications is not considered while designing the tribology tests. However, a number of experiments revealed that the influence of humidity is also material specific and, therefore, this factor needs to be critically considered. Considering the space-related applications for ceramic materials as well as to understand their wear properties in cryogenic environments, a number of chapters discuss the tribological properties of alumina, zirconia, and silicon carbide ceramics in liquid nitrogen. It has been demonstrated how the coefficient of friction and the wear rate can be scientifically explained on the basis of two effects: (1) cryogenic environment (cooling effect) and (2) material property effect, that is, in terms of elastic property, fracture behavior, thermal properties, and so on. One interesting observation is the limited tribochemical wear of SiC. In the future, similar experiments can be conducted to assess the tribological properties of silicon nitride ceramics under high-speed sliding conditions in liquid nitrogen. On the other hand, ceramics are globally recognized as materials for hightemperature applications. However, results of high-temperature tribological studies are rather limited in the published literature. The unavailability of test rigs that can be operated at temperatures above 1000°C often is a bottleneck for such studies. Room-temperature tribology data are never relevant for high-temperature applications, because extended testing at high temperature and reaction with the surrounding atmosphere can cause transition in wear mechanisms and often accelerate material removal. An improved understanding of the wear micromechanisms of new materials requires extensive and systematic understanding of the subsurface damage as well as analysis of the tribochemical layer or debris using some advanced characterization techniques, such as Raman spectroscopy, Fourier transform infrared spectroscopy (FT-IR), and so on. New techniques, such as focused ion beams (FIBs), can be used to investigate the fine-scale deformation or fracture of materials at various spatial distances across the worn surface Finally, careful component design requires the conceptual interaction and confluence of knowledge of mechanical engineering and materials science. To this end, the development of hybrid composites appears to be a better solution. For example, the compositional design of such composites can consider the addition of a second phase (e.g., MoS2), which can provide good lubrication properties; at the same time, the microstructure should be designed in such a way that it will have fewer cracks or defects with preferably finer grains. However, some of the tough ceramics such as silicon nitride-based materials often have elongated grain morphology. The component shape and tolerance in terms of dimension are equally important. All these factors will finally demand the selection of an appropriate processing and manufacturing route as well as process parameters. This is undoubtedly a challenging task for present and future generations of tribology researchers.
INDEX
Abrasion, 75–77, 95, 97, 163, 182, 193, 208, 248, 273, 285, 299, 322, 362, 375, 403, 418, 428, 433 micro, 320 parameter, 203, 207, 410, 417 two-body, 194 three-body, 194 Abrasive grooves, 335, 444 Abrasive scratches, 349, 479 Acetabular cup, 222, 234. See also Implant(s) Acid(s) lactic, 221 stearic, 143 Acid environment, 488, 494 Additive, 80, 101, 106 Adhesion, 73–75, 237, 273, 283, 378, 481 coating-substrate, 223 force, 57. See also Force(s) friction, 54. See also Friction interfacial, 16 interfacial energy, 424 theory, 348 work, 348 Adsorption chemical, 106 physical, 106 re-sorption, 228 Agglomerates, 496 Agglomeration, 418, 499 Al2O3, see Alumina Albumin, 297 Allergy, 288. See also Health
Alloy, 7 Co-Cr-Mo, 223, 234 metallic, 13, 454 Mg, 353 superalloy, 9 Ti-6Al-4V, 234, 355 titanium, 223, 234 Alloying, 118 Alumina, 134, 171, 201, 213, 226, 233, 235, 237, 242, 251, 254, 255, 260, 262, 273, 276–278, 283, 339, 408, 423, 451, 452, 459, 469, 481, 483, 487, 491, 498, 506 in cryogenic environment, case study, 439 glass-infiltrated, 213, 277, 284 glass-infiltrated, case study, 276 hydrated, 240 zirconia-toughened (ZTA), 13, 226 Aluminium hydroxides, 440, 444 Aluminosilicate, 129 Amelodentinal junction (ADJ), 258 Amontons’ laws, 49, 57 Amorphous phase, 494 Anisotropy, 8, 260 Annealing, 119 Antimicrobical properties, 272 Apatite, 229 Applications, 5, 7, 10, 14, 505, 506 high-temperature, 508 orthopedics, 248 tribological, 39, 155, 156, 505 Aqueous environment, 487
Tribology of Ceramics and Composites: A Materials Science Perspective, First Edition. Bikramjit Basu, Mitjan Kalin. © 2011 The American Ceramic Society. Published 2011 by John Wiley & Sons, Inc.
509
510
Index
Aqueous solution, 488, 496 acidic, 499 alkaline, 499 Archard equation, 413 Archard model, 458 Archard theory of wear, 73 Artificial hearts, 221 Artificial saliva. See Saliva Aseptic loosening, 218, 288. See also Health Aspect ratio, 8, 288 Asperity, 41, 52, 105 impacts, 51 Assay alkaline phosphatase (ALP), 218 biochemical, 218 comet (SCGE), 218 micronucleus, 218 MTT, 218 B4C, 134 Basal planes, 58 Bearing area, 192 Bending four-point, 28, 29 three-point, 28 Binder, 396 Bioceramics, 213 Biocompatibility, 213, 216–217, 272 Biocomposite(s), 213, 287 calcium phosphate-mullite, 288 HDPE-HAp-alumina, 235 hybrid, 213 polymer-ceramic, 213, 248 polymer-ceramic, case study, 233 Biointegration, 224 Biological cells, 218 Biomaterial(s), 215 calcium phosphate–based, 226 Blood, 221 Blunting, 277 Bond covalent, 8, 56, 73, 188, 194 debonding, 276 hydrogen, 73, 237 ionic, 8, 56, 73, 194 metallic, 56, 73, 194 Van der Waals, 73, 188, 237. See also Force(s) Bone, 214, 224 implants, 221
Boundary interface, 395. See also Layer Bovine serum albumin, 507 Bridging zone mechanisms, 35 Brittle materials, 23, 31, 77 Brittleness, 18, 33 Calcium phosphate, 224, 252, 287 biphasic (BCP), 227 Carbide(s) cemented, 320 mixed (WC/TiC), 408 secondary, 353, 356, 377, 386, 403 ternary, 185 Carbon-fiber-reinforced plastics (CFRP), 235 Carcinogenic response, 216 Case study (W, Ti)C-Co cermets, 407 alumina in cryogenic environment, 439 dental restorative materials, 251 glass-infiltrated alumina, 276 nanocrystalline yttria, 325 nanostructured tungsten carbide-zirconia nanocomposite, 338 polymer-ceramic biocomposites, 233 sialon ceramics, 167 silicon carbide in cryogenic environment, 469 tetragonal zirconia ceramics in liquid nitrogen, 454 Ti3SiC2 MAX phase, 185 titanium carbonitride-nickel cermets, 377 titanium diboride (TiB2), 197 tooth restorative materials, 251 transformation-toughened zirconia (TZP), 142 zirconia polycrystalline ceramics (Y-TZP), 325 Catheter, 221 Cell culture, 217 functionality, 218 Cement, 223 Ceramics, 7, 13, 505 biocompatible, 226 bioinert, 226 classification, 8 covalent, 34 feldspathic machinable, 253 fiber-reinforced, 12
511
Index
fillers, 237, 248 ionic, 34 low-fusing hydrothermal, 253 monolithic, 12, 134 nanostructured, 307, 320, 321 non-oxide, 9, 126, 132, 183, 321 oxide, 9, 321, 487 structural, 9, 11, 113, 439 traditional, 9 ultrafine-grained and nanostructured, 507 Cermet(s), 13, 197, 353, 356 (W,Ti)C-Co cermets, case study, 407 (W,Ti)C-Co mixed carbide, 358 composition, 403 mixed carbide, 407 TiCN, 356, 358, 377, 381, 400, 403 TiCN-WC-Ni, 346, 377, 378, 393, 398 TiCN-Al2O3-Ni-Mo, 379, 398 TiCN-Ni, 358, 377, 386, 387, 396, 398, 403 TiCN-Ni, case study, 377 TiCN-Ni-Mo, 378, 379, 396 TiCN-Ni-WC, 353, 398 WC–based, 134, 314 WC-Co, 353 WC-TiC-Co, 407, 418 Chemical reaction, 107 Chemical resistance, 134 Chemical stability, 13, 357 Chemical vapor deposition (CVD), 224 Chemistry, 4 Chewing, 251 Chipping, 462, 463 Cleavage steps, 451 Clinoenstatite (MgSiO3), 366 (MgSiO3·3H2O), 375 Coating(s), 12, 223, 276, 454 diamond-like carbon (DLC), 224, 454 self-lubricating, 454 Coefficient of friction, 50. See also Friction Coefficient of friction (values, data) acetal, 56 Ag, 55 Al, 55 alumina, 55, 136, 238, 272, 280, 442, 481 Au, 55 B4C, 136 BN, 55
CaO-MgO-Al2O3-SiO2, 272 CaO-MgO-Al2O3-P2O5-F, 272 cast iron, 55 Co, 55 Co-alloy, 55 Cr2O3, 55 Cu, 55 diamond, 55, 195 Dicor®, 272 Fe, 55 fullerene (C60), 56 glass-infiltrated alumina, 280 graphite, 56 HAp–based composites, 300 high density polyethylene (HDPE) (–based), 56, 238, 300 In, 55 K2O-B2O3-Al2O3-SiO2-MgO-F, 272 leaded brass (Cu, Zn, Pb), 55 Mo, 55 MoS2, 56 Ni, 55 Ni-alloy, 55 polyamide (PA) (nylon), 56 polytetraflouroethylene (PTFE) (teflon), 56 Pb, 55 Pt, 55 rubber (natural, synthetic), 56 rubber (silicone), 56 Si3N4, 55, 136 SiC, 55, 136, 470, 481 Sn, 55 steel (mild), 55 steel, 195, 272, 300, 386, 429 TaB2, 136 teeth (human), 272 Ti, 55 TiC, 55 TiCN cermets, 386, 387 (W,Ti)C-Co cermets, 136 TiN, 55 Ti3SiC2, 195 ultra high molecur wight polyethilene (UHMWPE), 300 W, 55 WC, 55, 136 zirconia (-based), 238, 300, 455, 481 Coefficient of thermal expansion, 276 Cold work, 437
512
Index
Composite(s), 8, 227, 234, 505 Al-alumina, 375 Al-SiC, 354, 364, 375 ceramic, 12 Cu-alumina, 375 HAp, 299, 507 HAp-Ti-6Al-4V, 224 HAp-alumina, 227 HAp-bioglass, 227 HAp-mullite, 302 HAp-TiO2, 227 HAp-zirconia, 227 metal matrix, 353 Mg-9Al-Zn matrix, 354 Mg-SiC, 353, 354, 362, 363, 375 nanostructured, 307, 320 particulate-reinforced, 362 polymer-ceramic, 234 TiB2-MoSi2, 198, 200 TiB2-TiSi2, 198 zirconia (Y-TZP)-ZrB2, 466 ultrafine-grained and nanostructured, 507 WC-Co, 339 WC-zirconia, 326 Composition, 169 design, 197, 378 sialon ceramics, 169 Conduction, 44 Conservation integral, 116 Contact, 39 bioadhesive, 216 elastic, 47 non-conformal, 44 plastic, 47 pressure, 44, 45, 296 single-asperity, 43 stress, 155 temperature, 60–61, 80, 151. See also Model mechanics, Hertzian, 39. See also Theory, Hertzian Contaminant, layers, films, 52, 56, 74 Convection, 44 CoO, 148 Core-rim grain, 358 Correlation length, 481 Corrosion resistance, 132, 507 Coulomb’s law, 50, 57 Crack(s), 96–98, 247, 284, 462, 463 binding, 129
bridging, 36, 113, 182 cone, 23 deflection, 35, 182, 317, 475 growth, 22, 35, 284 initiation, 78 lateral, 24, 25, 33, 98, 333–334, 417 length, 20, 33 long, method, 31 median, 24, 25, 31, 33, 97, 98 opening modes, 22 primary, 125 propagation, 20, 36, 78 short, method, 31 tortuous, 444 Crack tip, 18, 116 shielding, 125, 454 stress intensity, 125 Cracking, 45, 159, 164, 349, 404, 475, 483 brittle materials, 23 intergranular, 462, 465 micro, 463, 465, 466, 480. See also Microcracking transgranular, 465, 475 Crown, 252 Cryogenic environment, 13, 421, 449, 454, 466, 482, 508 Cryogenic technology, 423 Cryogenic temperatures, 423 Cryotribometer, 425, 437 Crystallographic orientation, 450, 451 Cu, 427, 433 Cutting, 71, 75, 94, 96 tools, 357 Cytotoxicity, 218, 242, 248 Deformation, 47, 437, 479 elastic, 103 high-temperature, 356 homogeneous, 34 plastic, 16, 75, 94, 96, 159, 193, 248, 320, 326, 378, 404, 424, 427, 428, 456 viscoelastic, 16, 57, 424 Delamination, 45, 78–79. See also Wear Densification, 309, 338, 358, 377 Density, 7, 10, 11, 126, 186, 273, 357, 410, 425, 469, 505 Dental implants, 221, 223, 251 Dental restoration, 277 Dental restorative materials, 272, 273 case study, 251
513
Index
Dentin, 251, 258, 273 Dentinal tubule, 252, 258 Dentin-enamel junction (DEJ), 251, 258 Dentistry, 214 Design, 5 Designing materials, 505, 506 Dislocation, 78 density, 450 glide, 34, 35 Dispersion, 36 Dissociation, 227 Dissolution, 198, 254, 298, 377, 492, 499 Dopant cation, 121 content, 154 distribution, 154 Drilling, 71 Ductile metals, 427 Ductility, 3 Durability, 7 Economic impact, 70 Elastic modulus, 7, 10, 26, 30, 37, 185, 247, 249, 296, 314, 466, 481, 505. See also Mechanical properties; Mechanical properties (data, values) Elastomers, 57 Electrical conductivity, 13 Electrochemical deposition, 224 Electrochemical effects, 500 Electrochemical forces, 500. See also Force(s) Electrochemical potential, 498 Electrochemical properties, 491 Electrostatic repulsion, 499. See also Force(s). Enamel, 251, 258, 273 Energy consumption, 7 dissipation, 65, 125, 202, 207, 381, 383, 451 kinetic, 95 Engineering, 4 ceramics, 9 surfaces, 52 systems, 7 Enstatite (MgSiO3), 373 Environment, 3 Enzyme, 218, 221
Erosion, 92–98. See also Wear brittle materials, 96–98 ductile materials, 94 Ethylene vinyl acetate, 234 Exfoliation, 58 Extracellular matrix, 217 Fabrication, 5, 7 Fatigue, 78–79, 273, 276, 283, 465, 466. See also Wear cracking, 299 long-term, 78 low-cycle, 78 static, 284, 285 Fe(OH)3, 282 Fe2O3, 191, 402 Fe2SiO4, 282 Fe3O4, 402 Fe9TiO15, 390, 392 FEM calculations, 154 Femoral ball, 234. See also Implant(s) Fiber, 8, 35–36 reinforcement, 35, 36 Fibroblast cells, 242 Film(s). See also Layer boundary, 54, 80, 81, 101, 106. See also Lubrication, boundary regime chemical, 56 thickness, 104. See also Lubrication, boundary regime Fish-scale pattern, 462, 463, 465, 466 Flaw size, 22 Flocculation, 492 Fluid flow, 92 Force(s) adhesive, 57 attractive, 499 cohesive, 496 electrochemical, 500 friction, 49 interatomic, 19, 56 repulsive, 499 tangential, 49 Van der Waals, 57, 58 Forsterite (Mg2SiO4), 373 Fracture, 208, 299, 374, 396, 450, 499, 506 brittle, 18, 273, 318, 322, 378, 439, 450, 452, 465, 482, 483 detachment, 500
514
Index
Fracture (cont’d) intergranular, 132, 168, 325, 326, 330, 334, 335, 347, 444, 451, 452 mechanics theory, 18, 24. See also Theory mixed-mode, 444 resistance, 131 transgranular, 444, 451, 452 Fracture strength, 29, 113, 276, 314, 339. See also Strength Fracture toughness, 11, 23, 31–33, 36–37, 77, 113, 126, 132, 163, 313, 314, 356, 377, 410, 418, 424, 455, 505. See also Mechanical properties; Mechanical properties (data, values); Toughness Fretting, 71, 81–92. See also Wear corrosion, 82 fatigue, 82, 375 friction logs, 92 maps, 89–90 mechanics, 84 modes, 82–84 regimes, 82, 86–89 wear, 81–82 Friction, 3, 8, 49, 163. See also Coefficient of friction; Coefficient of friction (values, data) DaVinci, 50 energy, 60, 65, 330 force, 49. See also Force(s) heating, 11, 49, 60, 81, 455 kinetic, 50, 55 laws, 49, 57 loss, 44 mechanisms, 51 of engineering materials, 54 static, 50 steady-state, 52, 56 Fullerenes, 58 Galling, 433 Genotoxicity, 213, 218 Gibbs Free energy, 373, 402 Glass, 7 45S5, 225 bioactive, 225 Bioverit®I, 225 ceramics, 221, 224, 251, 256, 262, 266, 271, 507 infiltration, 276
Nioverit®II, 225 phase, 129, 182 pyrex, 277 residual, 128 transition temperature, 224 Glass-infiltrated alumina. See Alumina. Gold, 253 Grain aspect ratio, 129 bonding, 132 boundary, 182, 334, 335, 444, 450, 469, 475, 482 bridging, 131 elongated, 131, 132, 168, 182, 187 equiaxed, 126, 168 growth, 338 pull-out, 75, 131, 168, 318, 325, 326, 330, 333, 334, 335, 444, 480 refinement, 317 shape, 126, 129, 167 size, 119, 145, 154, 167, 507 Graphite, 58, 188 Griffith’s theory, 20 Grinding, 71 Hardmetal, 356, 407 Hardening, 4 Hardness, 3, 10, 12, 26, 37, 73, 76, 126, 167, 185, 207, 247, 248, 273, 299, 308, 313, 314, 314, 317, 318, 334, 335, 349, 357, 403, 410, 418, 425, 437, 469, 481, 505, 505. See also Mechanical properties; Mechanical properties (data, values) HCl, 488, 493 Health allergy, 288 aseptic loosening, 218, 288 biocompatibility, 216 cytotoxicity, 218 genotoxicity, 213, 218 host response, 216 inflammation, 218 osteogenesis, 218 septic loosening, 218 Heat, 44, 60. See also Friction, heating dissipation, 439, 452, 482 generation, 62 steady-state flow, 63 transient flow, 63
515
Index
Hertzian theory, 44, 84, 296 Hexadecane, 143, 455, 456 HfC, 353, 356, 377, 398, 400, 403 HfO2, 402 High-tech materials, 6 High-temperature applications, 508. See also Applications Hip joint prosthesis, 222 replacement, 234 Histopathology, 220, 222. See also Health Hook’s law, 20 Host response, 216. See also Health Humidity, 58, 138, 142, 144, 156, 159, 160, 164, 167, 335, 440, 455, 508 Hydroxyapatite (HAp or HA), 220, 227, 235, 242, 252, 258, 287, 298, 300 bioceramics, 13 mullite-reinforced, 288, 289, 299 Impact angle, 92, 94, 95 Implant(s), 220, 221, 233, 288 loosening, 233. See also Health Indentation cracking, 31–33, 131 fracture theory, 97 strength, 121 toughness, 207 Vickers hardness, 121 Indenters, 31 Inertness, chemical, 318 Inflammation, 218 Interaction(s), 4 asperity-asperity, 47 chemical, 168 phonon-phonon, 449 physicochemical, 60 solid-liquid, 4, 109 Interatomic distance, 20 Interatomic force, 19, 56. See also Force(s) Interfacial properties, 15 Interfacial reactions, 227 Interlocking, 52, 182, 348 Iron oxide, 191, 282, 347, 387, 402, 498 Iron oxide debris, 263. See also Layer Iron-titanium oxide, 404 Iso electric point (IEP), 492, 498, 499 Junction growth, 56. See also Adhesion; Friction, mechanisms
Lambda parameter, 104 Layer, 188. See also Tribochemical layer apatite, 297 boundary, 51, 298 gel-like reaction, 254 hydrated, 293, 298, 490, 491, 496, 497 iron oxide, 347, 387. See also Iron oxide low-shear, 107 mechanically mixed, 433 occlusal, 258 protective, 163, 164, 444 sacrificial, 107 tribological, 491, 499 wear debris, 490–493, 494–497, 500 Liquid hydrogen, 440 Liquid nitrogen, 423, 430, 439, 451, 452, 456, 459, 464, 466, 469, 477, 480, 508 Liquid oxygen, 440 Lubricant, 56, 70, 101, 440 solid, 57, 454 Lubrication, 3, 70, 101, 134, 137, 449, 478, 508 boundary regime (BL), 106, 487 design, 109 elastohydrodynamic regime (EHD), 103 hydrodynamic regime (HD), 102, 487 mixed regime (ML), 105 regimes, 101 self-lubricating, 57 Machinability, 133, 185, 272 Machining, 9, 277, 506 Machining-induced damage, 276 Magnesia (Mg(OH)2), hydrated, 372, 375 Magnesium silicate, hydrous, 375 Maintenance, 70 Mastication, 251 Material(s) bioactive, 217 bioinert, 216 composition, 37 metallic, 10 properties, 3 transfer, 274, 403 Materials engineering, 4 Materials science, 4 Matrix, 8, 36 Matrix-fiber interface, 36 Mechanical engineering, 4
516
Index
Mechanical properties, 10, 18, 37, 145, 200, 215, 256, 260, 320, 363, 381, 408, 409, 413 Mechanical properties (data, values) Al-alloy, 10 alumina, 10, 26, 27, 135, 172, 278 B4C, 26, 135 cast iron, 10 Co-alloy, 26, 27 CrB2, 26, 27 glass ceramics, 256 glass-infiltrated alumina, 278 HAp-mullite composites, 296 high density polyethylene (HDPE), 10, 236 hard tissue (human), 215 Mg (-based materials) composites, 363 Mo, 26, 27 MoSi2, 26, 135 Nb, 26, 27 NbC, 26, 27 Ni-alloy, 26, 27 polyamide (PA) (nylon), 10 polyimide (PI), 10 polytetraflouroethylene (PTFE), 10 Si3N4, 10, 26, 27, 130, 135 sialon, 130, 169, 172 SiC, 10, 26, 27, 135 SiO2, 26, 27 S-sialon, 130, 172 steel, 10, 172 Ta, 26, 27 TaC, 26, 27 TaB2, 26, 27 TaN, 26, 27 TiB2 (-based materials), 26, 27, 135, 200 TiC, 26, 27, 135 TiCN cermets, 381 (W,Ti)C-Co cermets, 408 TiN, 26, 27 TiO2, 26, 27 TiSi2, 26, 27 W, 26, 27 WC, 26, 27, 135 WSi2, 26, 27 zirconia (-based), 10, 26, 27, 120, 121, 145 ZrB2, 26, 27, 135 ZrC, 26, 27
ZrN, 26, 27 ZrSi2, 26, 27 Mechanical properties measurements, 24 Melting point, 7, 8, 10 Metallurgy, 4 Metals, 7 Mg, 311, 375 Mg(OH)2. See Magnesia MgO, 362, 375 Mica, 251, 254, 271 Microcracking, 35, 121, 125, 163, 182, 299, 320, 322, 326 Microcracks, 273, 297 Microfracture, 142, 450, 452 Micromechanical modeling, 116 Microstructure, 6, 73, 113, 118, 126, 132, 208, 313 cubic, 155 design, 308 properties, 409 refinement, 320 tetragonal, 155 Milling, high-energy ball (HEBM), 311 Mo2C, 356 Model Archard’s temperature, 66, 178, 271, 392. See also Contact, temperature contact temperature, 62 fracture-mixed, 444 Kong-Ashby temperature, 67, 433 Robert’s wear, 418 Modeling, 507 Moist atmosphere, 366 Moisture, 375 MoO3, 379 Morphological characterization, 146 MoS2, 58, 188, 197 Mullite, 201, 277, 296, 299, 378 needles, 288 reinforcement, 287 Nanoceramics, 305, 325 alumina, 338 composites, 307, 325 monolithic, 311 yttria-stabilized zirconia, 330, 333 Nanocomposite(s) Al2O3-BaTiO3, 314 Al2O3-Nd2Ti2O7, 314 Al2O3-SiC, 313, 319, 326, 339, 347
517
Index
Al2O3-TiC, 314 Al2O3-TiO2, 327, 339 alumina-based, 313 ceramics, 338, 339 design, 313 Si3N4-BN, 339 Si3N4-SiC, 326 WC, 314, 393 WC-MgO, 317 WC-ZrO2, 338, 339, 340, 349 WC-ZrO2, case study, 338 yttria-stabilized polycrystalline zirconia, 339 zirconia-based, 317 ZrO2-WC, 317 ZrO2-ZrB2, 317 ZrO2-Al2O3, 317 Nanocrystalline ceramics processing, 309 Nanocrystalline materials, 338 Nanocrystalline reinforcement, 308, 309 Nanocrystalline Y-TZP, 347 Nanocrystalline α-alumina, 311 Nanoreinforcement, SiC, 330 Nanotribology, 106 Nanotubes, 58 multiwall carbon (MWCNT), 235 Nanowires, 58 Navier-Stokes equation, 102 Nb2O5, 402, 403 NbC, 353, 356, 377, 398, 400, 403 NbO, 403 NbO2, 403 NH4OH, 488, 493 Ni, 400 Ni-Cr, 355 NiO, 402 Nitrides, ternary layered hexagonal, 185 Nitriding, 4 Nitrogen (dry), 144 Nitrogen atmosphere, 326 Noise, 49 Octacalcium phosphate (OCP), 227 OH radicals, 240 Oil, 101 lubrication, 137 mineral, 143 paraffin, 138, 440, 455, 456 viscosity, 103
Orthopedics, 214 applications, 248. See also Applications joints, 295. See also Implant(s) Osseointegration, 242, 273. See also Health tests, 249 Osteoblasts, 218 Osteogenesis, 218. See also Health Oxidation, 80–81, 185, 347, 397, 398, 401–403, 404 high-temperature, 401 layer formation, 375 reaction, 403 resistance, 358 wear, 80–81, 374, 375, 378, 427, 428, 433, 437. See also Wear Oxide, 52, 56, 477 debris, 269, 293. See also Wear debris hydrous, 375 Particle(s) amorphous, 298 erosion, 92–98 hardness, 94 nano, 58 shape, 94 size, 94 strength, 94 wear, 52, 74. See also Wear debris Particulates, 8, 35 Peclet number, 66, 458 Perfluoro-polyalkylether, 143 pH, 488, 492, 494, 498, 499, 507 Phase diagram S-sialon, 129 Ti-B, 133 Physical properties (data, values). See Mechanical properties (data, values) Physical science, 213 Physics, 4 Physiological conditions, 217 Physiological environments, 228 Physiological fluids, 221 Pitting, 47 Plastic flow, 74, 452 Plasticity, 185, 194, 435 Platelet crystals, 128 Plowing, 75, 76, 95, 182, 299, 433, 437. See also Abrasion Poisson’s ratio, 30 Polishing, 71
518
Index
Polymethyl methacrylate (PMMA), 235 Polyamide (nylon), 57 Polyether ether ketone (PEEK), 235 Polyethylene (PE), 57, 213, 255, 424 high-density (HDPE), 233, 242 ultra-high-molecular-weight (UHMWPE), 220, 234 Polylactide (PLA), 235 Polymer(s), 7, 10, 233, 424 Polyoxymethylene (POM, Acetal), 57 Polytetrafluoroethylene (PTFE), 16, 57, 424, 454 Porcelain, 253 Powder(s), 127, 309 premixed, 410 solid solution, 418 starting, 418 Precipitation, 377, 499 Pressure-viscosity coefficient, 103 Process zone mechanisms, 35 Prostheses joint, 506 Protein, 217, 221, 229, 287, 507 Pull-out, 167, 182, 347, 407. See also Grain, pull-out crystal, 251 Pyrophosphates, 218 Pyroxene ((Mg,Fe)SiO3), 366 Radiation, 44 Real contact area, 42, 60, 62 Reinforcement, 8, 37, 113, 227, 249, 317 nano-SiC, 326 particulate, 35 self-reinforcement, 131 Repulsive force, 499. See also Forces Resistance curve (R-curve), 31, 116, 132 Reynolds equation, 102 Rolling, 71 Roughness, 39, 47, 54, 104, 192, 492 Rule of mixture, 393 Running-in, 47, 51, 52, 138, 273, 329, 375, 443, 471, 482 Saliva, 251, 257, 262, 273 Sealing, 506 Secondary phase, 35, 198 Septic loosening, 288. See also Health Shear modulus, 451 strain. See Strain
strength. See Strength stress. See Stress Si3N4, 126, 134, 167, 168, 171, 198, 278, 311, 326, 335, 373, 379, 408, 439, 470, 487, 506 Sialon, 13, 126, 169, 171, 179 compositionally tailored, 172 case study, 167 S-sialon. See S-sialon SiC, 13, 134, 201, 278, 313, 335, 354, 362, 373, 423, 480, 481, 483, 487, 506 in cryogenic environment, case study, 469 Silica, 191, 366, 373, 476 amorphous, 282, 483 hydrated, 373, 375 layer, 174, 475 nanoceramics, 338 Silicides, 206 Silicon carbide. See SiC Silicon nitride. See Si3N4 Simulated body fluid (SBF), 223, 233, 242, 248, 287, 288, 289, 295, 296, 298 Sinterability, 133 Sintering, 37, 127, 197, 207, 309 additives, 197, 207 aids, 142 field-assisted technique (FAST), 309, 325 hot isostatic pressing (HIP), 308, 309, 507 hot pressing route, 144 liquid phase, 199 pressure-less, 225, 309, 316, 407 spark plasma (SPS), 308, 325, 334, 338, 410, 507 techniques, advanced, 313, 507 three-stage, 410, 418 Sliding, 49, 71 Slip planes, 33 systems, 34 Slippage, 188 Smearing, 374, 418, 432, 437 Solubility, 74, 198, 492, 497, 499 Spalling, 163, 248, 274, 349, 407, 418, 462, 463, 466, 480, 482 S-phase. See S-sialon Spherulitic-dendritic crystals, 254 Spinning, 71
519
Index
S-sialon, 127, 167 Ba-doped, 127, 182 Sr-doped, 127 Staphylococcus epidermidis, 226 Steel, 171, 186, 256, 263, 339, 347, 354, 355, 378, 427, 454 stainless, 234, 424 Stent, 222 Stick-slip, 51. See also Fretting, regimes Stiffness, 125, 335. See also Mechanical properties Strain. See also Mechanical properties energy, 20 hardening, 74 incompatibility, 334 shear, 125, 451 viscoelastic, 247 Strength, 3, 8, 12, 126, 276, 308, 318, 356, 505. See also Mechanical properties cohesive, 19 compressive, 10, 11, 27–28, 111, 450 flexural, 28–30 fracture, 29, 113, 276, 314, 339 shear, 107 tensile, 28, 450 yield, 437 Strengthening, 276 Stress compressive, 32, 35, 276, 277 concentrations, 506 dilatational-hydrostatic, 153 field, 18 Hertzian, 44. See also Theory plasticity-induced, 318 residual, 33, 126, 276, 284 residual factor, 33 shear, 47, 437 tensile, 334, 506 von Mises, 153, 437 yield, 440 Stress-corrosion cracking, 497 Stress-intensity factor, 22, 23, 116 Stress-strain behaviour, 28 Stribeck curve, 107–109 Stribeck parameter, 107 Structure chain, 57 core-rim, 377 crystal, 134 cubic, 469
grain, 199 hexagonal, 469 layered, 58 Subsurface damage, 52 region, 297 stress, 39 Surface, 3, 39 asperity slope, 42 coating, 276 cracks, 277. See also Crack(s); Cracking fatigue, 374, 378. See also Fatigue finish, 57, 73 hardening, 451, 452, 482 layer, 106. See also Layer roughness. See Roughness stress, 39. See also Stress topography, 41 Surface charge, 488, 499 Surface energy, 20, 109, 450 Surgery, 233 Synthesis, high-frequency induction-heated combustion, 410 System-dependent property, 507 Ta2O5, 402 TaC, 353, 356, 377, 398, 400, 403 Tangential force, 49. See also Force(s) Tearing, 22, 251 Technical diagnostics, 70 Techniques, advanced, characterization, 508 Teeth. See Tooth Temperature, 3, 298, 392, 393, 397 bulk, 61, 68 contact, 60, 151, 507. See also Model energy, 65 flash, 61, 66, 68, 430, 437, 458, 466, 477, 483 gradient, 60 maps, 63 rise, 62 Test in vitro, 217 in vivo, 217 Tetracalcium phosphate (TTCP), 227 Theory adhesion, 348 Archard’s wear, 73 delamination, 45
520
Index
Theory (cont’d) fracture mechanics, 24 Griffith’s, 20 Hertzian, 44, 84, 296 indentation-fracture, 97 Rabinowicz, 348. See also Adhesion Therapy, 216 Thermal capacity, 60 conductivity, 13, 60, 65, 143, 314, 358, 424, 437, 449, 452, 466, 469, 481 convection, 65 diffusivity, 60 energy, 62 expansion, 132, 314, 358 expansion coefficient, 506 properties, 10 radiation, 65 severity, 477 shock resistance, 185, 126, 356, 357, 378, 466, 506 Thermodynamic data, 283 Thermoplastics, 57 Thermosets, 57 Third body, 194. See also Layer; Wear debris Tri calcium phosphate (TCP), 227, 288 Ti, 186, 355, 427, 430 Ti/HAp/bioactive glass, 225 Ti3B4, 134 Ti3SiC2, 195 MAX phase, case study, 185 TiB, 134 TiB2, 13, 132, 197, 200, 311 hexagonal, 134 case study, 197 TiC, 198, 356, 411 TiN, 198, 311 TiO2, 191, 226, 311, 330 nanoreinforcement, 330 layers, 430 TiSi2, 197 Tissue, hard, 215 Tolerances, 73, 357, 508 Tooth human, 224, 251, 254, 257, 260, 262, 273 natural, wear mechanisms, 272 restorative materials, case study, 251
Topography, 40, 54. See also Roughness Toughening mechanisms, 33–37, 113, 131, 316 Toughness, 7, 12, 33–37, 73, 113, 121, 154, 167, 207, 334, 466. See also Fracture toughness Toxicity, 147, 216, 248, 273, 288 Transfer films, layers, 56, 73. See also Film(s); Layer Transformability, 124, 154 Transformation hydrothermal, 143, 497–499 martensitic, 114, 125 phase, 35, 162, 322, 455 strength, 142 stress-induced, 116, 125, 153 tetragonal-to-monoclinic, 114, 121, 163, 334, 454, 498 toughening, 12, 35, 113, 116, 118, 153, 154, 454 toughness, 142 zirconia phase, 149, 460 zone, 33 Transition crysralline-to-amorphous, 298 metals, 13, 185 regime(s), 89, 105 Translucency, 272 Tresca criterion, 437 Tribochemical layer, 57, 80, 81, 193, 254, 273, 274, 276, 283, 376, 404, 407, 418, 475, 477, 483, 490–492, 496, 498, 499, 508. See also Layer Tribochemical reaction(s), 3, 15, 49, 142, 178, 191, 282–284, 298, 362, 372–375, 378, 387–390, 397, 398, 401–403, 483, 492 Tribochemical wear, 80–81, 160, 164, 168, 182, 208, 285, 322, 338, 349, 371, 375, 378, 387–393, 469, 473, 482, 483, 507, 508 Tribochemistry, 3, 80 Tribolayer, 298, 395–396, 404, 432, 433, 490–497, 499. See also Layer; Tribochemical layer Tribological applications, 5–9, 11–16, 70, 71, 155, 156, 505. See also Applications contact, 9, 44. See also Contact
521
Index
Tribology, 3, 7, 14, 39, 49, 70, 101, 505 Tribooxidation, 138, 346. See also Oxidation Ultrasonic method, 30 Vacuum, ultra-high, 58 Van der Waals force, 57, 58 Velocity accommodation, 91 Vibration, 49 Vickers indentation, 24, 31, 33 Viscosity, 449 Water, 16, 137, 168, 440, 455, 456, 487, 488, 492, 506 distilled, 254, 298 lubrication, 487, 496, 499, 506. See also Lubrication Waviness, 40 WC, 134, 311, 346, 353, 356, 377, 393, 398, 400, 403, 407, 411 WC-CO, 147, 149, 151, 202, 204, 316, 321, 407, 408 Wear, 5, 70–100, 505 abrasive, 75. See also Abrasion adhesive, 73. See also Adhesion attrition, 471 classification, 72 cryogenic, 459, 466 debris, See Wear debris delamination, 45, 78–79, 159, 194, 254, 299, 325, 330, 334, 496 diffusive, 407 electrochemical, 498 erosive, 92. See also Erosion fatigue, 78–79. See also Fatigue fretting, 81–92. See also Fretting in vitro, 213, 217 in vivo, 213 loss, 163 lubricated, 455 map, 398, 430, 432 mechanism, 70, 72, 152, 157 micromechanisms, 296 mild, 348, 449 model, 72. See also Theory oxidative, 80–81. See also Oxidation resistance, 3, 37, 142, 144, 207, 413 severe, 348, 449
three-body, 75 tribochemical, 80–81. See also Tribochemical wear tribomechanical, 37, 159, 164, 182, 273, 330, 333, 335, 428 two-body, 75 Wear debris, 72, 194, 205, 229, 233, 248, 274, 282, 293, 348, 349, 387, 418, 449, 459, 466, 494 agglomerated, 270 amorphous, 299 detachment, 500 layer, 494, 498, 499. See also Layer needle-like, 366 platelet-like, 347 polycrystalline, 299 submicron, 347, 348 Wetting, 109 Whisker(s), 8, 35, 228 pullout, 36 reinforced ceramic, 12 reinforced composites, 35 WO3, 148, 311, 346, 379, 402, 403, 411 Work-hardening, 78, 96 Work-softening, 78 WS2, 58 Young’s modulus. See Elastic modulus Yttria, 118 content, 121 distribution, 122 dopant, 330 doped, 311 nanocrystalline, case study, 325 stabilisation, 122, 155 hydroxide, 163, 164 Zeta potential (ZP), 488, 499 Zirconia, 13, 113, 156, 233, 237, 245, 288, 339, 378, 423, 459, 460, 464, 466, 469, 481, 483, 487, 493, 498 coatings, nanostructured, 326 cubic (c-ZrO2), 114, 154, 460 fully-stabilised (FSZ), 114 hot isostatic pressing (HIP), 121 metastable orthorhombic (o-ZrO2), 460 monoclinic (m-ZrO2), 226, 328, 460, 497 nanostructured, tetragonal (t-ZrO2), 327, 334
522
Index
Zirconia (cont’d) nontransformable, tetragonal (t′-ZrO2), 115 orthorhombic (o-ZrO2), 466 partially-stabilised (PSZ), 114, 226 polycrystals (TZPs), 114 rhombohedral (r-ZrO2), 115 stabilisation, 117 stabilized, 335, 507 tetragonal (t-ZrO2), 114, 460 tetragonal (t-ZrO2) in liquid nitrogen, case study, 454
transformation-toughened tetragonal polycrystalline (TZP), case study, 142 yttria-stabilized tetragonal polycrystalline (Y-TZP), 142, 144, 145, 155, 163, 311, 325, 327, 330, 425, 455 yttria-stabilized tetragonal polycrystalline (Y-TZP), case study, 325 ZnO, 226 ZrB2, 311 ZrO2. See Zirconia Zr-hydroxide, 162, 163, 164, 497