Advanced Inorganic Fibers: Processes - Structures - Properties - Applications

ADVANCED INORGANIC FIBERS Processes - Structures - Properties - Applications

MATERIALS TECHNOLOGY SERIES Series editor: Renee G. Ford The Materials Technology series is dedicated to state-of-the-art areas of materials synthesis and processing as related to the applications of the technology. By thorough presentation of the science underlying the technology, it is anticipated that these books will be of practical value both for materials scientists and engineers in industry and for engineering students to acquaint them with developments at the forefront of materials technology that have potential commercial significance. Ceramic Injection Molding Beebhas C. Mutsuddy and Renee G. Ford Hardbound (0 412 538105) Cryochemical Technology of Advanced Materials Yu. D. Tretyakov, N.N. Oleynikov and O.A. Shlyakhtin Hardbound (0412 639807) Modelling of Materials Processing Gregory C. Stangle Hardbound (041271120 6) Porous Materials Kozo Ishizaki, Sridhar Komarneni, Makota Nanko Hardbound (0412711109) Functionally Graded Materials Yoshinari Miyamoto, Wolfgang A. Kaysser, Barry H. Rabin, Akira Kawasaki, Renee G. Ford Hardbound (0412 607603)

ADVANCED INORGANIC FIBERS Processes - Structures - Properties - Applications Contributors:

FREDERICK T. WALLENBERGER Manager, Advanced Technology PPG Fiber Glass Research Center, Pittsburgh, Pennsylvania

ROGER NASLAIN Professor, University of Bordeaux Director, High Temperature Structural Composites Laboratory Pessac, France

JOHN B. MACCHESNEY Fellow, Bell Laboratories Lucent Technologies, Murray Hill, New Jersey

HAROLD D. ACKLER Lawrence Livermore National Laboratory Livermore, California

Editor: FREDERICK T. WALLENBERGER

KLUWER ACADEMIC PUBLISHERS Boston / Dordrecht / London

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Library of Congress Cataloging-in-Publication Data Advanced inorganic fibers : processes--structures --properties--applicationsl contributors, Frederick T.Wallenberger. . .[et al.]; editor,Frederick T. Wallenberger. p. em. -- (Materials technology series) ISBN 0-412-60790-5 1. Inorganic fibers I. Wallenberger, Frederick T., 1930TA418 .9.F5 A38 1999 620.1921--dc21

99-046026

Copyright @) 2000 by Kluwer Academic Publishers . All rights reserved . No part of this publication may be reproduced, stored in a retrieval system or transmitted in any form or by any means, mechanical, photocopying , recording , or otherwise, without the prior written permission of the publisher, Kluwer Academic Publishers, 101 Philip Drive, Assinippi Park, Norwell, Massachusetts 02061 Printed on acid-free paper. Printed in the United States of America

TABLE OF CONTENTS SECTION I. INTRODUCTION F. T. Wallenberger 1 1.1 1.2 1.3 1.4

FIBERS FROM THE VAPOUR PHASE The most important phase isthe liquidphase Afibre by any name isstill afiber Biographic Sketches of the authors Acknowledgements

3 3

4 6 7

SECTION II. FIBERS FORM THE VAPOUR PHASE F. T. Wallenberger

2

2.1

2.2

2.3 2.4

SHORT FIBERS, WHISKERS AND NANOTUBES Advanced vapor phase processes 2.1 .1 Evolution ofa technology 2.1 .2 Crystal growth and phase transformations (a) Vapor-liquid-solid (VLS) growth (b) Vapor-solid (VS) growth 2.1 .3 Metal catalyzed chemical vapor deposition (a) Reaction chemistry (~ Controlled whisker growth (c Whisker morphology (~ Generic whisker properties 2.1.4 Laser ablation ofwhisker precursor alloys 2.1 .5 Hot fiber chemical vapor deposition 2.1 .6 Chemical vapor infiltration 2.1 .7 Carbothermal reduction (a) Pyrolytic processes (b) Chemical mixing processes (c) Self-propagating high-temperature synthesis 2.1 .8 Plasma and related processes (a) Arc discharge processes (b) Laser vaporization and ion bombardment Advanced liquid phase processes 2.2.1 Self-assembly ofsilver nanowires 2.2.2 Whiskers from organic solvents 2.2.3 Whiskers from mesopitch Advanced solid phase processes 2.3.1 Micropillars by lithography and etching Selected fiber structures and properties 2.4.1 Silicon whiskers and nanowhiskers 2.4.2 Silicon carbide whiskers and nanowhiskers

11 11 11 12 12 13

15 15 18

19

20 20

21 22

23 23 23 24 24 24

25

26 26

27 28

29 29 30 30 34

vi

2.4.3 2.4.4

2.5

Short graphite fibers Carbon nanotubes (a) Structures (b) Properties Selected fiber products and applications 2.5.1 Silicon whiskers and nanowhiskers 2.5.2 Silicon carbide whiskers and nanowhiskers 2.5.3 Short carbon and diamond fibers Short carbon fiber composites Diamond/carbon fiber composites 2.5.4 Carbon nanotubes

36 36 37 39 39 40 41 41 41 42

CONTINUOUS OR ENDLESS INORGANIC FIBERS Continuous vapor phase processes 3.1 .1 Laser assisted chemical vapor deposition (a) The rceneric process concept (~ The ow pressure process (c The high pressure process (~ Automatic process control 3.1.2 Conventional chemical vapor deposition (a) Commercial hot filament CVD process (b) Experimental CVD and PVD processes 3.1 .3 Chemical vapor infiltration processes (a) CVI ofcarbon fibers with silicon oxide (bj CVI ofboron oxide fibers with ammonia (c CVI ofpolyborazine fibers with ammonia 3.1.4 Laser vaporization ofcarbon-metal mixtures Selected structures and properties 3.2.1 Hit and low pressure LCVD fibers (a Reactor pressure vs. growth rate (~ Tip temperature vs. ~roperties (c Side growth versus ipgrowth ~ ~ Versatility versus whIsker processes 3.2.2 Commercial hot filament CVD fibers (a) Sheath/core boron/tungsten fibers (bj Sheath/core versus pure boron fibers (c Sheath/core silicon carbide/carbon fibers Important CVI and PVD fibers 3.2.3 3.2.4 Structure - property commonalties (a) Straight, coiled and tubular structures (b) Fiber strength, modulus and toughness Selected products and applications 3.3.1 BIW and SiC/C fiber reinforced composites 3.3.2 Rapid evaluation ofnew fibers by LCVD (a) Ultrahigh temperature fibers (b) High temperature sensor fibers 3.3.3 Rapid PrototyPinp ofmicroparts by LCVD (a) Evolution 0 rapid prototyping (~ Laser chemical vapor deposition (c Photonic band-gap microstructure ~ ~ The future ofvapor phase processing

47 47 47 47 49 53 54 55 55 56 59 59 59 60 60 60 61 61 62 63 63 65 66 66 67 68 69 69 69 70 70 71 71 73 73 74 74 75 75

f~

3

3.1

3.2

3.3

34

vii

SECTION III. FIBERS FROM THE LIQUID PHASE F. T. Wallenberger with a chapter by H. Ackler and J. MacChesney

4 4.1

4.2

4.3

4.4

CONTINUOUS MELT SPINNING PROCESSES Important melt forming processes 4.1 .1 Princ~les offiber formation (a) ehavior ofviscous melts (li] Behavior ofinviscid melts (c Generic fiber forming processes Structure ofmelts and fibers 4.1 .2 (a) From melts tofibers (li] Fiber structure versus modulus (c Fiber structure versus strength Forming glass fibers from strong melts 4.2.1 Downdrawing from solid preforms (a) Structural silica fibers (li) Optical silica fibers Melt spinning from strong silicate melts 4.2.2 4.2.3 Structural silicate glass fibers (a) Product design parameters (li) General and special purpose fibers Forming glass fibers from fragile melts 4.3.1 Glass fibers from fragile silicate melts 4.3.2 Melt spinning from supercooled melts (a) Single and double crucible processes (li) Single and bicomponent fluoride fibers 4.3.3 Updrawing from supercooled melts (a) P/odrawing oftellurite ~/ass fibers (li) pdrawing ofalumina e glass fibers Hybrid fiber forming processes 4.3.4 4.3.5 Quaternary calcium aluminate fibers (a) Fiber properties (b) Potential applications Forming amorphous fibers from inviscid liquids 4.4.1 Attainment offiber forming viscosities 4.4.2 Rapid solidification (RS) processes (a) Amorphous metal ribbons (b) Products and applications 4.4.3 Inviscid melt spinning (IMS processes (a) Principles ofjetand fi erformation (b) Principles ofincreasing the jet lifetime 4.4.4 Oxide fibers from containerless, laser heated melts 4.4.5 Metal fibers in a reactive environment 4.4.6 Oxide glass fibers inareactive environment 4.4.7 Mechanism ofjetsolidification 4.4.8 Cryogenic fibers from liquefied gasses Growing single crystal fibers from inviscid melts 4.5.1 Edge defined film fed growth (a) Growth ofsapphire fibers (b) Process versatility 4.5.2 Laser heated float zone growth (a) Growth ofsingle crystal fibers

6

4.5

81 81 81 81 84 85 87 87 88 91 92 92 92 92 92 93 94 94 95 95 97 97 97 97 98 98 100 101 101 102 103 103 103 104 105 105 106 107 107 108 110 111 113 113 113 114 114 115 115

viii

4.5.3 5 5.1

5.2 6

6.1

6.2

6.3

6.4

(b) High T« superconducting fibers The future ofsingle crystal oxide fibers (a) Single crystal sapphire fibers (b) Other single crystal oxide fibers

116 118 118 119

CONTINUOUS SOLVENT SPINNING PROCESSES Dry spinning ofsilica glass fibers 5.1.1 Process concepts 5.1 .2 Pure silica fibers from water glass solutions 5.1 .3 Ultrapure silica fibers from sol-gels Silica fibers by other processes

123 123 123 124 126 128

STRUCTURAL SILICATE AND SILICA GLASS FIBERS General purpose silicate glass fibers 6.1 .1 Commercial fiber forming processes 6.1 .2 Commercial commodity 91ass fibers (a) Evolution ofborosil/cate E.glass fibers (b) Boron- and fluorine-free E.glass fibers 6.1 .3 Structures and properties (a) Mechanical properties (b) Other fiber properties 6.1.4 Commercial products and applications Special purpose silicate glass fibers 6.2.1 Hir strength - h~h temperature fibers (a Process an products (b) Properties and applications 6.2.2 High modulus - high temperature fibers 6.2.3 Ultrahigh modulus glass-ceramic fibers (a) Process and products (b) Properties and applications 6.2.4 Fibers with high chemical stability (a) Chemical resistance ofglass fibers (bJ Alkali resistant ~/ass fibers (c Acid resistant g ass fibers 6.2.5 Other special purpose fibers (a) Fibers with low dielectric constants (~ Fibers with high densities and high dielectric constants (c Fibers with very high dielectric constants (~ Fibers with super- and semiconducting properties (e) Fibers with bone bioactive glass compositions Non-round, bicomponent and hollow fibers 6.3.1 Silicate glass fibers with non-round cross sections (a) Processes and structures (b) Products and applications 6.3.2 Structural bicomponent silicate 91ass fibers (a) Sheath/core and side-by-s/de bicomponent fibers (~ HoI/ow sheath/core silicate glass fibers (c HoI/ow porous sheath/core silicate glass fibers (~ HoI/ow superconducting sheath/core glass fibers (e) Solid side-by-side bicomponent glass fibers High temperature silica glass fibers Value-in-use ofsilica glass fibers 6.4.1 6.4.2 Ultrapure silica fibers from solid preforms 6.4.3 Ultrapure and pure silica fibers from solutions 6.4.4 High silica fibers by leaching ofborosilicate fibers

129 129 129 130 130 131 132 132 133 134 136 136 136 139 140 141 142 144 145 145 146 148 149 149 150 151 152 153 153 154 154 155 156 156 156 158 158 160 162 162 163 164 165

ix

7 7.1 7.2

7.3

7.4 7.5 7.6

7.7

OPTICAL SILICA FIBERS (H. Ackler and J. MacChesney) Introduction Principles ofoptical transmission Wave guide physics 7.2.1 (a) Step index fibers (b) Graded index fibers 7.2.2 Optical loss (a) Scattering (b) Absorption 7.2.3 Dispersion Birefringence 7.2.4 Fabrication ofoptical fibers 7.3.1 Fabrication ofpreforms Doublecrucible method 7.3.2 Outside vapor deposition (OVD and VAD) 7.3.3 7.3.4 Modified chemical vapor deposition (MCVD) (a) Chem~alequmbria (b) Thermophoretic deposition and sintering 7.3.5 Plasma chemical vapor deposition (PCVD) Fiber drawing process 7.4.1 The drawing tower 7.4.2 Protective fiber coatings Sol-gel processing Applications ofoptical fiber devices 7.6.1 Optical amplifiers 7.6.2 Fiber gratings as mirrors and filters 7.6.3 Strainsensor and other applications Summary and outlook

169 169 169 169 171 172 172 173 173 174 179 180 181 181 183 185 187 189 190 191 191 192 193 194 195 196 197 198

SECTION IV. FIBERS FROM SOLID PRECURSOR FIBERS R. Naslain

8 8.1

8.2

CERAMIC OXIDE FIBERS FROM SOL·GELS AND SLURRIES General considerations 8.1 .1 The generic sol-gel process (a) The starting materials (~ The gelation step (c The drying step ~ ~ The calcination and sintering steps Alumina and alumina based fibers 8.2.1 General considerations 8.2.2 Processing ofalumina based fibers (a) Polycrystalline alumina fibers (~ Transition alumina fibers (c Mullite and related fibers (~ Alumina-zirconia fibers 8.2.3 Structure and microstructure (a) Transition alumina fibers (b) Mullite and related fibers

205 205 205 205 206 207 207 207 207 209 210 211 212 215 216 216 216

x

(c) Corundum and related fibers Mechanical properties (a) Atroom temperature (b) At high temperature 8.2.5 Physical properties Applications 8.2.6 Zirconia based fibers 8.3.1 General considerations 8.3.2 Processing of zirconiabased fibers (a) Fibers from zirconia sols (b) Fibers from polyzirconoxanes 8.3.3 Properties and applications Yttrium aluminumgamet (YAG) fibers 8.4.1 General considerations 8.4.2 Processing of YAG fibers (a) From diphasicgels fbJ From polymer precursors (c From YAG powders Properties and applications 8.4.3 8.4.4 Applications

218 21 9 219 222 224 225 225 225 226 226 226 227 227 227 228 228 228 228 228 229

CARBON FIBERS FROM PAN AND PITCH General considerations 9. 1.1 History ofcarbon fibers 9.1.2 Elemental carbon 9.1.3 Classification ofcarbon fibers Processing ofcarbon fibers 9.2.1 Principles offiber formation 9.2.2 From polyacrylonitrilebased precursor fibers (a) Nature of theprecursor (~ Spinning of PAN based precursor (c Stretchm~ (~ Stabiliza Ion (~ Carbonization ( Post heat lfeatment 9.2.3 From pitch based precursor fibers (a) Nature ofpitches (~ The carbonaceous mesophase stage fc ~inning and stabilization ~ ~ arbomzation and graphitization Structure of carbon fibers 9.3.1 Structural parameters 9.3.2 Microtexture (a) PAN based high tenacity carbon fibers fbJ PAN based hi~h modulus carbon fibers (c Mesopitch (M ?based carbon fibers Properties of carbon fibers 9.4.1 Mechanical Properties (a) Youn~ 's modulus (~ Tensl e strength fc Compressive strength ~ ~ High temperature properties 9.4.2 Thermal and electrical properties (a) Thennalexpans~n (b) Transportproperties

233 233 233 233 235 235 235 237 237 237 237 237 238 239 239 239 240 243 245 245 245 247 247 247 247 250 252 250 253 256 256 257 257 258

8.2.4

8.3

8.4

9

9.1

9.2

9.3

9.4

xi

9.5

9.4.3 9.4.4

Oxidation ofcarbon fibers Coated carbon fibers Applications

259 261 261

10

SILICON CARBIDE AND OXYCARBIDE FIBERS

265 265 266 267 269 269 270 272 272 272 275 275 275 275 276 276 276 279 280 281 283 284 284 287 287 288 291 291 293 295

10.1 General considerations 10.2 Preparation ofSi-C-O fibers 10.2.1 The Yajima process 10.2.2 Melt spinning ofPCS fibers 10.2.3 Stabilization and curing 10.2.4 Pyrolysisof PCS fibers 10.2.5 Related Si-C-O (Ti) fibers 10.3 Preparation ofoxygen-free Si-C fibers 10.3.1 From radiation cured PCS precursor fibers 10.3.2 From infusible PCS precursor fibers 10.4 Preparation ofquasi-stoichiometric SiC fibers 10.4.1 Pyrolysis of PCS precursor fibers under hydrogen 10.4.2 Pyrolysisofboron doped PCS precursor fibers 10.4.3 From extruded SiC powder/polymer mixtures 10.5 Structure ofsilicon carbide fibers 10.5.1 Silicon oxycarbide fibers 10.5.2 Silicon carbide fibers 10.6 Thermal stability ofsilicon fibers 10.6.1 Silicon oxycarbide fibers 10.6.2 Silicon carbide fibers 10.7 Mechanical properties ofSiC fibers 10.7.1 Atroom temperature 10.7.2 Athigh temfeeratures (a) Tensi e tests (b) Creep tests (c) Bend stress relaxation test 10.8 Oxidation ofsilicon carbide fibers 10.9 Transport properties ofSiC fibers 10.10 Applications 11

11.1 11.2

11.3

SILICON NITRIDE AND BORIDE BASED FIBERS General considerations Si-C-N-O and Si-C-N fibers 11 .2.1 Processing (a) From porvsilazane (PSZ) fibers (b) From poycarbosilazane (PCSZ) fibers 11.2.2 Structure and properties (a) Fiber structure (~ Thermal stability (c Mechanical pro erties ( ~ Oxidation resisrance (e) Other properties Si-N-Oand Si-N fibers 11 .3.1 Processing (a) From Yajima type polycarbosilane reS) fibers {bj From pemydrOtO%SilaZane (PHPS Jfibers {c From other poysiazane fibers

299 299 299 299 299 300 301 302 302 304 304 306 306 306 306 307 308

xii

11.3.2

11.5

Structure and properties (a) Thermal properties (bJ Mechanical properties (c Other properties Si-B-O-N, Si-B-N and Si-B-N-C fibers 11.4.1 Processing (a) From perhrodropogsilazane (PHPSZ) fibers (b) From trich orosilyamino-dichloroborane (TADB) fibers 11.4.2 Structures and properties (a) Structure and thermal stability (b) Mechanical properties Applications

308 308 309 309 309 309 310 311 311 311 311 311

12

APPLICATIONS OF CARBON AND CERAMIC FIBERS

315 315 316 316 320 322

ACRONYMS GLOSSARY

331 335

INDEX

341

11.4

12.1 Fiber applications 12.2 Composite applications 12.2.1 Polymer matrix composites 12.2.2 Metal matrix composites 12.2.3 Carbon and ceramic matrix composites

SECTION I INTRODUCTION F. T. Wallenberger This book serves as an introduction to advanced inorganic fibers and aims to support fundamental research, assist applied scientists and designers in industry, and facilitate materials science instruction in universities and colleges. Its three main sections deal with fibers which are derived from the vapor phase such as single crystal silicon whiskers or carbon nanotubes, from the liquid phase such as advanced glass and single crystal oxide fibers, and from solid precursor fibers such as carbon and ceramic fibers.

Contents FIBERS FROM THE VAPOR, LIQUID AND SOLID PHASE

1.1 1.2 1.3 1.4

The most important phase isthe liquid phase Afiber by any name is still a fiber Biographic sketches ofthe authors Acknowledgments

CHAPTER 1 FIBERS FROM THE VAPOR, LIQUID AND SOLID PHASE F. T. Wallenberger The book describes advanced inorganic fibers, focuses on principles and concepts, analyzes experimental and commercial processes, and relates process variables to structures, structures to fiber properties and fiber properties to end-use performance. In principle, there are discontinuous or inherently short, and continuous or potentially endless, fibers. Short fibers range from asbestos fibers, which were described as early as 300 BC to carbon nanotubes which were discovered in 1991 [1] and have been fully described in 1999 [2]. Continuous inorganic fibers range from silicate glass fibers which were reported in 1630, to vapor grown boron fibers which were reported in 1995 [3], single crystal germanium fibers [4] and amorphous yttrium aluminum garnet fibers [5] which were reported in 1998. Even continuous cryogenic hydrogen and argon fibers [6] were recently reported. 1.1 The most important phase is theliquid phase Some short as well as discontinuous fibers can be grown from the vapor phase, some formed from the liquid phase, either a viscous melt or a viscous solution, and others yet are derived from a solid precursor or green fiber. The present book subordinates the discussion of individual orgroups of fibers tothe functional hierarchy of these process concepts. Almost all short fibers, which are derived from the vapor phase, grow by a vapor-liquid-solid (VLS) phase transformation, including single crystal silicon whiskers and carbon nanotubes. Only rarely does the growth of short fibers or whiskers occur by a vapor-solid (VS) phase transformation, and the evidence for this type of phase transformations is often difficult to obtain experimentally. Silicate glass fibers, which command by far the largest sales volume in the market, are derived from the liquid phase, a viscous melt. Ultrapure silica fibers are either derived from a melt, which isdowndrawn from a preform, orthey are dry spun from a viscous solution. While a high melt or solution viscosity seems to be a general prerequisite for fiber formation, it is possible to form glass fibers, such as YAG oraluminate glass fibers, from melts. Carbon fibers, as well as ceramic oxide and carbide fibers, which have a combined sales volume well below that ofglass fibers, are derived from solid precursor orgreen fibers. These precursor fibers are in turn derived from a liquid phase, e.g., from a viscous solution orfrom a viscous polymer by dry or melt spinning, respectively. By virtue of its organization, this book is uniquely able to pay equal attention to the formation, structures and properties of the

4

Chapter 1

functional carbon and ceramic fibers and to those of the nonfunctional green or precursor fibers from which they are derived. Not only is a liquid phase a key process element during the formation of nearly all inorganic fibers, but it must have a solution ormelt viscosity oflog 2.5(or 316) to log 3.0(or 1000) poise atthe forming temperature. Even the viscosities of YAG oraluminate compositions which are only 200 IJm are generally known as rods. Fibers with diameters ranging from 100 to 200 IJm are either known as very large diameter structural fibers (e.g., boron/tungsten fibers) or optical fibers (e.g., silica fibers). Fibers with diameters ranging from 1 to 25 IJm are generally known as microfibers, and those with diameters generally ranging from 1to 25 nm are known as nanofibers.

Chapter 1

5

Table I. Inorganic fibers from the vapor phase, the liquid phase and from solid precursor fibers Fiber Compo- Process Precursor Fiber length form nents status phase Rods> 200 um diameter Silicate glass rod endless solid one comm . liquid/melt Silicate glass tube endless hollow two comm . liquid/melt exper . liquid/melt rod short solid one Superconductor Coarse fibers> 100 urn diameter comm . vapor phase filament endless solid two Boron/tungsten waveguide endless solid two comm . liquid/melt Silica - optical fiber endless solid two comm . vapor phase Silicon carbide/C fiber endless solid one exper. liquid/melt Superconductor Fine fibers >1 urn diameter fiber endless solid one comm . liquid/disp. Alumina - cryst. fiber endless solid one comm . liquid/melt Aluminate - glass Aluminum fiber endless solid one exper . invsicid melt fiber short solid one exper . liquefied gas Argon - cryogenic Boron fiber endl ess solid one exper. vapor phase one comm. vapor phase Carbon - structural fiber short solid Carbon - structural fiber short solid one exper . liquid/sol'n Carbon - structural fiber endless solid one exper . vapor phase Carbon - structural fiber endless solid one comm . solid fiber Diamond / tungsten fiber endless solid one exper . vapor phase Fluoride fiber endless solid two exper . liquid/melt fiber endless solid one exper. vapor phase Germanium - s.c. Hafnium carbide whisker short solid one exper. vapor phase ribbon endless solid one comm. liquid/melt Metal alloys Oxide - s. crystal fiber endless solid one comm . liquid/melt Oxide - ceramic fiber endless solid one comm . solid fiber Oxide - glass fiber endless solid one exper. inviscid melt Silica - structural fiber endless solid one comm . liquid/melt Silica - structural fiber endless solid comm . liquid/sol'n one Silica - structural fiber endless solid one comm . solid fiber Silica - structural fiber endless solid one comm . solid fiber Silicate - glass fiber endless solid one comm. liquid/melt hollow fiber endless hollow two comm . liquid/melt Silicate - glass Silicate - glass porous fiber endless porous two exper. liquid/melt whisker one comm . vapor phase Silicon - s. crystal short solid Silicon - s. crystal micropillar short one exper . bulk solid solid Silicon - amorphous fiber endless solid one exper . solid fiber Silicon carbide whisker short solid one exper. vapor phase Silicon carbide fiber endless solid one exper. vapor phase Silicon carbide fiber endless solid one comm. solid fiber Steel - melt spun fiber endless solid one exper. inviscid melt Superconductor fiber endless hollow three exper . liquid/melt YAG - s. crystal fiber endless solid one comm . liquid/melt fiber YAG - polycryst. endless solid one exper. solid fiber YAG - amorphous fiber endless solid one exper . liquid/melt Zirconia - cryst. fiber endless solid one ~.~uid / sol'n Nanofibers >1 nm diameter Boron nitride nanotube short hollow two exper . vapor phase Carbon fullerene pipe endless hollow two exper. vapor phase Carbon nanotube short hollow two exper. vapor phase Carbon fullerene rope endless hollow two exper . vapor phase Carbon nanotubule hollow two short exper . vapor phase Silicon nanowire short solid one exper . vapor phase Silicon quantum wire short one solid exper . vapor phase Silicon carbide nanowhisker short solid one exper. vapor phase Silver/DNA nanowire short two solid exper. liquid/sol'n Composition (alphabetical)

Trival fiber name

Chap . #

2.2.2 3.2.2 7.3.2 3.2.2 4.5.2 8.4.1 4.3.3 4.4.6 4.4.6 3.2.1 2.5.4 2.4.3 3.2.1 9.2.1 2.4.6 4.3.2 3.2.1 2.6.2 4.4.2 4.5.1 8.1.2 4.4.7 6.4.1 6.4.3 6.4.2 6.4.4 6.1.2 6.3.2 6.3.2 2.6.2 2.3.2 3.2.1 2.3.3 3.2.1 10.2.1 4.4.6 6.3.2 4.5.2 8.5.2 4.4.4 8.4.2 2.2.7 2.2.7 2.2.7 2.2.7 2.2.7 2.2.7 2.2.7 2.6.3 2.3.1

6

Chapter 1

Specialty fibers have a non-round, e.g., ribbon, dumbbell, or trilobal cross section, and bicomponent fibers often consist of two concentrically arranged materials, a core of one material or a hollow core, and a sheath or cladding ofa second material. Furthermore, fibers with round and non-round cross-sections, and fibers with bicomponent structures, can either be continuous (practically endless) or discontinuous (short). In general, the latter have high aspect ratios (length divided by diameter ratios >1000). Experimental fibers, which are made by potentially continuous processes, have been counted as potentially continuous fibers in groups of chemically related fibers are known by their generic names, e.g., boron fibers, silicon whiskers or carbon nanotubes. Fiber producers often protect the commercial identity, quality and reliability of their products with tradenames; e.g., Dacron is a registered trademark owned by Du Pont toprotect its polyester fibers. Individual mountain climbers name the peak which they scaled for the first time, and scientists often give trivial names to fibers they synthesized for the first time, for example, fullerene pipes, buckytubes, or quantum wires. Such nicknames reflect the sense of discovery that always prevails in a new and rapidly growing field . In the scientific context of this book, generic and specific compositions are more meaningful than trade names. Therefore, compositional descriptions have been used throughout the book to characterize a given fiber, notably indiscussions which relate structures to properties. Trade names have been used only when absolutely necessary. No attempt has been made to suppress the diversity of trivial names as they appear inthe literature. A fiber byany name is still a fiber. In summary, the book introduces a unified view of advanced inorganic fibers to the aspiring materials science student and attempts to foster cross-fertilization among the experts inthe field .

1.3 Biographic sketches ofthe authors Fred Wallenberger is an expert in the fields of inorganic, polymer-organic and natural fibers. He got his Ph.D. degree from Fordham University, was a Research Fellow at Harvard, and joined the staff of Pioneering Research Laboratory, Du Pont Fibers, where he contributed for over three decades to the commercialization of new fibers through intrapreneurial research, project management and technology transfer. Subsequently, he became a Research Professor at the University of Illinois in Urbana-Champaign and a Visiting Professor at the University ofCalifornia in Davis, and assisted entrepreneurial high technology businesses with organizational advice, technical assistance and license negotiations. He has published over 100 papers, several in Science, and recently joined the staff of the Fiber Glass Research Center, PPG Incorporated as Manager, Advanced Technology. Roger Naslain received his Ph.D. degree at the University of Bordeaux and spent one year at the General Electric Corporate Research Center in Schenectady, NY. Since then he pursued research in composite materials, first as a group leader at CNRS Laboratory of Solid State Chemistry, and then as a manager of the Institute of Composite Materials, a technology transfer center. He is now manager of the Laboratory for Thermostructural Composites located in the Bordeaux area, and professor of materials science at the University of Bordeaux. He has published more than 200 papers, has received fifteen patents and has edited several books inthe field offiber and composites technology.

Chapter 1

7

John MacChesney is a materials scientist with a BA degree from Bowdoin College and a Ph.D. degree from the Pennsylvania State University. He has spent his professional career at Bell Laboratories working on glasses for electronic oroptical use. A pioneer infiber optics, he is credited with the invention of the MCVD process to make fibers, and is a principal in the development of sol-gel silica for fiber use. He has published about 100 papers and an equal number of patents. He is a Fellow of Bell Laboratories, and a member of the National Academy of Engineering. Harold Ackler, Lawrence Livermore National Laboratory, got a MS degree in Materials Science and Engineering at the University of California in Berkeley, and a Ph.D. degree in Ceramics at the Massachusetts Institute of Technology. His research, while at Bell Laboratories, Lucent Technologies, has focused on the processing of optical fiber preforms via sol-gel methods, planar photonic devices, and glasses with non-linear optical properties.

1.4 Acknowledgments Dr. Wallenberger gratefully acknowledges the initial encouragement for writing the book from the late Norman Kreidl, pioneer and teacher; James Nottke, formerly Director, Pioneering Research Laboratory, Du Pont Fibers; and Zhao Jiashiang, Director, Beijing Research Institute for Materials and Technology. He enjoyed the searching discussions with Eugene Givargizov, head of the Crystallography Laboratory, Russian Academy of Sciences in Moscow; Paul Nordine, President, CRI in Evanston IL; Gary Tibbetts, Staff Scientist, General Motors Research in Warren, MI; Robert Feigelson, Stanford University; Austen Angell, Arizona State University; and with Norman Weston, Consultant, Lewes, DE. The completion of the book would not have been possible without the support from Jaap van der Woude, formerly Director, PPG Fiber Glass Research Center, and with that of Norman Weston who reviewed the manuscript and prepared the appendix. Professor Naslain is grateful for valuable advice, data and illustrations from I. Mochida, Kyushu University; A Oberlin, J. B. Donnet and X Bourrat, CNRS; J. L. White, University of California, San Diego; R. J. Diefendorf and D. D. Edie, Clemson University; J. C. Lewis, Union Carbide; J. Economy, University of Illinois; H. Ichikawa, Nippon Carbon; 1. Yamamura, Ube Industries; J. Lipowitz, Dow Corning; K. Okamura, Osaka Prefecture University; R. M. Laine, University ofMichigan; M. D. Sacks, Universityof Florida; H. P. Baldus, Bayer AG; J. DiCarlo, NASA-Lewis; R. Tressler, Penn State University; A. R. Bunsell, Ecole des Mines, Paris; P. Olry, SEP/SNECMA; J. Dunogues , University ofBordeaux and R. Pailler, CNRS. REFERENCES [1} [2J [3J [4} [5} [6J

S.lijima, Carbon nanotubes, Nature, 354, 56(1991). L. C. Venema, J. W. C. Wildoer, J. W. Janssen , S. J. Tans, H. L. J. Temminick Tuinstra, L. P. Kouwenhoven and C. Dekker, Imaging electron wave functions as quantized energy levels in carbon nanotubes, Science, 283, 52-55 ([999). F. T. Wallenberger, Rapidprotoptying directly from the vapor phase, Science, 276, 1274-1275 (1995). F. T. Wallenberger and P. C. Nordine, Potentially continuous single crystal germanium fibers bylaser assisted chemical vapor deposition, in preparation (1998). J. K. Weber, J. J. Felton, B. Cho and P. C. Nordine, Glass fibres of pure and erbium- or neodymium-doped yttria-alumina compositions, Nature, 393, 769-771 (1998). R. Aliaga-Rossel and J. Bayley, A cryogenic fiber maker forcontinuous extrusion, Rev. Sci. Instrum., 69 [6}, 2365-2368 (1998).

8

(7] [8) [9]

Chapter 1

M. S. Dresselhaus, G. Dresselhaus and P. C. Eklund, Science of Fulferenes and Carbon Nanotubes, Academic Press, San Diego, CA (1996). A.Kelly, Editor, Concise Encyclopedia of Composite Materials, Pergamon, London (1994). V. L. Kostikov, Editor, Fibre Science and Technology, Chapman & Hall, London (1995).

[10] P. W.Johnson, Ceramic fibers and coatings, advanced materials for the twenty-first century , Public ation NMAB-494, National Academic Press, Washington, DC (1998) .

SECTION II FIBERS FROM THE VAPOR PHASE F.T. Wallenberger Advanced inorganic fibers fall into two categories: (1) discontinuous or short fibers and (2)continuous fibers or atleast potentially continuous fibers. Chapter 2 deals with short fibers from the vapor phase butalso introduces short liquidand solid-phase derived fibers. There isa better fit forthese fibers here than in achapter oncontinuous fibers.

Contents 2

SHORT FIBERS, WHISKERS AND NANOTUBES 2.1 Advanced vapor phase processes 2.2 Advanced liquid phase processes 2.3 Advanced solid phase processes 2.4 Selected fiber structures and properties 2.5 Selected fiber products and applications

3

CONTINUOUS OR ENDLESS INORGANIC FIBERS 3.1 Continuous vapor phase processes 3.2 Selected structures and properties 3.3 Selected products and applications

CHAPTER 2 SHORT FIBERS, WHISKERS, AND NANOTUBES Fred Wallenberger Short needle shaped, inorganic fibers occur in nature, or can be synthesized by a variety of experimental and commercial processes. If these fibers are filamentary single crystals, they are called whiskers. If however they are polycrystalline or amorphous, they are called short fibers. 2.1 Advanced vapor phase processes The recorded history of short fiber technology starts over two thousand years ago with asbestos fibers and reaches into the future with silicon nano-whiskers and carbon nanotubes. Asbestos is derived from the solid phase, but today, the most important short inorganic fibers are derived from the vapor phase. 2.1.1

Evolution ofa technology

The evolution of modern vapor phase processes starts with metal catalyzed chemical vapor deposition and ends with laser vaporization (see Table I). Most vapor phase processes require metal particle catalysts; some proceed without the addition of metal particles. The growth temperatures range from 100 to 4000·C. The length of silicon nanowires is 3700°C, graphitic carbon fibers at >2800°C, turbostratic fibers between 1400 and 900°C [7] and amorphous carbon fibers which contains the element desired in the product, e.g. , Si, as well as the metal catalyst required for whisker growth, e.g., Fe. Ablation produces a vapor of Si and Fe that quickly condenses into liquid Si-rich nanoclusters that become supersaturated. The Si phase grows and crystallizes as single crystal Si nanowire or nanowhisker by a vapor-liquid-solid growth mechanism. As produced, the nanowire has a concentric sheath/core structure with an amorphous Si02 sheath and a single crystal Si core. The sheath, which may be caused by residual oxygen in the reactor and the iron in the tip of the nanowire, can be removed by etching in hydrofluoric acid. The final bare nanowire shows only Si with traces ofoxygen by EDX analysis [74J. The nanowire length is >1~m, the outer nanowire diameter is 17.1 ± 0.3 nm and the average diameter of the crystalline core ofthe nanowire is7.8 ± 0.6 nm. Sio.9Feo / (solid phase)

laser ablation

9 Si + Fe (vapor phase) ----~ 9 Si + Fe nanoclusters (liquid phase)

crystal growth

)

(7)

Sheath / core SiO ] / Si nanowire with Fe tip (solid) ~ Single crystal silicon nanowire (solid phase)

Except for initially producing a sheath/core Si02/Si nanowhisker, the laser ablation process parallels the metal catalyzed chemical vapor deposition process (Chapter 2.2.3). In this process, the Si that is desired is generated by chemical vapor deposition and dissolved in molten metal droplets, e.g., Au or Fe. The molten alloy droplets, e.g., SiAu, which result in this process sequence, give rise to the growth of single crystal Si micro-whiskers by a similar overall VLS phase transformation. The laser ablation process has been demonstrated so far only for Si and Ge nanowires [74J, but it isclear that it isa new generic tool for growing crystalline nanowires. Thus, it should be possible to make nanowires or nanowhiskers of SiC, GaAs, BbTe3and BN in this way and perhaps, inthe presence ofatomic hydrogen, even diamond nanowires [74). 2.1.5 Hot fiber chemical vapor deposition Boron, silicon carbide, diamond and other materials can be deposited by chemical vapor deposition on the surface of hot wires or hot fibers. If a minimal vapor deposit is applied, the process will modify only the surface of the fiber and produce a coating, while leaving its core functionality unchanged. If, however, a thick vapor deposit is applied, the process will create a new and very large diameter fiber that has the functionality of the sheath and a sacrificial core. The hot fiber (Wire) CVD process has been commercially used for 30 years to produce continuous sheath/core bicomponent boron/tungsten and silicon carbide/carbon fibers. Since they are continuous fibers, they are discussed in Chapter 3.3. More recently, this process was used to produce discontinuous, i.e., short, experimental sheath/core diamond/carbon fibers by depositing a thick diamond sheath on short pieces ofa potentially carbon fiber.

22

Chapter 2

Figure 6.Growth ofdiscontinuous sheath/core bicomponent diamond/carbon fibers. Duration ofdiamond deposition at 1000·C - 12 hours (left) and 72 hours (right). Courtesy of J. M. Ting and M. L. Lake, Applied Sciences, Incorporated, Cedarville, OH

Short diamond/carbon whiskers (Figure 6), the first truly discontinuous sheath/core fibers [28], were made bya two step process. The short vapor grown carbon core fibers were produced by pyrolysis of H2/CH4 mixtures in the presence of iron catalysts [25]. These vapor grown carbon fibers were then ultrasonically polished, and diamond was deposited by a microwave plasma-enhanced chemical vapor deposition technique [28]. 2.1 .6 Chemical vapor infiltration A process that appears to proceed by a metal particle catalyzed chemical vapor infiltration and a vapor-liquid-solid phase transformation [24] was found to yield well-defined short amorphous and polycrystalline silicon nitride fibers reported to have very high strength. These fibers were up to 5 mm long, had smooth surfaces, diameters ranging from 1450·C. A process variant [9] yields SiC whiskers >1350·C in a fixed bed percolated by a hydrogen flow. The addition of iron above 1450·C affords submicron whiskers ending with a silicon rich droplet. The iron seems to evaporate and condense below 1450·C leaving behind whiskers with silicon rich tip >1450·C. These processes use the same starting materials as the rice hull processes but they also use a metal particle catalyst. As a result, they are believed to proceed by a VLS phase transformation. (b) Chemical mixing processes Silicon carbide whiskers can also be synthesized by carbothermal reduction of silicon nitride [10]. Silicon nitride decomposes >1300·C, silicon melts at 1410·C, and reacts with graphite. Whisker formation in this process is initiated >1400·C and can be completed between 1550

24

Chapter 2

and 1650·C. With the addition of a metal catalyst, distinct metal droplets were found in the tips of the whiskers [10) suggesting a VLS phase transformation, but none by VS phase transformation without the addition ofparticulate catalysts. SiJN.J (s) >130(fC ) 3 Si(s)+2 N 2 (g) (11a) Si(l)+C(s)

>140(fC

(11b)

)SiC(s)

If this carbothermal process is brought to only partial completion (Equation 11a and 11b), a homogeneous mixture of silicon carbide whiskers and silicon nitride powder [10) is obtained which can be fired directly to yield whisker reinforced ceramics. Silicon carbide reinforced alumina composites and silicon carbide whisker reinforced zirconia composites [31) are also products of the "chemical mixing process". The whisker growth rate in the zirconia process can be accelerated by adding metal particle catalysts such as cobalt chloride, thus potentially facilitating a VLS phase transformation.

(c) Self-propagating high temperature synthesis Very pure single crystal SiAlON whiskers [75) were recently made by an inexpensive selfpropagating high temperature synthesis (SHS), a process that has earlier resulted in SbN4 whiskers [75). While many routes are available for the production of silicon nitride whiskers, this seems to be the first method capable of yielding single crystal oxynitride whiskers. The synthesis is performed in a pressurized water-cooled stainless steel reactor vessel that is 1 meter long and has a capacity of30 liters [75). A homogeneous powder mixture ofsilicon (86 wl.%), alumina (8%), aluminum (1%), silicon nitride (5%) and a trace of pure ammonium fluoride in nitrogen is raised to a pressure of 100 atmospheres, and ignited. Pure single crystal SiAION whiskers with diameters of 2 I.1m are formed in the reaction wave, having a temperature >2000·C (Equations 12, a-d). (12a)

NH .J F ------+ HF + NH 3 nSi+n(a I 2-b I 6)N 2 +(b l 3)NH 3

~(SiNuHh )n

(12b)

AI 203 + Al ------+ 3Ato

(12c)

(3 - z )(SiNuHh)n + z AIO ~ nSi3 _ zAI:0:N.J_: +

(12d)

[an(3 - z) -( 4 - z)]NH J + 112[bn(3- z) - 3an(3 - z) + 3(4 - z))H 2

Some whiskers have a hollow, tubular structure; none have metal droplets at their tips. This and other factors suggest [75) that whisker growth does not proceed by a VLS, but by a VS transformation whereby the ammonium fluoride catalyst plays an important role in the growth process. 2.1 .8 Plasma and related processes Carbon and graphite fibers with diameters of 0.3-3.0 I.1m, and multishell carbon and silicon nitride nanotubes with diameters of 3-20 nrn, have been shown to grow by metal catalyzed chemical vapor deposition. Carbon nanotubes will also grow by arc discharge, carbon ion bombardment and laser discharge processes. In each case, there is the option of adding metal catalysts to the process, thus facilitating a more controlled VLS phase transformation and therefore amore uniform product with higher yields. (a) Arc discharge processes

Chapter 2

25

Carbon nanotubes were discovered in 1991 by a carbon arc discharge method [11] whereby a DC current of 150 Alcm 3 is applied with a voltage set at 20 V inhelium at a pressure of 50 torr [17]. The arc discharge is generated at >3700°Cbetween two carbon rods, 1 mm apart. The positive electrode is consumed and a complex deposit forms at the negative electrode. The outer hard shell of the deposit is removed, and its soft core contains aligned bundles of 10 to 100 multishell nanotubes. The bundles are separated by sonication in alcohol [17]. The individual multishell nanotubes have outer diameters ranging from 2 to 20 nm. Carbon nanoparticles remain present after sonication and are selectively burned away in the presence of oxygen, a step that also consumes a portion of the nanotubes [17]. This basic arc discharge process gave only about a 1-% overall yield ofcarbon nanotubes. Carbon nanotubes are also formed in an arc discharge process, when carbon black [32] or graphite [33] is covered with transition metal nanoparticle catalysts. Multishell boron nitride nanotubes with inner diameters of 1-3 nm and lengths of 200 urn) diameter carbon fibers [8] with low growth rates «10 ~m/s) when a C02 laser is used. Small «20 urn) diameter carbon, boron and silicon fibers [1-2] [4] [7] are produced with equally low growth rates at C,H, SiH,/NH,

Fiber diameter, um Low Avg. 6 19 10 63 59 70 15 45 13 120 21 45

Growth Rate, um/s Avg. High 625 1100 125 331 18 35 460 500 75 125 338 740

54

Chapter 3

The experimental high pressure process provides cwpower up to 270 mW. The beam from a one watt, TEM oo mode, cwNd-YAG laser (emission wavelength 1.064IJm) is passed through a polarizer, a Linconix laser power stabilizer, a variable neutral density filter, a beam expander, and focused with a 10 em focal length lens into the fiber growth reactor [2] [12]. Laser power ranges from 0 to 200 mW with a stability of one mW, as measured outside the reaction chamber. The laser beam was focused onto a point inside a reaction chamber, where localized heating promotes vapor deposition in the direction of the laser. The reactor could be moved parallel and orthogonal to the laser beam direction. Motion parallel to the laser beam was driven at selected rates by a computer operated by a stepper motor. Growth of the fiber at a location where the laser beam converges toits focal pointbecomes self-regulating [2]. (d) Automatic process control In the experimental high pressure process, the gas flows, reactor pressure, and laser power are set atconstant values and fiber growth is initiated by opening a shutter on the laser beam. The fibers grow from the carbonized edge of a paper substrate held inside the reactor by a removable probe. Fiber growth occurs spontaneously when the laser beam is turned on or when the edge of the paper substrate is drawn near to the laser focal point by moving the reactor. Growth of the fiber at a location where the laser beam was converging to its focal point provided an intrinsic method forrate control in vapor-solidgrowth [2). Thus, if the deposition rate exceeds the reactor translation rate, the fiber tip would grow to a point further from the beam waist, thereby cooling the tip and reducing the deposition rate. If the deposition rate is less than the growth rate, the fiber tip will bedrawn closer to the beam waist where its temperature is higher and the growth rate increases to match the pulling rate. The fiber therefore adjusts its position to achieve a temperature that made the pulling and growth rate equal. It is evident that the HP-LCVD process affords considerable latitude in the design and execution of fiber growth and kinetic studies. The discovery of the self-regulating mechanism made it possible tofabricate uniform two meter long fibers [2] bya semi-continuous version of the batch process (Figure 6). The results confirm that a continuous process is feasible in the high pressure regime, and should yield optimal growth rates exceeding 1 mm/s or about equal to the rate by which commercial sapphire fibers are produced by a flux method (see Chapter 6). Recently, a method was described for the real-time measurement of growth rates and feedback control of three-dimensional laser assisted chemical vapor deposition [11]. This method allows the accurate reproduction of high quality films, fibers, and three-dimensional structures. High aspect ratio axisymmetric forms of desired shape and microstructure were grown from vapor phase precursors by this method. Three-dimensional rods, cones, hyperboloids, and spheroids ofpyrolytic graphite, nickel, iron, and nickel-iron superalloys were obtained from ethylene, nickel tetracarbonyl, iron pentacarbonyl, and mixtures of nickel and iron carbonyls, respectively. To control the process [11), a measure of the volumetric growth rate was obtained from specific emission spectra generated during the heterogeneous reaction, and direct feedback control of the reaction rate was realized byusing this growth rate to modulate the laser power in real time. By this feedback method, layered and continuous prototyping is possible on a microscale since real time compensation for growth rate perturbations can be made. The

55

Chapter 3

study was carried out atpartial pressures with growth rates up to 45 IJm!s. While the process is potentially continuous even at partial pressures, the growth rates are too low to be of commercial significance. 3.1 .2 Conventional chemical vapor deposition Mass transfer in metal catalyzed and in laser assisted CVD processes is driven by highly localized temperature gradients. The relatively small area ofeither a hot molten metal particle or of a hot laser focus affords whiskers [4] or continuous fibers, respectively [2] [18-19]. The transfer of an equal mass from the vapor to the solid phase in a conventional chemical vapor deposition results in a thin coating over the relatively large area of a hot surface, i.e., that of a flat complex shaped composites part.

1-----Laser beampath

Centralcavity Mechanism to support fiber growth

• '-"-"-"-r'!

F'--: .;;,:._ •• ...J

"-"L_,

ii :

, I

._.._.._.. _!._~ Reactant inlet

Outlet

Figure 6. Drawing of the high pressure LCVD reaction chamber. Courtesy of Dr. P. C, Nordine, Container1ess Research Inc.. Evanston, IL.

(a) Commercial hotfilament CVD process

Conventional chemical vapor deposition produces a coated fiber when a thin coating is uniformly deposited by this method over the fiber surface. Such a coating affords an insignificant diameter increase of the resulting fiber and the functionality of the product continues to be that of the coated fiber. It merely provides an enabling, e.g., an oxidation resistant, function. Chemical vapor deposition, however, produces a large diameter sheath! core fiber, when a thick coating is uniformly deposited over the surface of a small diameter fiber (Figure 7). The fiber does not grow by lengthening (i.e., directional growth) but by thickening (i.e., side growth). The diameter of the fiber increases in this process to up to 10x without change in its length, and the functionality changes from that of the core to that of the sheath.

56

Chapler3

Continuous boron/tungsten fibers were the first high performance fibers tobe designed, about 40 years ago and commercialized about 30 years ago to meet the demanding end use requirements for resin and metal matrix composites in aircraft and sporting goods markets. Using the same process, silicon carbide/carbon fibers were also commercialized. Both fibers continue to represent important niche products. Experimental boron/carbon and silicon carbide/tungsten fibers were also developed. The structures, properties, and applications of these fibers are discussed inChapter 3.2.2. Boron/tungsten (BIW) fibers are produced by vapor deposition of boron (Equation 1) on the surface of a practically endless, electrically heated tungsten filament having a diameter of 121Jm. The reaction chamber (Figure 7) is a closed system with mercury seals on both ends and two segments [24-25]. The surface of the tungsten filament isdecontaminated by heating it to ~1350°C [30] in the first (short) segment in a reducing atmosphere. The cleaned tungsten filament and the reaction mixture of boron trichloride and hydrogen are then passed through the second (long) segment, the reactor. The electrically heated tungsten filament causes boron trichloride to decompose and deposit boron on the filament surface. Since each boron fiber requires its own reaction chamber, commercial production requires literally hundreds ofreaction chambers which are linked toa common gas supply, a gas mixer and gas regenerator. An individual reactor may be about 2 mlong. The single individual high density tungsten filament enters the reactor with a diameter of 12 IJm and a low-density boron/tungsten fiber with a diameter of 100 or 140 IJm exits. Since it passes through the reactor within a minute or two, the throughput may be as high as 1.0 m/min. The materials cost for the sacrificial tungsten wire dominates the overall cost of manufacture. A less expensive carbon filament was used for awhile as a substitute for the tungsten wire. Figure 8 contrasts a boronltungsten fiber with a pure boron fiber. Aside from the ~10x difference in diameter, the differences in surface texture are noteworthy. The surface of the pure boron fiber made by high pressure LCVD is smooth. Its strength is 7.5 GPa and its modulus is 400 GPa. In contrast, the surface texture of the boronltungsten fiber is "nubby". Its strength is3.6 GPa and its modulus is 400 GPa. In summary, the tensile strength ofboron fibers isrelated totheir surface uniformity. Silicon carbide/carbon (SiC/C) fibers are also commercially produced by side growth in a single stage hot wire CVD process, using silane or tetrachlorosilane, hydrogen and methane as the reactants [25]. A good balance between vapor phase reaction and deposition ofsilicon carbide from the vapor phase is obtained when the carbon filament is resistively heated to 1200°C. The bicomponent fiber reaches 140 IJm in this process, has a density of 3.0g/cm 3, strength of3.45 GPa, and modulus of400 GPa. Apyrolytic graphite coating may be applied to the carbon fiber to modify the interface between the core and the sheath, and various commercial coatings, designated as SCS, may be applied to modify the outer surface of the silicon carbide sheath. A tungsten wire can be used as a commercial alternative instead of the carbon fiber core. (b) Experimental CVD and PVD processes This subchapter deals with experimental sheath/core fibers made by conventional as well as plasma enhanced chemical vapor deposition, and by plasma enhanced physical vapor deposition.

57

Chapter 3

Payout Pyroolytic Graphite deposition section

H2 --+J"-+...,

o a. B E

20

Tungsten

g

:5c

Boron deposition section

40

~

tii

§ '0 C

-

~

c !!! 1i>

~

'iii c

~

~

240 kPa «20 psig) 1.38

200 0

200

400

600

800

lCVD chamber pressure

Figure 10. Effect of reactor pressure on the strength of boron fibers. Redrawn from F. T. Wallenberger, P. C. Nordine and M. Boman. Inorganic fibers and microstructures directly from the vapor phase, Composites Science and Technology,S, 193-222 (1994).

62

Chapter 3

The second example (Table II) deals with LCVD silicon fibers. In the high pressure LCVD process, single crystal silicon fibers were obtained with tip temperatures above 1400°C, high growth rates (>500 IJm/s) and VLS phase transformation. And, polycrystalline silicon fibers were obtained with tip temperatures between 600 and 1400°C and intermediate growth rates (12-500 IJm/s) under conditions where VLS or VS transformations could occur. Amorphous silicon fibers were obtained with very low tip temperatures (525°C), low growth rates and VS phase transformation. In the LP-LCVD process, single crystal silicon fibers were obtained with high tip temperatures and low growth rates (~1 IJm/s) and polycrystalline silicon fibers with lower tip temperatures, low growth rates (::;1 IJm/s) and VS. Table II. Growth of High and Low Pressure LCVD Silicon Fibers [2] [4] [5] [14] [15] TipT.,OC HPandLP >1400 1400-600 >525

Rate. um/s HP LP >500 ~1 >12 :51

1

Fiber morphology HPandLP Single crystal Polycrystalline Amorphous

Strength. GPa

HP

HP

4.0

3.0-6.7

Modulus. GPa HP LP 180 164 130

The relationship between process variables (e.g., tip temperature) and structures is reminiscent of that which governs the growth of metal particle catalyzed chemical vapor deposition (Chapter 2). It follows traditional patterns. Single crystal silicon fibers have relatively high strength and relatively high stiffness. Polycrystalline fibers have lower stiffness (modulus) than single crystal silicon fibers. Single crystal silicon fibers made by the lowpressure process were occasionally found to have a polycrystalline overgrowth . The latter serves as a stress riser and is responsible for the variable strength levels which were observed. (b) Tip temperature vs. properties

The third example (Table III) deals with HP-LCVD carbon fibers and illustrates the same overall relationships [13] [16]. Depending upon growth conditions and feed gas chemistry [16], these fibers were very strong and graphitic when formed at high tip temperatures; thickened and brittle when formed atintermediate tip temperatures; orvery flexible and elastic when formed at low tip temperatures. Graphitic LCVD carbon fibers had the highest strength (3.0 GPa) and modulus (::;180 GPa), and flexible carbon fibers the lowest strength (::;0.4 GPa) and modulus « 30 GPa). Flexible high pressure LCVD carbon fibers could be readily bent toradii with curvatures of ::;1 mm. The force required was much lower than that required for (a) equal diameter HP-LCVD boron fibers having a modulus of >275 GPa [12] or (b) equal diameter intermediate modulus or 1M carbon fibers having a modulus of250 GPa [16]. These qualitative relationships parallel those observed for silicon fibers, where single crystal fibers were formed with the highest tip temperatures and had the highest strength and stiffness, and where amorphous fibers were formed atthe lowest tip temperature and had the lowest strength and modulus. Table Carbon Fiber Type Commercial 1M Fiber "Graphitic" Fibers "Brittle" Fibers "Flexible" Fibers

m.

Mechanical Properties of HP -LCVD Carbon Fibers Strength, GPa - 3.5 0.5-3.0 NA 0.2-0.4

Elongation, % - 1.5 >1.5 NA

Modulus, GPa -250.0 28-180 NA 6 IJm and a strength levels 7.6 GPa, i.e. , with 1.6x the maximum strength at 0.06-0.04x the diameter of the former. High specific properties (strength or modulus divided by density) are

Chapter 3

67

very important properties since nearly all transportation composites are weight sensitive. In these terms (Figure 13), average low diameter single component boron fibers were nearly as strong and stiff as VLS SiC whiskers. Their average specific strength was 1.2x that of commercial 1M carbon fibers, 1.7x that of commercial boron/tungsten fibers and 2.1 x that of commercial Nicalon SiC fibers. Their average specific modulus was 2.3x that of Nicalon SiC fibers, and comparable tothat of 1M carbon orboron/tungsten fibers. Finally, pure boron fibers grow by tip growth or lengthening and therefore have a smoother surface and higher strength than boron/tungsten fibers, which grow by side growth or thickening. Carbon whiskers made by metal particle catalyzed CVD, may serve as an analogy. Initially they grow by tip growth, but an additional carbon sheath is obtained by side growth, or thickening. Tip-grown carbon fibers are stronger than tip-grown carbon fibers with a secondary, side grown carbon sheath. Also, the temperature in the small laser focus is more readily controllable than that of a practically endless hot wire, i.e. another factor favoring tip over side growth. Specific modulus, 108 in 2.4

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-

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Glass &:lZllass ceramic . fibers ..... Strong high high

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high

high

.....

Fragile

high

low

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Fragile

high

low

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Inviscid

.....

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Metal ribbons & metal fibers ..... ..... very Inviscid low ..... ..... very Inviscid low

metal alloy ribbons steel, other metal fibers

Generic fiber forming processes

Environment

downdrawing >LT downdrawing >LT downdrawing MP melt spinning >MP

by rap id solidification by jet surtace stabilization

cold quench wheel reactive gas

staI ceramic fibers

very low very , other fibers >MP fiber owth low "Liquidus temperature (LT), melting point (MP), rapid solidification (RS)process

In the scientific literature, the relationship between melt viscosity and temperature is often expressed by comparing the increase in log viscosity of a fiber forming melt with the glass transition temperature (Tg) divided by melt temperature (T), both in Kelvin (K). Among the viscous mells, this approach distinguishes between strong and fragile melts (2). Silica, the ideal strong melt, exhibits a perfect straight line relationship. Other strong melts, such as those of silicates exhibit a near straight line relationship. Fragile melts of certain aluminates or fluorides deviate noticeably from the ideal straight line behavior. The glass transition temperature, however, is nota process-related factor and this relationship is not useable in this context. The most important factor in the analysis or design of a process is the liquidus temperature (TL). The original concept that distinguishes strong from fragile melts (2) is valuable, and it can be modified by focusing on the liquidus temperature (TL), rather than on the glass transition temperature (Tg) . Thus, a linear change in mell viscosity (in poise or Pa.s) is compared, without normalizing for reciprocal temperature, with a linear change in melt temperature (in ·C orOF). Accordingly, a strong fiberforming melt has a high viscosityabove

Chapter 4

83

and below the liquidus temperature, specifically eXhibiting only small changes in viscosity above the liquidus temperature. And a fragile melt has a relatively high viscosity below the liquidus temperature [3-4] [9] [17] but a low viscosityabove the liquidus temperature. E-glass, is a commercial general-purpose fiber. It isderived from a strong silicate melt, has a log 3.0 (poise) fiber forming temperature of 1185°C, a liquidus temperature of 1050°C, and a ~ T between forming and liquidus temperatures of 135°C(Figure 1, #1). Because of its high ~ T, this melt isconsidered to be not only a strong but also a long melt. S-glass (Figure 1, #5), a commercial specialty fiber in the familyof high strength (HS) fibers is derived from a strong but short silicate melt. It has a log 2.95 (poise) fiber forming temperature of 1565°C, a liquidus temperature of 1500°C, and a ~ T between forming and liquidus temperatures ofonly 65°C(Figure 1, #2).

-

-----

10,000

LT FT FV .---

9900 9800 \ 1100-7000 C;;

cil

0-

II Q) til

1000 800 700

~

600

Q) til

500

0.

400

'0

~ 'iii 0

u

s

til

.

0 0

900

'0

0. 0

Liquidus temperature Fiber forming temperature Fiber forming viscosity Conventional melt Supercooled melt

.

0 0

1

FT\\FT ..

.. 2l.. .\ 3

j

.0

·· 6 ·.

~

0

4

0

FT

3.0113.886 3.000

0.

FT\

o'

3.991

'0

0 0

0

3.995

Q) til

0

~

FV

. ··

4.000

~ 'iii 0 o

til .:;;

01 0

...J

2.500

300 200 100

2.000

0

0.000

1100

1200

1500

1600

Temperature, °C

Figure 1. Relationshipbetween linear viscosity (10 poise = 1Pa.s) and temperature. Glass fibers requirea fiber forming viscosity (FV) ranging from log 2.5to log 3.0poise. The corresponding temperaturesare the fiber forming temperatures (FT). Glass fiberswhich are formed from strong viscous melts include Eijlass(#1) and Sij lass (#5). Their liquidus temperatures (LT), the highest temperatures atwhich crystalscan form are ::::50°Cabove their forming temperature. Glass fibers which are formed from frag ile viscous melts which have useful forming viscosities only below their liquidus temperatures include an yttria modified silicate fiber (#2) and a quaternary aluminate fiber (#3). The viscosity of these melts must be raised by supercooling before fiber formation can commence. Glass fibers which can be formed frominviscid melts include a binary aluminate (#4) and a YAG melt (110), but special viscositybuilding fiber forming processes are required toraise theirviscosity from 300 poise.

84

Chapler4

Only experimental fibers have been produced from fragile melts. Yttria modified silicate glass fibers (Figure 1, #2) have fragile and short melts [19]. They have a log 3.0poise fiber forming temperature of 1280°C, a liquidus temperature of 1390°C, and fibers can be formed 110°C below the liquidus temperature of the melt, but the molten jet must be efficiently cooled in an especially designed bushing and cooling tower. Quaternary aluminate glass fibers hsve also fragile and short melts [8-9). These melts have a liquidus temperature of 1350°C, and fibers can be formed at 1300°C, i.e., below the liquidus temperature, by updrawing them from a super-cooled and carefully stabilized melt (Figure 1, #2). When crystallization occurs with E-glass melts, the rate ofcrystal growth is low (Figure 2, #1). The crystal growth rate of S-glass mells is somewhat faster (Figure 2, #2). When crystallization occurs with a modified silicate melt (Figure 2, #3), the crystal growth rate is higher than that ofS-glass mells and with yttria modified melts it isnearly uncontrollable. Thus, fibers are best formed from strong silica and silicate melts by down-drawing them from preforms or from conventional bushings. Glass fibers can be formed from fragile melts, but they must be super-cooled to a viscosity in the log 2.5 and log 3.0poise range. Glass fibers 8.0 6.0

~ 4.0 :1.

J 2.0

o/

1

1100

1200

1300

1400

1500

Temperature, °c Figure 2. CrystallizaUon rates offiber forming melts. (1) Commercial E-glass melt, (2) commercial S-glass melt, (3) experimental silicate glass melt, (4) yttria modified silicate glass melt. Redrawn from V. E. Khazanov, Y. I. Kolesov and N. N. Trofimov, Glass fibers in Fibre science and technology, pages 15·230, Chapman and Hall, London (1995).

(b) Behavior ofinviscid melts The most common inorganic materials including oxides and metals are polycrystalline, and have high and well-defined melting points. Their melts are inviscid above the melt temperature and their melt viscosities are comparable to that of light to heavy motor oil at

Chapler4

85

room temperature. When they solidify they will almost instantly revert topolycrystalline solids. Crystallization can be prevented by increasing the quench rate without increasing the viscosity or by increasing the viscosity without increasing the quench rate. Some oxides and metals can be rapidly solidified to yield glasses by increasing their quench rate. For example, continuous metal glass ribbons can be rapidly quenched on the surface of a cold quench wheel [59), but no self-supporting glass or metal fibers have as yet been made by rapid solidification by increasing the quench rate from 104 to 10611500°C are downdrawn from the surface melt of an appropriate preform (Figure 3, top right). For details regarding fibers from strong melts see Chapter 4.2. Alumino-silicate glass fibers can be drawn from their fragile melts in an otherwise conventional fiber drawing process (Figure 3, top, left), but with extremely careful control of the cooling temperature [19]. Bicomponent heavy metal fluoride glass fibers can be drawn from their supercooled fragile melts (Figure 3, top, left) which are contained in a double crucible bushing [7). Quaternary calcium aluminate glass fibers can be updrawn from their supercooled fragile melts (Figure 3, bottom, left) which are maintained ata carefully controlled temperature [8-9). These fibers can also be downdrawn from preforms (Figure 3, top right) but stringent process control is required [41]. For detailsregarding fibers formed from fragile melts see Chapter 4.3 Continuous glassy metal ribbons can be formed with high quench rates from their inviscid melts by a rapid solidification process [60) that is akin to a generic bushing process (Figure 3, top left), except that the extruded ribbon must be rapidly cooled on the surface of a cold quench wheel. Continuous aluminate glass fibers and metal wires [10-12) and continuous amorphous YAG fibers [731 can be melt spun from inviscid melts by increasing the jetlifetime

86

Chapter 4

without increasing the quench rate. The former can be formed from the inviscid jet by chemically stabilizing its surface (Figure 3, bottom right) and the latter can be formed from the inviscid, argon-levitated, laser-heated melt in a containerless process (Figure 3, center). For details regarding these fibers see Chapter 4.4. Continuous single crystal oxide fibers can be grown from their inviscid melts byrelatively slow updrawing processes (Figure 3, bottom left) such as laser assisted [13] or flux assisted [14] crystal growth. For details regarding growth ofsingle crystal fibers see Chapter 4.5.

B

Preform Ring Heater

Windup

D

E

Melt

+

I-

NOZZI~bbon I

Inert gas

F

Laser \ I Focus- - \ I

I

Melt - -

Quench I \ wheel: I

.

,

(Enlarged)

,

I

.

\

'Ribbon

N~"'~

Fiber- -

Windup

tJ

Quench wheel

Figure 3. Generic fiber forming processes. Drawing fibers from a bushing (A), downdrawing fibers from a preform (B), updrawing fibers from a supercooled melt (C), rapid solidification of a metal ribbon on a quench wheel (D), extruding fibers from a bushing into a chemically reactive environment (E), and drawing fibers from an acoustically levitated, therefore, containerless melt (F). Redrawn and expanded from F. T. Wallenberger and S. D. Brown, High modulus glass fibers for new transportation and infrastructure composites and for new infrared uses, Composites Science and Technology, 51 , 243-263 (1994).

87

Chapter 4

4.1 .2 Structure ofmelts and fibers Fiber forming melts may appear to be homogeneous, but are often immiscible mixtures of a primary major phase with at least one secondary nano-, rnicro-, or even macrodomain [5J. The resulting fibers may appear to be amorphous (literally featureless) when examined by xray analysis, but sensitive tools, e.g., small angle x-ray scattering [57J or NMR [15J often reveal ordered domains from ambient temperature to 2000 K [15] in solid fibers and melts. A description of immiscible domains [5J in terms of coordination numbers, network formers, network modifiers and intermediates is possible, but remains devoid of physical meaning as long as their physical dimensions cannot be tied tomechanical properties. (a) From melts tofibers

This lack of technical information is regrettable, especially since the structure of a melt not only predetermines the structure ofthe resulting fiber [4] but also its mechanical properties. A conceptual overview is inorder. In a strong melt, any secondary domain, nanodomain or any other form of order such as nano-or microcrystals, will dissolve gradually with increasing temperature and decreasing viscosity. In a fragile melt any secondary domain seems to dissolve gradually below the liquidus, but then more rapidly above the liquidus temperature, thereby causing the discontinuous change in melt viscosity with temperature above the liquidus temperature that isobserved. In an inviscid melt, a highly crystalline material is instantly transformed into a highly fluid melt with a very low viscosity, t.e., generally c=:tc=oC:=:=>C=»c:::=>c==>C=> C:::=>C::>C=>C::::=>C=>C:::=:>c::::::>c::::::>C:::=> c::::::>c:::>c=>~c=>c::::::>c::>c=:oc::::::>c::::::>

c::>c::>c::::::>c:::>c::>c=>c::>c=> A c=:o c=> c::::::> c:::=> c::::::> c=> c::::> c=o c=:> c::::::> C=>C::>C:::=>C=>C:::=:>C::::::>c:=>c=>

c:::>c:::=>c::::::>c=>c::>c:=>c:::=>c:=>c:=> C::::::>C:=:::OC::>C:::=>C=::>C::>C:=:=>c=>C:::> c=»c:::=>c=>c:=>e::::::-c:::=>c:=>c:=>

Melt flow

Figure 4. Schematic relationship between the melt uniformity and the structure of resulting fibers. (A) Represents a melt with a uniform second phase flowing horizontally in a fumace toward the bushing, and a glass fiber formed from this melt. Itsrelatively uniform structure can beenvisioned to have high strength. (B) Is a melt with a non-uniform second phase, and a resulting fiber which can bereadily envisioned to possess low strength.

88

Chapter 4

Alternatively, an inviscid melt can be rapidly quenched or solidified to yield an amorphous glass. Analogously, an inviscid jet, the molten precursor ofa fiber, can be rapidly supercooled and stabilized, under certain conditions, atafiber forming viscosity oflog 2.5 tolog 3.0 poise. In a stationary strong, fragile or inviscid melt, the nanodomain(s) are most likely spherical, and have dimensions ranging from relatively small to relatively large. In a dynamic melt, as for example that in a commercial fiber forming process, the nanodomains will most likely directionally elongate as the melt flows toward a bushing (Figure 4). When fibers are pulled from the tiny orifices orbushing tips, the lId ratio of the second phase domains will most likely further attenuate and, upon solidification, form a directional, second solid phase. A composite fiber structure that results may even appear as if it were homogeneous and amorphous (featureless) by older instrumental analysis including x-ray diffraction. Newer experimental methods, such as low angle x-ray scattering, will yield quantitative and confirmative answers In any event, a uniform second phase in a typically inhomogeneous oxide glass melt is expected to yield a more uniform, Le., stronger fiber than a non-uniform second phase, and the lId ratioofthe secondary domain may additionally also affect the fiber modulus. Specifically, silica melts and fibers have a uniform anisotropic network structure. The uniformity of this network structure causes a linear increase of the log melt viscosity with decreasing melt temperature. The anisotropy of the network structure, however, affords a relatively low fiber modulus. Two examples, admittedly extreme cases, may conceptually document the effect of disrupting the uniform anisotropic silica network structure by the addition ofother oxides. For example: (1) the uniformly anisotropic silica structure can disrupted by the addition of alumina and magnesia, i.e., by the formation of a ternary eutectic, Si02-AbOrMgO. Melt viscosity, liquidus temperature and forming temperature drop, the crystallization potential at the liquidus and fiber modulus increase, and a high strength, high modulus HS-glass fiber solidifies. (2) The relatively uniform structure of a ternary high strength SiOrAb03-MgO fiber can be further disrupted by the addition of CaO and either 5-20% boron oxide or 16-17% lithium oxide In both cases, the formation of a secondary, liquid or nanocrystalline, nanophase, is most likely "the initial stage ofstructural ordering" [57). This ordering process stops when the fiber solidifies. No pertinent information is available for melts ofhigh strength (HS) glass fibers, but the formation ofa submlcro-heteroqeneous [57) melt structure has been documented by small angle x-ray scattering for a specific borosilicate composition. It increases with decreasing temperature or increasing viscosity. In summary, the structure of a melt predetermines the structure of a given fiber as well as of itsproperties. Amorphous glass fibers, nanocrystalline glass-ceramic fibers, polycrystalline ceramic fibers, and single crystal fibers possess different levels ofstructural order and uniformity, therefore vastly different mechanical properties. (b) Fiber structure versus modulus

The modulus or stiffness of a fiber reflects its structural or internal order [3]. Figure 5 compares the fiber modulus of silica-alumina and calcia-alumina based composition ranging from 100% silica (or calcia) to 100% alumina. The fiber modulus increases with increasing structural order, Le., from 41 to 125 GPa for amorphous glass fibers to 125 to 250 GPa for nanocrystalline glass fibers, 250 to 400 GPa for polycrystalline ceramic fibers, and 405 to 410

89

Chapter 4

GPa for single crystal alumina fibers. The overall trend reflects the weight percent increase in alumina. Differences in each category reflect super-imposed effects of other oxides (not shown) orprocess-induced differences. Squares represent fibers drawn from strong silica and silicate melts having 0-25% alumina. Triangles represent fibers drawn from fragile aluminate melts having 30-45% alumina, as well as a fragile yttria-modified silicate melt. Inverted triangles represent aluminate fibers spun from inviscid melts having 54-81% alumina. Circles represent sol-gel orslurry spun fibers with 64-99% alumina and single crystal sapphire fibers (100% alumina) made by slow crystal growing processes. Oxide fibers 500-,.---------------------, Single crystal fibers

Saphikon

()

400-+-----------------------1

Fiber-FP ()

Polycrystalline fibers

fu

CIl

en

A/max

300-

Safimax

()

0

:J

"5 'C

o

E

.9l 'w c: ~

o Nexte/480

DSialon-2 Nanocrystalline fibers

200-

Nexte/440

0

OSialon-2

m

100

[]AO

\l 0

Sialon-1

61MS-81

1\ RIMS-54

Y-10

_ .JZ-OO

OA/tex

6Z-32 60-46

OS-Glass DE-Glass

6IMS-54

.6,IMS-80

Amorphous fibers

0 - + - - - -I - , - - - - - , - - - - -I , - - - - - ,I - - - - - 1 o 20 60 80 100 Alumina, wt. %

Figure 5. Fiber composition versus modulus. This illustration correlates the increase in alumina content in the silica-alumina and/or calcia-alumina system with the measured increase in fiber modulus, and observed form of crystallinity. Redrawn and enlarged from F. T. Wallenberger, The structure ofglasses, Sdence, 267, 1549 (1995).

90

Chapter 4

E-glass and S-glass are generic names for two commercial glass fibers. Astroquartz (AO), Nextel, Altex, Safimax, Fiber FP and Saphikon are tradenames for commercial fibers with compositions ranging from 100% silica to 100% alumina. Among the experimental fibers, Z refers to zinc oxide modified silicate and aluminate glass fibers. Z-OO has no alumina, Z-32 has 32% alumina. Sialon refers tooxynitride or nitride modified silicate glass fibers, 0-46 toa quaternary non-silica aluminate glass fiber with 46% alumina, Y-10 toa silicate glass fiber with 10% yttria. IMS refers to inviscid melt spun aluminate glass fibers with 54, 80 and 81% alumina, respectively; and RIMS refers toa redrawn IMS fiber with 54% alumina. AQ, a silica glass fiber, has the lowest modulus (69 GPa) among the silicate fibers derived from strong melts (Figure 5) because it has a highly uniform but anisotropic network structure. Incorporating 10% alumina into itsnetwork structure yields E-glass (aborosilicate fiber) with a modulus of 72 GPa, and incorporating 15% alumina yields S-glass (a magnesium aluminum silicate fiber) with a modulus of 84 GPa. While deferring an analysis of nitride (Sialon), zinc oxide (Z) modified compositions, and inviscid melt spun (IMS and related) glass fibers, the modulus of the other fibers shown in Figure 5 continue to increase with increasing alumina content. This can be seen from 0.46, a quaternary aluminate glass fiber with 46% alumina (110 GPa) which is based on a fragile melt, to those of ceramic fibers ranging from Nextel (140 GPa) toSaphikon (410 GPa) which can no longer be melt spun. Generically, alumina is a powerful modulus modifier. Specifically, alumina increases the crystallization potential of a given melt and therefore the internal order and the modulus, and that of the resulting fibers. Glass fibers with very high moduli become glass ceramic fibers. In ceramic fibers, alumina continues to increase the modulus by the same mechanism and the morphology changes from polycrystalline to single crystal. Superimposed on the effect of alumina are the contributions of other modifiers as shown in Figure 5, i.e., that of nitride (Sialon), yllria (Y), and zinc oxide (Z), and that of inviscid melt spinning (IMS) and related processes. Yttria is a more powerful modulus modifier than ZnO, whereby 10% yttria, as in Y-10, raises the modulus ofa typical S-glass composition from 85 to 130 GPa. All Sialon fibers (Figure 5) were melt spun. Those having moduli ranging from 125 to 140 GPa are known to have an amorphous, glassy structure. Sialon fibers with moduli ranging from 140 to 170 GPa are known to have a nanocrystalline, glass ceramic structure. The dramatic increase in modulus (from 125 to 248 GPa) at about the same silica and alumina level is known to be due to a proportionate increase in surface tension rather than crystallinity. Sialon fibers with the moduli ranging from 170 to 248 GPa are therefore most likely glass ceramic, and not polycrystalline ceramic fibers Inviscid melt spinning (IMS) from aluminate melts having viscosities of 81% alumina could no longer be processed by inviscid melt spinning [17]. Specifically, IMS-54, a fiber with 54% alumina had the lowest modulus (44 GPa). Redrawing IMS-54 at a high temperature [171 yielded a nanocrystalline glass fiber (RIMS-54) with twice the modulus. In summary, the crystallization potential ofalumina can be suppressed by rapid solidification and recovered by heat treatment

91

Chapter 4

(c) Fiber structure versus strength

Strength is a measure of the structural (internal) uniformity and surface uniformity. Table II shows the average tensile strength of sixteen glass fibers. Their strengths range from 5.57 GPa for a military optical glass fiber (FOG-M) to0.37 GPa for a highly porous, high silica fiber obtained byleaching E-glass with hydrochloric acid. All fibers have diameters ranging from 4 to20 urn (19), except FOG-M (62) and the binary calcium aluminate fibers [17], both of which have a diameter of>1 00 urn. Surface flaws or non-uniformities tend to reduce the strength of individual fibers, and directionality or spin orientation, as inferred from birefringence measurements, tends to increase the strength of individual fibers. No study, however, is known that relates spin orientation and/or surface uniformity of a wide range of fiber types, such as those shown in Table II, totheir relative fiber strengths. Surface uniformity and spin orientation are important factors but, even if their relative relationships were fully documented, the up to 15-fold differences instrength among the sixteen generic fibers shown inTable II cannot be attributed to these factors alone.

Table II. Yam tensile strength versus composition Generic glass fiber Military fiber optics glass fiber, FOG-M Magnesium aluminosilicate, 10% MgO Magnesium aluminosilicate, 15% MgO Ultrapure silica fiber, Astroquartz Zn/Ti magnesium aluminosilicate Sodium calcium aluminosilicate E-type aluminum borosilicate Copper aluminum borosilicate Borate glass fiber Lead silicate glass fiber Phosphate glass fiber Sodium silicate glass fiber Calcium aluminate, 54% CaO Pure silica fiber from waterglass Calcium aluminate, 80% CaO Porous silica fiber from E-glass

GPa 5.57 4.80 4.00 3.50 3.20 2.75 2.70 2.70 1.90 1.55 1.50 1.10

0.95 0.85 0.50 0.37

Reference [62] [19] [19] [19] [19] [19] [19] [19] [19] [19) [19] [19] [11)

[4) [11)

[4]

A major reason for the observed differences in fiber-to-fiber strength must be therefore be sought in differences in the uniformity of the internal structures. A uniform network structure such as that of FOG-M silica fibers translates into high tensile strength. A highly nondirectional (random) arrangement of calcium oxide and aluminum oxide in a rapidly solidified binary calcium aluminate fiber translates into very low tensile strength. The ultralow strength ofporous high silica glass isdue to its porosity. Furthermore, the differences in the strengths of amorphous glass. nanocrystalline glass ceramic and polycrystalline ceramic fibers are attributable to differences in the internal order of the fiber structure. For example, the presence of a uniformly distributed minor nanocrystalline (second) phase in a major amorphous (primary) phase of a glass-ceramic

92

Chapler4

fiber may not affect its strength while the weak interface bonding between the crystals in a polycrystalline ceramic fiber is certainly responsible for its very low strength. 4.2 Forming glass fibers fromstrong melts

Two generic fiber-forming processes are generally used to fabricate glass fibers from strong melts; downdrawing from solid preforms and conventional melt spinning, a process that consists ofdrawing fibers from bushings. 4.2.1 Downdrawing from solid preforms Small diameter, structural silica glass fibers and large diameter, bicomponent optical silica fibers are downdrawn from the surface melt of a solid preform. The melt temperatures needed to contain the melt exceed the capability ofpractical ceramic and bushing materials. (a) Structural silica fibers

For structural applications, pure silica fibers are pulled from high purity silica rods as shown in Figure 3, top right. Gas flame or electrical furnaces are used to soften and melt the ends of the preform rods sufficiently for drawing. Each individual silica filament obtained in this downdrawing process has a diameter of 9 urn, and yarns and/or rovings with up to 4800 filaments are then made available to the trade, where they are used either as woven reinforcement fabrics, oras yarns, rovings, and threads for various specialty uses. Pure silica fibers are amorphous glass fibers despite the fact that they are often, and quite incorrectly, called "quartz fibers". They are suitable for high temperature applications. Silica glass fibers have a moderately high tensile strength (3.4 GPa), low modulus (69 GPa) and very high continuous service temperature (1050°C), softening point (1670°C) and volatilization temperature (>2000°C). Chemical analysis shows them to be 99.95+ % pure. The technology, properties, and applications ofsilica glass fibers are analyzed inChapter 6. (b) Optical silica fibers

For optical applications, bicomponent fibers are pulled from the ends of >2 meter long and >0.6 meter diameter preform rods having a silica sheath (clad) and a waveguide core. The preform is fabricated by depositing the waveguide substrate by chemical vapor deposition within a tube of a high silica glass oron a removable rod known as bait. The windup of each fiber is individually controlled by a double laser interferometer and therefore by its diameter, not by a constant windup speed. Individual fibers are covered with protective organic coatings to preserve their inherent strength and to prevent abrasion that might occur during handling. Fiber optics technology and applications ofoptical glass fibers are discussed inChapter 7. 4.2.2 Melt spinning from strong silicate melts Glass-forming silicate melts approach the ideal requirements of a strong melt such as that of silica, but they do so at much lower temperatures. As a result they can be formed by a conventional melt spinning process. Thisdiscussion includes only continuous glass fibers and excludes discontinuous glass fibers such as glass wool made by a centrifuge process.

Chapter 4

93

Structural silicate glass fibers have high melt and forming viscosities (1000 or log 3.0poise) at temperatures ranging from >1100 to 500 compositions could be fiberized on a conventional melt fiberizing unit [17J. Thisprocess is not suitablefor fragile melts. 4

3

LT-200

-----.().-.. Selectedternary and quaternary calcia-alumina melts

2 (J)

til

...

Q.

*

--

LT __-0••

... ...

Cl

0 ....J

0 LT-200 -1

-2 30

LT LT+100

40

50

60

70

80

90

100

Alumina concentration, wt.% Figure 7. Calcia-alumina-silica system. This illustration compares the melt viscosity of the inviscid binary aluminate melts from 60%CaO to 75% alumina. Addition of MgO and CaO produces quatemary aluminate melts, raises theviscositysignificantly andproduces a fragile melt with high viscositiesbelow the liquidus temperature and low viscosities above theliquidus temperature. Redrawn from F. T. Wallenberger and S. D. Brown, High modulus glass fibersfor newtransportation and infrastructurecompositesand fornew infrared uses, Composites Science and Technology. 51 , 243-263 (1994).

100

Chapter 4

Monofilament

Hepurge Water-cooled

_\~OO_C_OI __---, "

1

G

Graphite guide spool

eQ G e c e e

Furnace no. 1

Hepurge Me"ed glass charge in platinum boat

o e e Q e e

Alumina support ring ~ Zirconia support ring

1/2 RPM

Figure 8. Continuous updrawing process. Redrawn from T. F. Schroeder, H.W. Carpenter and S. C. Camiglia, High modulus glasses based onceramic oxides, Technical Report R-8079, Contract N00019-69-C-Q150, US Navy Department, Naval Air Systems Command, Washington, DC, December 1969.

4.3.4 Hybrid fiber forming processes The most important glass forming systems contain elements from the sixth, or chalcogenide, column of the periodic table which includes oxygen, sulfur, selenium, and tellurium. Oxygen containing (i.e., oxide) glasses are insulators; the others tend to be semiconductors. Some melts are inviscid and/or fragile such as the tellurite melts described in the previous chapter. Others are viscous, and whether they represent strong or fragile melts, they are difficult to fiberize. A hybrid process for forming chalcogenide glass fibers [65] has been described that uses elements of downdrawing from preforms and fiberizing through bushings. Specifically, a cylindrical chalcogenide preform isvertically inserted into acylindrical crucible furnished with a nozzle inits bottom plate. The crucible is heated only in the vicinity of the nozzle, and a fiber is continuously drawn from the nozzle at a forming temperature that corresponds to a melt viscosity oflog 3 poise.

101

Chapter 4

Heavy metal fluoride fibers require a fiber forming process that relies on a supercooled melt, while certain fluorophosphate melts can be fiberized by pulling fibers from the melt using a conventional bushing process [66]. Fluorophosphate glass fibers are difficult to pull from the melt because the temperature range around the fiber forming viscosity oflog 3 poise isnarrow and the crystallization rate of the melt is high. Nonetheless, glass fibers were obtained with a diameter of 27 IJm and a tensile strength of 334 MPa for producing reinforced visible-IRtransparent poly (chloro-trifluoroethylene) composites. 4.3.5 Quaternary calcium aluminate fibers Quaternary calcium aluminate glass fibers made by updrawing from a supercooled fragile melt offer superior mechanical properties and sapphire-like infrared transmission spectra. (a) Fiber properties The highest reported pristine strength (8.3 GPa) was obtained (Table VII) with low silica 44.3% Ab03'48.7% CaO - 3.5% MgO - 3.5% Si02 glass fibers when they were updrawn from supercooled melts in an induction furnace [37]. When updrawn in an oxyacetylene furnace, they had a somewhat lower strength (4.2GPa), but it was still higher than that of E-glass (3.5 GPa). The modulus (110.3 GPa) of these fibers was 1.5x that of E-glass. The lowest modulus of any calcium aluminate fiber shown in Table VII was 10% higher than that of Eglass. The stiffest fibers shown in Table VII, a group of zinc oxide modified calcium aluminate glass fibers, had moduli ofup to 122.7GPa, or1.7x the stiffness ofE-glass. With two exceptions, all fibers shown in Table VII had strong hydroxyl bands in their infrared transmission spectra. The exceptions are two recent hydroxyl-free compositions made bythe Davy process [39]. One is a low silica composition and the other is a non-silica composition (46.2% Ab03- 36.0% CaO - 4.0% MgO - 13.8% BaO). An in-depth analysis of the physical properties offibers shown in Table VII isavailable [36-37]. Table VII. Modulus of updrawn calcium aluminate fibers [4, 6-8,31) Modulus GPa 72.0 79.3 103.4 108.9 109.6 109.6 110.3 110.3 110.3 111.7 115.8 122.7 248.0

5.2 5.3 4.0

Na,o,K,O 13.8% Baa 30.0%ZnO

3.5

5.6

25.7% PbO 25.0%ZnO 20.0%ZnO

Process Route and Comments Melt spun through bushing [21J Updrawn low-silica fiber [17] Melt spun through bushing [17] Updrawn low-silica fiber (17) Updrawn non-silica fiber [17] Updrawn high Zno fiber (17) Updrawn low-silica fiber (17) Updrawn low silica fiber (17) Updrawn high PbO fiber (17) Updrawn high Zno fiber (17) Updrawn high ZnO fiber [17] U drawn ZnO/L' 0 fiber 17

102

Chapter 4

100 90

-(-;--_. . I.

80 ;,!! 0

c:

0

'iii en

70 60

'E

50

c:::

40

en

III

t=

., ·· ·,·· ··,

, ..

" " "'. ' ..'

·· ···.'

30 20

"

10 0 0

~

··, ., .. . .,

345

2

6

7

Wavelength, 11m _ . - Ca aluminate

-

-

Sapphire

•••••• Quartz

Figure 9. Spectral transmission of calcia-alumina glass fibers. Redrawn from F. T. Wallenberger, N. E. Weston and S. A.Dunn, Melt spun calcia-alumina fibers: infrared transmission, J. Non-Cryst. Solids, 12[1),116-119 (1990).

Their resistance to alkaline media exceeds that of commercially available AR silicate glass fibers (Chapter 6) having a zirconia content of up to 15% [20). Hydroxyl-free quaternary calcium aluminate glass fibers (Figure 9), e.g., non-silca fibers containing 46.2% Ab03 36.0% CaO - 4.0% MgO - 13.8% BaO, afford sapphire-like infrared transmission properties. (b) Potential applications

The commercial potential of updrawn quaternary calcium aluminate glass fibers was tested in two stages. In the early 1960s, they were evaluated because they yielded higher moduli than those which could then be achieved with silicate glass fibers [36-37]. Timing for this development coincided with the onset ofthe commercial development ofcarbon fibers and no new aluminate orsilicate glass fiber was commercialized until 1995. In the early 1990s, a renewed evaluation of quaternary calcium aluminate glass fibers was triggered by their sapphire-like optical properties [17] [42-43]. Properly coated aluminate fibers might afford sapphire-like sensor performance applications at a more affordable cost than single crystal sapphire fibers, but unlike the latter, they would be limited to ambient and moderately high temperatures [431. Calcium aluminate fibers are not commercially available, but so far they offer valuable models for important structure-property relationships [4] regarding strength, modulus, and optical properties.

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4.4 Amorphous fibers from inviscid liquids Melts of metals as well as crystalline ceramic oxides have low viscosities. All solidify at a sharp melting point and their viscosity above the melting point increases rapidly and then reaches a viscosity comparable to that of motor oil at room temperature, i.e., log 25 urn in thickness. Solder and brazing ribbons have solidus temperatures ranging from 770 to 1130°C and liquidus temperatures ranging from 925 to 1150°C. Applications include aerospace components, electric motors, and brake pads [60].

4.4.3 lnviscid melt spinning (IMS) processes In processes with conventional quench rates of -104 K1s, the inviscid liquid (melt) or resulting jet is carefully under- or supercooled to its narrow fiber forming range (log 2.5 to log 3.0 poise). Three types of fibers have been made by variants of this process: continuous optical yttrium aluminum garnet (YAG) glass fibers, continuous aluminate glass fibers and steel fibers.

106

Chapter 4

(a) Principles of jet and fiber formation

The key principle governing the formation and breakup of a liquid jet is well known [12]. A liquid jet is unstable with respect to viscosity, diameter, and surface tension. A jet has a tendency to break up due to axisymmetric surface pressures, and produces Rayleigh waves, orperiodicvariations ofincreasing amplitude in the jetdiameter (Figure 12). Ultimately, these diameter variations cause the jet to break up into separate droplets, and they, in turn, crystallize and form shot. For any process to yield a continuous fiber from a low viscosity liquid, the transient viscosity ofthe melt exiting the bushing tip orspinneret orifice must quickly reach the fiber forming level oflog 2.5-log 3.0poise before jetinstability (formation of Rayleigh waves) and potentially disruptive crystal growth can occur.

Figure 12. Straight fiber and frozen Rayleigh waves. The straight, cylindrical fiber represents an inviscid melt spun calcium aluminate glass fiber that had been surface stabilized with particulate carbon. The frozen Rayleigh wave structure represents a calcium aluminate fiber that was not surface stabilized and solidified while it was in the process ofbreaking upinto droplets and shot.

For fiber formation (Equation 1), a jet must have a lifetime (t) sufficient for the viscosity to reach log 2.5-log 3.0poise before the onset of turbulence orbreakup [12]. The viscosity (11) of the jet is the key factor since it depends exponentially on the temperature. The diameter (D) ofthe jetisthe next most important factor. Thus, a reduction of the viscosity (11) shortens the jet lifetime (t) and limits the attainment of low diameter fibers. Surface tension (y) and density (p) are less important, i.e., less sensitive totemperature. t

= J4 [( pD 3 / r)' 12 + [31]D / r ))

dn / d t: TL < T"

L dVc / dr

(1 ) (2) (3)

Chapter 4

107

The conditions (Equation 2) under which oxide or metal fibers can be formed from inviscid melts is also defined by the dynamics of the relationship inthe forming zone between the rate of solidification or the change of the viscosity (11) and the rate of crystal growth (Vel with time (T). Any change in melt viscosity, even if only of the surface viscosity, of a mol en jet will disproportionately affect its lifetime. If the lifetime is short, the jet will, on cooling, break up into Raleigh waves, liquid droplets and crystalline shot, and if it is long, the jet will instead solidify into a continuous fiber on cooling. (b) Principles ofincreasing the jet lifetime One inviscid melt spinning process, the containerless laser heated melt process (Chapter 4.4.4) is believed tofacilitate the formation of fibers by increasing the viscosity of the inviscid melt (and jet lifetime) at a normal quench rate of 104 Kls, i.e., without increasing the quench rate to -108 Kls. Two older inviscid melt spinning processes require a chemically reactive environment and are believed to facilitate fiber formation only by increasing the surface viscosity of the jet. They are a metal fiber forming process (Chapter 4.4.6) and an oxide glass fiber forming process (Chapter 4.4.7). The theory governing inviscid melt spinning by chemically modifying the jet and fiber surface surface isdiscussed in Chapter 4.4.8. 4.4.4 Oxide fibers from containerless, laser heated melts Containerless liquid phase processing has been successfully used to achieve deep undercooling of molten oxides to a temperature significantly below their liquidus temperatures where the viscosity is sufficiently high for fiber pulling. Containerless conditions eliminate heterogeneous nucleation by containers (crucibles, spinning cells or bushings, and precious metal tips) and help deep undercooling [73). One may assume the fiber forming viscosity, although not reported, was between log 2.5and log 3.0poise.

Optical pyrometer

Fiber _ Stepper motor

---- Levitated molten oxide _ Stinger

Stepper motor controller

Computer

Figure 13. Container1ess laser-heated melt process. Redrawn from J. K. Weber, J. J. Felton, B. Cho and P. C. Nordine, Glassfibres of pure and erbium- or neodymium-doped yttria-alumina compositions, Nature, 393, 769-771 (1998).

108

Chapter 4

In this process (Figure 13), molten samples of yttrium aluminum garnet (YAG: YJAb01Z), 3 mm in diameter, were levitated in a flow of argon gas. The levitated material was completely melted ina continuous wave COz laser heating beam. The viscosity ofthe inviscid melt above the melting point was 0.05 Pa.s (0.5 poise). The laser beam was then blocked, resulting in a cooling rate of -250 K1s. At a pre-selected fiber forming temperature between 1600 and 1660· C, i.e.,a temperature that iswell below the liquidus temperature, fibers were pulled from the levitated droplet by rapidly introducing and withdrawing a 100 IJm diameter tungsten wire "stinger". At higher temperatures the stinger pulled out of the melt without forming a fiber, while crystallization was likely tooccur atlower temperatures [73]. Glass fibers, up to 0.5minlength and 5 to 30 IJm in diameter were pulled from the apparatus at rates of 1.0-1 .5 m/s before crystallization terminated the process. The fibers had a homogeneous appearance and smooth surfaces. They were transparent, highly flexible, and x-ray amorphous. Glass fibers of this kind, which were also made from erbium- or neodynium-doped ytlria alumina compositions, would therefore expand the range of fiber lasers into the mid-infrared. The economics and scalability of the new process are not known. The materials cost and the cost ofoperating a laser process are probably about the same for an amorphous YAG sensor fiber made by the containerless laser heated melt process and a for single crystal YAG sensor fiber made by laser heated pedestal growth (Chapter 4.5.2). And both are containerless processes. However the higher process speed may favor the laser heated melt process (1 .5 m/s) over the laser heated pedestal process (1 mm/s). 4.4.5 Metal fibers ina reactive environment Commercial wire drawing processes produce metal wires with round cross sections but they are highly energy and labor intensive. Wire drawing falls outside the scope of this book. Commercial rapid solidification processes yield amorphous metallic ribbons. lnviscid melt spinning yields metal fibers by a chemically assisted jetstabilization process. In the inviscid melt spinning process [10) [51), steel wires are formed by the same mechanism as glass fibers. In this case, the process shown in Figure 14 depends on the presence of silicon in the steel formulation and on the presence of carbon dioxide in the process environment. A commercial pilot production unit based on the schematic process diagram shown in Figure 14 consisted of a 0.9 m long furnace, a 1.5 m long cooling column and a windup [51). The furnace contained a crucible holding 50 kg of steel, and was heated with a 4 kHz power unit supplying 70 KW. An essential ingredient in the cooling medium was carbon dioxide. Wires with diameters of 100-200 IJm can be made at 1500·C with speeds of 10-20 mIs, respectively. At a rate of 15 mIs, the spinning of a 165 IJm diameter wire lasted 4 hours. In this pilot process, unbroken wire was obtained for periods in excess of 1 hour which represents a continuous length ofabout 80 km. As shown in Equations 4 to 6, carbon dioxide diffuses into the surface of the inviscid molten jet. Its concentration decreases with increasing distance from the surface, but wherever it finds silicon that is evenly distributed throughout the melt (and therefore the surface of the liquid jet), it forms silicon dioxide and carbon which cause a steep increase in the surface

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viscosity. Carbon may further react to form carbides. The surface skin is neither a sheath nor a film. (4)

Si+20~Si02

2 Si + 2 CO ~~ SiO

2

+ 2C

(5) (6)

Si+C~SiC

ESCA analysis confirms the presence of oxidized silicon and oxidized iron in the wire surface or skin [51). The silica peak gradually disappears at a depth of 100 nm, giving way to the peaks ofiron and silicon. Analysiswith a CAMECA ion analyzer shows that the intensity of Si peaks decreases to naught between 17 and 55 nm from the surface of a 165 IJm diameter steel wire. The results parallel those noted for carbon with aluminate fibers. Silica is the most viscous inorganic material known, especially at 1500°C. It is the ideal surface viscosity builder to increase the lifetime of a hot inviscid steel jet long enough to prevent formation of Rayleigh waves and shot. Other viscosity builders can be formed in-situ at the spinning temperature to stabilize a given molten metal jet [10). They must have a higher melting point than the metal, and be insoluble in the molten jet[10).

-

_

Reactive gas

Figure 14. Inviscid melt spinning process (schematic drawing). Redrawn from F. T. Wallenberger, N. E. Weston and S. A. Dunn, Inviscid melt spinning: as-spun amorphous alumina fibers, Materials Leiters, 2[4]121-127 (1990).

Inviscid melt spinning is considered to be a potentially viable alternative to wire drawing [51) for making steel wires for radial automobile tires, but a prior product development did not reach beyond the pilot plant level. Using silica steels, the complex chemistry (Equations 4-6) produce also minute amounts of iron oxides which were detected by ESCA [51), and are a potentially undesirable trace byproduct. The challenge [4) remains to fine tune the chemistry ofthis process before commercial development.

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4.4.6 Oxide glass fibers ina reactive environment Continuous binary calcium aluminate glass fibers can also be formed by inviscid melt spinning. In this case, carbon particles which are formed by the decomposition of propane enter into the surface of the molten jetand raise its surface viscosity, a process that lengthens the lifetime of the jetand prevents itsbreakup. In this process [8) [10), alumina and calcia are placed in a tungsten crucible having an orifice in itsbottom plate. The crucible isplaced ina furnace and the surrounding air is replaced with argon. The oxide powder is melted and the mell is maintained 100°C above its melting point under a mild vacuum (Figure 3, bottom right, and Figure 14). The fibers are spun from the oxide mell after introducing propane below the crucible and by increasing the argon pressure above the crucible. Propane decomposes on the hot surface of the mollen jet, forms carbon particles which enter into the surface of the jet, and sometimes deposits an additional secondary carbon sheath on the fiber surface after it solidifies. These fibers contained 51 .5- 80.2% alumina and 43.5 to 19.8% calcia, and occasionally 3.8 to4.0% silica and/or 0.1 to 7.5% magnesia (Table VIII). Pristine strengths ranged from 0.16 to 1.05 GPa (vs. 3.5GPa for E-glass) and moduli ranged from 41 .1 to61.2 GPa (vs. 72.5 GPa for E-glass). The moduli [11] were much lower than those of updrawn high viscosity calcium aluminate fibers, suggesting much lower internal order. By SEM, the fracture patterns were typical ofglass fibers, and fibers with up to 80% (!) alumina were amorphous.

Table Vlll, Properties of inviscid melt spun calcium aluminate glass fibers (11) Al,O, 54.0 54.6 54.8 59.0 60.8 66.8 67.3 80.2

Fiber Composition (Wt % ) CaO MgO 39.0 3.0 39.0 2.5 38.9 2.5 40.8 0.2 39.1 0.1 33.2 34.2 19.8

Pure silica glass fiber Borosilicate E-glass

SiO, 4.0 3.9 3.8

100.0 54.0

Spin T. °C 1700 1500 1500 1500 1700 1800 1700 1800

Diam . J.1IIl 170 190 216 190 225 102 167 117

Strength GPa 0.5 0.7 0.6 0.7 0.5 0.9 0.8 1.1

Modulus GPa 50.8-68.9

1750 1200

100 10

3.5 3.5

69.0 72.0

46.6-61,2 45.9-56.4 41.0-54.3

Some fibers have a secondary carbon sheath which may be up to 600 nm thick [17). Not surprisingly, these fibers are black. The secondary overgrowth is not an integral part of the fiber [11-12]. It does not affect the fiber properties, and can be peeled or burned off. Other fibers have no carbon sheath. These fibers are translucent. Carbon that is present in the fiber surface or skin does not affect its transparency, whether the fiber had originally no carbon sheath as-spun, orwhether the sheath had been removed from the as-spun fiber [12). Sputtered neutral mass spectrometry (SNMS) depth profiles document that carbon is present in the skin of all fibers to a depth of about 50 nm (Figure 15), whether a given fiber has a secondary carbon sheath overgrowth or not [11). X-ray photoelectron spectroscopy (XPS)

Chapter 4

111

showed that carbon in the surface or skin [12) consists of carbide (51%) carbon (41%) and carbonate (8%). Infrared depth profiling by diffuse reflectance infrared spectroscopy (DRIFT) provided further important insights [52). It showed that carbon alters the oxygen environment of the aluminum atoms near the fiber surface from octahedral to tetrahedral coordination and promotes the generation of carbonaceous species such as ethers and esters in addition to carbonates and carbides [52) which have also been found with XPS [12). In summary, a typical aluminate fiber is spun between 1500 and 1700·C (Table VIII). Carbon enters into the skin of the still liquid jet. A secondary carbon sheath is obtained when more carbon is present in the reactive propane environment than needed to stabilize the liquid jet, but after the fiber is solidified «500·C). At higher temperatures it would oxidize (burn off). Thus, only some fibers have a secondary carbon sheath. If it is formed, it results from side growth, a secondary growth mode, that has already been discussed with regard to the growth ofcarbon whiskers in acarbon vapor environment (Chapter 2.1 .1). 100.------:::::=========j 90

80

70

s!:!. 5060 ~

:; LL

oe

40 30

20 10 50

100

150

200

Depth, nm

Figure 15. SNMS depth profile ofa translucent calcium aluminate fiber. Redrawn from F. T. Wallenberger and S. D. Brown, Highmodulus glass fibers fornew transportation and infrastructure composites and fornew infrared uses. Composites Science and Technology, 51 , 243-263 (1994).

Inviscid melt spun calcium aluminate glass fibers have low strength (0.5-1.1 GPa) and moduli (46-58 GPa). Low strength and low stiffness can be attributed to the random structure frozen into the fibers during rapid solidification. As a result, they are not likely to become composite reinforcing fibers, despite their excellent alkali resistance which they share with quaternary calcium-aluminate fibers [9). 4.4.7 Mechanism ofjetsolidification Continuous aluminate glass fibers are formed in the presence of propane, but not in its absence. A viable mechanism of jet stabilization must therefore explain (1) the function of carbon which enters into the surface orskin of the molten jet, (2) the function of carbides and carbonates which are instantly formed in the molten jet surface, and (3) the increase in tetrahedral from octahedral coordination of aluminum atoms in the surface [51) before the fiber solidifies and secondary overgrowth with carbon can occur.

112

Chapter 4

The principle governing jet formation and jet breakup [12] has already been discussed in Chapter 4.2.2. A liquid jet is unstable and will degrade into Rayleigh waves (Figure 12, right) and then droplets. The stability of a jet (Equation 1)depends on liquid density, jetdiameter, melt viscosity, and surface tension. The lifetime of a jet is the time required for the melt to traverse the continuous length of the jet before the onset of Rayleigh waves (12). A jet of a silicate glass oran organic polymer melt has a high viscosity (>W Pa.s) and a lifetime greater than 10° seconds; it can be spun or drawn from the melt by conventional methods, and it solidifies well before itcan form Rayleigh waves. Table IX. Properties of inviscid calcium aluminate jets [12] Alum. content 51.5% 54.6% 66.8% 80.2%

Fiber structure amorph amorph amorph amorph

Melt temp.

Spin temp.

1415 1390 1650 1830

1500 1500 1700 1900

0c.

0c.

Jet/Fiber Diam . urn 375 190 105 118

Melt density

g/cm 2}0 2.70 2.68 2.68

Surface tension mN/m 680 680 625 575

Melt viscosity Pa.s 0.34 0.55 0.14 0.06

Unassisted jet life, sec 1.4x10" 8.7x10·' 2.0xlO'3 1.7xlO'3

A jet of an aluminate melt with >50% alumina has a low viscosity «1 Pa.s) and a calculated lifetime less than 10.2 seconds (Table IX). If ejected into an ambient, neutral environment it will form Rayleigh waves and droplets (or shot when they freeze) rather than uniform continuous fibers. Equation 1 shows that viscosity is the major factor in determining jet lifetime; surface tension is a secondary factor [12). Any increase in the surface viscosity of the molten oxide will disproportionately increase the jet lifetime from that calculated for an unassisted jet (2.0x10·3 seconds) to that calculated for an assisted jet (2.0x10·1 seconds) which must have been obtained since continuous fibers were obtained. Particles, especially shaped particles are known to increase the viscosity of a suspension, following Mooney-Einstein [53). Thus, carbon particles (12) may enter into the surface or skin of a liquid inviscid jet and increase its viscosity sufficiently and long enough to facilitate its solidification and fiber formation. For this mechanism to be viable, three conditions must be fulfilled. (1) The increase in the jet surface viscosity must afford a stabilized (assisted) jet lifetime that at least matches the jetcooling time. (2) The assisted lifetime resulting from the viscosity increase must be comparable to the actual (unassisted) lifetime of a typical silicate fiber such as E-glass. (3) The surface viscosity increase needed to achieve this lifetime must be realistically achievable by carbon insertion inthe jetsurface. The pyrolytic production of carbon has been said to create a "snowstorm of large, flat molecules containing the hexagonal ring structure of graphite" [54), Le., flakes or flat aggregates of smaller particles. They enter into the surface of the molten jet and act as viscosity builders, where they and their instant reaction products such as carbides can be detected by ESCA. Since jetgeometry and surface forces tend toconstrain particle formation into planar structures parallel to the jet surface, the rheological treatment for flakes is appropriate. The viscosity (11) of a Newtonian fluid containing solid, suspended particles, relative tothat (110) of the suspending fluid, isgoverned by Equation 7 [53].

In (71 1710) = kEf:z l [1- f2 1 f ilii

(7)

The Einstein coefficient (kE) for incorporated particles depends on particle shape, f2 is the volume fraction offiller, and fm isthe maximum packing fraction for the flakes. The viscosity of

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113

the suspending fluid (bulk molten) oxide is 0.14 Pa-s, and the required 36.6 Pa-s viscosity for the surface layer of the jetcan be attained with flake shapes having UDbetween 4 and 9, and volume fractions of solids between 0.32 and 0.45. The combined volume fraction of the solids in the oxide skin detected by SNMS was calculated (12) to be 0.508, a value well above that needed forrheological jetstabilization. E-glass jets with a diameter comparable to that of the 66.8% alumina jet can bemelt spun at temperatures ranging from 1100 to 1480°C(11) where they have a viscosity between 6.1 and Table X. Chemical stabilization of inviscid calcium aluminate jets Jet/Fiber Compo (21) (105 }lIlt Diameter) CaO-A},O, (66.8%) E-glass

Melt Behavior Inviscid Strong

Temp. °C 1700 1360

Viscosity, Pa.s Bulk Surface 0.14 36.6 31.7 31.7

Jet Lifetime, sec Unassisted Assisted 1 2.0x10·' 2.6xlO· 2.6x10"

816 Pa-s and a calculated lifetime between 8.8x10·2 and 1.2x10' seconds (Table X). The assisted lifetime of the 66.8% alumina jet (2.6X1 0 1 seconds) is therefore well within the range of the unassisted E-glass jet lifetimes (12). In summary, the stabilization of inviscid aluminate jets can be attributed to the increase in surface viscosity due to suspension of solid carbon particles (as well ascarbides and carbonates) in the molten oxide surface (12). The viscosity (rheology) controlled jetstabilization appears to be accompanied bythe observed change [52) from octrahedral totetrahedral coordination ofaluminum atoms near the surface. 4.4.8 Cryogenic fibers from liquefied gasses

A viable process forthe formation ofcontinuous, self-supporting fibers such ashydrogen from

liquefied gases has emerged over the past two decades (74). Like all prior process iterations (74), it appears to be an inviscid melt spinning process (IMS) and not a rapid solidification (RS) process. The first step in this process consists of increasing the viscosity, presumably to log 2.5 to log 3.0 poise, bydecreasing the temperature, e.g., forliquid hydrogen, to 29 K. The second step consists of extruding a liquid jet into a closed system and tosolidify the resulting fiber at a 10 K lower chamber temperature, l.e., at 19 K.

4.5 Growing single crystal fibers from inviscid melts Continuous single crystal fibers can be grown from inviscid melts by two relatively slow processes: the edge defined film fed growth (EFG) process (13) and the laser heated float zone (LHFZ) or laser heated pedestal growth (LHPG) process (14). Both offer growth rates of to0.3-0.7 mm/s [13-14). 4.5.1

Edge defined film fed growth

Edge defined film fed growth (EFG) is a commercial process (13) that facilitates the fabrication of continuous void free single crystal oxide fibers (Figure 16) from tungsten or other growth orifices.

114

Chapter 4

(a) Growth of sapphire fibers

This process yields commercial single crystal sapphire fibers. A liquid pool from which the continually growing filamentary crystal iswithdrawn isformed on top ofa planar surface of the orifice and fed by capillaries which extend down through the orifice into a liquid reservoir. The crystal shaping or edge definition is maintained by the geometry of the top surface of the orifice and the fulfillment of a contact angle of 4 ~m/s [67]. In summary, the laser heated float zone (LHFZ) method [67], inparticular the traveling solvent zone melting (TSZM) configuration [68], is a highly effective technique to grow centimeter long crystals of high Te and other low dimensional cuprates. High temperature superconducting fibers, wires, tapes, and ribbons have also been made by the powder-in-tube method. These are sheath/core bicomponent fibers or ribbons having a protective metal sheath and a functional core consisting of an appropriate multiple oxide material. In this method [32] the super conducting powder iscontinuously introduced into a metal tube and the filled tube isdrawn by conventional wire drawing methods. Although the powder-in-tube method did yield the highest current density[64] reported so far (>1,000,000 Alcm 2, 77 K, 0 T), it is not reviewed here in detail because it is a metal drawing, not a melt forming process. High temperature superconducting sheath/core bicomponent fibers have also been made by introducing the superconducting material into the core of hollow glass fibers as they are formed under the bushing (Chapter 6.32). 4.5.3 The future ofsingle crystal oxide fibers Continuous sapphire fibers (Chapter 4) and continuous sheath/core bicomponent silicon carbide/carbon fibers (Chapter 3) offer impressive performance as reinforcing fibers and in ceramic and metal matrix composites. Here are some noteworthy commonalties and differences. (a) Single crystal sapphire fibers

Sapphire fibers are hard, strong and scratch resistant to most materials and provide excellent wear surfaces. They can withstand higher pressures than polycrystalline alumina since they lack the grain boundary interface breakdown of the latter. Sapphire fibers transmit ultraviolet, visible, infrared and microwaves and serve as excellent wave guides between 10.6 and 17 microns, and offer durable and reliable IR transmission. By virtue of their high thermal conductivity they can be rapidly heated and cooled. EFG sapphire fibers melt sharply at 2050°C and maintain measurable strength at extreme temperatures [13]. Table XII shows tensilestrength as a function of test temperature from 25 to 1500°C. The room temperature strength, 3.57 GPa, is low for a single crystal fiber but typical for sapphire fibers, irrespective ofprocess. For a single crystal fiber, room temperature strength should be approaching the theoretical value, which is>10 GPa. In fact, its strength is only 40% of an about equal diameter polycrystalline sheath/core bicomponent silicon

Chapter 4

119

carbide/carbon fiber (Chapter 3). This deficiency ofsingle crystal sapphire fibers still needs to be corrected. The tensile strength ofsapphire fibers at 1500·C in this example (Table XII) is 0.55 GPa. In addition, isolated literature values report strength levels of0.40 GPa up to1900·C. These are impressive results since they refer to an oxidative environment. Strength levels of0.80 GPa Table XII. Strength of EFG sapphire fibers at elevated temperatures Fiber test temperature (0C) 25 400

800

1094 1500

Average tensile strength (GPa)

3.57

2.08 1.85 1.03

0.55

Standard deviation (GPa) 0.66 0.49

0.34 0.14

0.09

were observed for silicon carbide/carbon fibers at 1600°C, but the thermal stability of silicon carbide, except for single crystals, is lower than that of single crystal oxides in an oxidative environment especially inprolonged use above 1400°C. (b) Other single crystal oxide fibers

Continuous single crystal oxide fibers, including sapphire, have a number of property advantages over comparable polycrystalline oxide fibers (see Chapter 8). They include microstructural stability at high temperatures, retention of high elastic moduli at high temperatures, and creep resistance. But because of the high diameters, single crystal oxide fibers made by today's processes cannot be woven and must either be wound or used as inserts. Further improvements of the high temperature creep behavior are therefore being sought. The goal is an optimum continuous single crystal oxide, irrespective of process [13] [14] [4849], or new continuous single crystal silicon carbide fibers by laser assisted chemical vapor deposition (Chapter 3) having low diameters «15IJm), near theoretical strength at room temperature, and low creep, high strength and high strength retention at 1600 to 2000°C in oxidative environments. Key opportunities also exist for single crystal fibers for high T, superconductor [67] and for optoelectronic applications [71]. REFERENCES [1) [2J [3] [4] [5] [6]

[7]

D.R. Uhlmann, A kinetic treatment ofglass formation, J Non-crysl. Solids, 7,337-348 (1971). A. Angell, Relaxation in liquids, polymers, and plastic crystals - stronglfragile patterns and problems, J. NonCrystalline Solids, 131-133, 13-31 (1991). F.T.Wallenberger, The structure ofglasses, Science, 267,1549 (1995). F. T. Wallenberger, Melt viscosity and modulus ofbulk glasses and fibers - challenges forthe next decade, in "Present state and future prospects of glass science and technology", Kreidl Symposium, Triesenberg, Liechtenstein, July 3-8, 1994, Glasstech. Ber. Glass Sci.Technology 70C, 63-78 (1997). A. K. Vareshneya, Fundamentals of inorganicglasses, Acad. Press, Boston (1994). H. Tokiwa, Y. Mimura, T. Nakai and O. Shinbori, Fabrication of long single-mode and multi-mode fluoride glass fibers bythe double crucible technique, Electronics Letters, 21 [24], 1130-1131 (1985). M. L. Nice, Apparatus and process forfiberizing fluoride glasses using a double crucible and the compositions produced thereby. US Patent 4,897,100, Jan. 20, 1990.

120 [8J [9J [10) [11] [12] [13] [14] [15] [16] [17J [18) [19] [20] [21] [22] [23) [24) [25) [26] [27] (28) [29J (30) [31] [32] [33] (34) [35) [36] [37] [38]

Chapter 4 G. Y. Onoda, Jr. and S. D. Brown, Low silica glasses based on calcia-aluminas, Journal of the American Ceramic Society, 53(6), 311-316 (1970). F. T. Wallenberger, N. E. Weston and S. D. Brown, Calcia-alumina glass fibers: drawing from super-cooled melts versus inviscidmelt spinning, Mat. Letters, 11 (89), 229-235 (1991). R. E.Cunningham, L. F. Rakestraw and S. A. Dunn, Inviscid mett spinning of filaments, in Spinning wire from motten metal, J. Mottern and W. J. Privott, ed.; AIChE Symposium Series, 74(180), 20-32 (1978). F. T. Wallenberger, N. E. Weston and S.A. Dunn, Inviscid melt spinning: as-spun amorphous alumina fibers, Materials Letters, 2 [4]121-127 (1990). F.T.Wallenberger, N.E.Weston, K. Motzfeldt, and D. G.Swartzfager, Inviscid melt spinning ofalumina fibers: chemical jetstabilization, Journal oftheAmerican Ceramic Society, 75[3], 629-639 (1992). R. S, Feigelson, Growth of fiber crystals, in Crystal growth of electronic materials, E. Kaddis, Ed.,127-145, ElsevierScience Publishers, London (1985). J. Monbleau, Single crystal technology, Product Bulletin, Saphikon Inc., Milford, NH (1994). B. Cote, D. Massiot, F. Tantelle and J. P, Coutures, 27AI NMR spectroscopy of aluminosilicate melts and glasses, Chemical Geology, 96, 367-370 (1992). A. L. Greer, Metallic glasses, Science, 267,1947(1995). F.T.Wallenberger and S.D. Brown, High modulus glass fibers fornew transportation and infrastructure composites and fornew infrared uses, Composites Science and Technology, 51, 243-263 (1994). F. T. Wallenberger, New melt spun glass and glass-ceramic fibers forpolymer and metal matrix composites, in High performance composites: commonalty of phenomena, K. K. Chawla, P. K. Law and S. G. Fishman, eds., The Minerals, Metals and Materials Soc., 85-92 (1994). V. E. Khazanov, Yu, I Kolesov and N. N. Trofimov, Glass fibers, in Fibre science and technology, V. I. Kostikov, Editor, 15-230, Chapman and Hall, London (1995). K. L. Loewenstein, The manufacturing technology of continuous glass fibres, third, completely revised edition, Elsevier, Amsterdam (1993). P. K. Gupta, Glass fibers forcomposite materials, Chapter 2 in Fibre reinforcements forcomposites materials, A.R. Bunsell, ed., Composite Materials Series 2,Elsevier, Amsterdam, 19-71 (1988), R. A. Houpt, R. M. Potter, T. D. Green, D. P. Aschenbeck and C. Berdan, II, Dual-glass fibers and insulation products therefrom, US Patent 5,431,992, July 11, 1995. J. E. Loftus, C. R. Strauss and R. Houston, Method for making fibers by causing one glass to flow around another asthey arespun, US Patent 5,539,596, June 25,1996. W. L. Eastes, D. A. Hofmann, J. W. Wingert, Boron-free glass fibers, International Patent Application, W096/39362, December 12, 1996. J. F. Sproull, Fiberglass composition, US Patent 4,542,106, September 17, 1985. H. Kaplan-Diedrich and G. H. Frischat, Properties of some oxynitride glass fibers, Journal of Non-Crystalline Solids, 184,133-136 (1993). J. Kobayashi, M. Oota, K. Kada and H. Minakuchi, Oxynitride glass and the fiber thereof, US Patent 4,957,883. September 18, 1990. K. Suganuma, H. Minakuchi, K. Kada, H. Osafune and H. Fujii, Properties and micro-structure of continuous oxynitride glass fiber and itsapplication toaluminum matrix composite, J. Mater. Res., 8 [1], 178-186 (1993). F. T. Wallenberger, S. D. Brown and G.Y. Onoda, ZnO-modified high modulus glass fibers, Journal of NonCrystalline Solids, 152,279-283 (1993). P. F.Aubourg and W. W. Wolf, Glass Fibers, in Advances in Ceramics, Vol. 18, Commercial Glasses, pages 51-63, D. C. Boyd and J.F. MacDowell, editors, American CeramicSociety, Westerville OH (1986). F. Rossi and G. Williams, A new erainglass fiber composites, Paper presented at the 28 th AVK Conference, Baden-Baden, Germany, pages 1-10, October 1-2,1997. U. Balchandran, A. N. Iyer, P. Haldar and L. R. Motowidlo, The powder-in-tube processing and properties of Bi-223, J. Metals, 45[9], 54-67 (1993) F. M. Costa, R. F. Silva andJ. M. Vieira, Influence onepitaxial growth ofsuperconducting properties of LFZ BiSr-Ca-Cu-O fibres, Part I.,Physica C, 289, 161 -170 (1997) and Part II., Physica C, 289, 171-176 (1997). T. D. Erickson and W. W. Wolf, Glass composition, fibers, and methods for making same, US Patent 4,026,715, May 31 ,1977 F. T. Wallenberger, N. E. Weston and S. D. Brown, Infrared optical tellurite glass fibers, Journal of NonCrystalline Solids, 144 (1),107-110 (1992). S. D. Brown and G. Y. Onoda, Jr., High modulus glasses based onceramic OXides, Report R-6692, Contract NOw·65-Q426-d, US Department ofthe Navy, October 1966. G. Y. Onoda, Jr. and S. D. Brown, High modulus glasses based onceramic oxides, Report R-7363, Contract N00019-67-C·301, US Department ofthe Navy, February 1968. T. F. Schroeder, H. W. Carpenter and S. C. Carniglia, High modulus glasses based on ceramic oxides, Technical Report R-8079, Contract N00019·69·C-Q150, US Navy Dept., Naval Air Systems Command, Washington, DC, December 1969.

Chapter 4

121

[39] J. R. Davy. Development of calcia-alumina glasses for use in the infrared spectrum, U. S. Patent No. 3,338.694 (1967), Glass Technology. 19 [2], 32-36 (1978). (40) R. Maddison, Calcia-aluminas, Product Bulletins WB37A and WB39B, Sassoon Advanced Materials LTD, Dumbarton. U.K. (1994). [41] P. R. Foy, T Stockert, J. Bonja. G. H. Sigel, jr., R. McCauley, E. Snitzer and G. Merberg, Meeting Abstracts, American CeramicSociety, 94th Annual Meeting, Presentation 7-JXV-92, Minneapolis, MN, April 12-16, 1992. [42] F. T. Wallenberger, N. E. Weston and S. A. Dunn, Melt spun calcia-alumina fibers: infrared transmission. J. Non-Cryst. Solids. 12[1),116-119 (1990). [43] F. T. Wallenberger, N. E. Weston and S. D. Brown, Melt processed calcia-alumina fibers: optical and structural properties. inGrowth of materials for infrared detectors, R. E. Longshore & J. Baars. Eds., Proceedings ofthe SPIE, Society ofPhoto-Optical Instrumentation Engineers, Bellington, WA, Volume 1484, 116-124 (1991). [44] F. T. Wallenberger, Design factors affecting the fabrication of fiber reinforced infrastructure composites, Annual Wilson Forum. Santa Ana, CA, March 20-21, 1995; in Applications of Composite Materials in the Infrastructure, 1-10 (1995). [45] F. T. Wallenberger, High modulus glass fiber reinforced composites for currently emerging infrastructure applications. Proceedings, ASCE Materials Engineering Conference, San Diego, California, November 14-16, 1994. [46] M. Matsukura, Z. Chen, M. Adachi and A. Kawabata, Growth of potassium lithium niobate single-crystal fibers bythe laser-heated pedestal growth method, Jpn. J. Appl. Phys., 36. Part 1, No. 9B. 5947-5949 (1997). [47] J. T. A. Pollock, Filamentary sapphire - The growth of void-free sapphire filament at rates upto 3.0 em/min, Journal ofMaterials Science, 7, 786-792 (1972). [48] T. Mah, T. A. Parthasarathy, M. D. Petry and L. E. Matson, Processing, micro-structure, and properties of AI203-YJAlsO,2 (YAG) eutectic fibers, Ceramic Engineering and Science Proceedings, 622-638, 17th Ann. Conference onComposites and Advanced CeramicMaterials, Am. Ceram. Soc.,Westerville OH (1993). [49] K. J. McClellan, H. Sayir, A. H. Heuer. A. Sayir, J, S. Haggerty and J. Sigalovsky, High strength, creep resistant Y203-stabilized cubic Zr02single-crystal fibers, Ceramic Engineering and Science Proceedings, 651 659, 17th Ann. Conf. on Composites and Advanced Ceramic Materials, Am. Ceram. Soc., Westerville, OH (1993). [50] R. S. Feigelson, D. Gazit, D. K. Fork and T. H. Geballe, Super-conducting Si-Ca-Sr-Cu-O fibers grown bythe laser-heated pedestal growth method, Science, 240,1642-1645 (1988). [51] J. M. Massoubre and B. F. Pflieger, Small diameter wire making through solidification of silicon steel jet, in Spinning wire frommolten metal, J. Mottern and W. J. Privott, eds.:AIChE Symposium Series, 74 (180), 48-57 (1978). [52] F. Fodeur and B. S. Mitchell, Infrared studies of calcia-alumina fibers, J.Am. Ceram. Soc., 79 [9] 2469-2473 (1996). [53] M. Mooney, The viscosity of a concentrated suspension of spherical particles, J. Colloid Science, 6 (2), 162170 (1951). [54] R. J. Diefendorf and E. R. Stover, Pyrolytic graphites: how structure affects properties, Metal Progress. 81 [5J. 103-108 (1962). [55J S. A. Dunn and E. G. Paquette, Redrawn inviscid melt-spun fibers, Advanced Ceramic Materials 2. 804 (1987). [56] B. S. Mitchell. K. Y. Yon. S. A. Dunn and J. A. Koutsky, Phase identification incalcia-alumina fibers crystallized from amorphous precursors, Journal ofNon-crystalline Solids, 152, 143-149 (1949). [57J V. V. Golubkov, A. P. Titov and E. A. Porai-Koshits, On the structure of lithium borate glasses according to small angle scattering data. Soviet Journal ofGlass Physics and Chemistry, 18 [2], 122-129 (1992). [58] D. E. Polk and B. C. Giessen, Amorphous orglassy materials, inRapidsolidification technology source book, edited byR. L.Ashbrook, 213-247, American Society forMetals, Metals Park, Ohio (1983). [59] H. H. Liebermann, Rapidly solidified al/oys. Marcel Decker, Inc.,New York (1993). [60] H. H. Liebermann, MetglaS® product bulletin, Allied Signal, Parsippany, NJ (1993). [611 D. M. C. Narashima, Planar flow casting ofalloys, US Patent 4,142,571 (1979). [62J J. E. Ritter and J. D. Helfinstine, Atougher fiber for the FOG-M, Photonics Spectra, 8. 90-93 (1967). [63] H. Miao, J. C. Dietz, L.A. Angurel, J. I. Pena and G. F. dela Fuente, Phase formation and micro-structure of laser floating-zone grown BSCCO fibers: reactivityaspects, SolidState lonics, 101-103, 1025-1032 (1997). [64J G. Geiger, New record for super-conducting wire, The American CeramicSociety BUlletin, 74(12), 19 (1995). [65] J. Nishii, I. Inagawa, T. Yamagishi. S. Morimoto and R. lizuka, Process for producing chalcogenide glass fiber, US Patent, 4,908,053, March 13, 1990. [66) H. Lin. W. L.Dechent, D. E. Day and J. O. Stoffer, Preparation and properties of mid-infrared glass fibers and poly(chlorotrifluoroethylene) composites, J. Mater. Sci., 32, 6573-6578 (1997). [67] A. Revcoleschi and J. Jegoudez, Growth oflarge high-Tcsingle crystals by the floating-zone method: a review. Progress inMaterials Science, Vol. 42, 321-339 (1997). [68J P. H. Keck and M.J. E. Golay, Phys. Rev.,39, 1297 (1953).

122

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[69] W. Que and S. Lim, Evaluation ofmicro-structure characteristics oflithium niobate single-crystal fiber with magnesium-ion-indiffused cladding, J. Am. Ceram. Soc., 80, (11) 294548 (1997). (70] J. I. Peiia, H. Miao, R. I. Merino, G. F. de la Fuente and V. M. Orera, Polymer matrix synthesis of zirconia eutectics for directional solidification into single-crystal fibers, Solid State lonics, 101-103, 143-147 (1997). (71] W. M. Yen, Preparation of single-crystal fibers, in Insulating maferials foropto-electronics, F. AguIl6-Lopez, Editor, World Scientific, Singapore (1995). (72] W. Jia, H.Yuan, L. Lu, H. Liu and W. M. Yen, Phosphorescent dynamics in SrAI204:Eu 2',Dy3+ single-crystal fibers, J. Luminescence, 76 &77, 424428 (1998). (73] J. K. Weber, J. J. Felton, B. Cho and P. C. Nordine, Glass fibres of pure and erbium- or neodymium-

30

-. ~ r-. -,

v-R-glass

r-. G r-, r-,

c: ~

in .9! '0; 20 c:

~

E-glass--"

~

10

1\ ,1\ ,

\ \

o

100

200

300

400

500

,

600

700

Temperature Co

Figure 5. Strength and modulus of R-glass, a typical HS-glass fiber, atelevated temperatures, Drawn from data contained in French patent 1,435,0739, issued in 1963 to Saint Gobain Company, Chambery, France,

In another tensiletest, R-glass is exposed for prolonged periods of time to a specific elevated temperature before its strength is tested. At 750· itstensile strength drops within the first 400 hours to 1 GPa (or1/4th its room temperature strength), but then remains nearly constant for at least the next 600 hours of exposure. A temperature of 750·e is well beyond the reach of either ofthe two general purpose E-glass fibers. In summary, high strength fibers have superior high temperature resistance, superior impact damage resistance, high strength retention at elevated temperatures, and intermediate moduli. The formation of a nanocrystalline structure which facilitates the attainment of high service temperatures tends toreduce strength, but in most cases the strength loss is minimal.

(b) Properties and applications HS-glass fibers fill an important niche in the composite reinforcement market between E-glass and the lowest cost carbon fibers. They are used in a great variety of applications in the aerospace and aircraft industry, automotive industry, electrical and electronics industry, sporting goods industry, and in military markets. The higher cost of other inorganic high

140

Chapter 6

strength fibers can be reduced by hybridization, Le., by designing composites reinforced with mixed HS-glass/carbon, HS-glass/aramid, or HS-glass/boron fibers. In the aerospace and aircraft market HS-glass fibers are used because of their high specific properties and relatively low cost, e.g.; in satellite components, motor cases, nose cones, aircraft flooring, cargo liners, and radome skin sheets. Their use in helicopter rotor blades benefits from the high damage resistance (toughness) of HS-glass fibers. In automotive composites, HS-glass fibers are used because of their high specific strength in pressure vessels and their superior toughness in leaf springs. Their use in compressed natural gas cylinders isone ofthe fastest growing applications. In the electrical and electronics market, HS-glass fibers are used as strength members for optical fiber cables, printed circuit boards, and other cables, but mostly when high temperatures are involved. In the sporting goods market, they are used for product differentiation and because of their toughness in tennis racquets, squash racquets, fishing rods, surfboards, windsurfing masts, skis, archery bows, arrow shafts, and racing yachts. In military markets, HS-glass fibers are used in aircraft fuel tanks because of their specific strength, and inrigid armor and helmets because oftheir superior ballistic impact resistance. 6.2.2 High modulus - high temperature fibers HM-glass fibers have modulus ranging from 100 to 132 GPa, and otherwise offer about the same properties as HS-glass fibers, such as high strength and superior resistance to high temperature. Some experimental HM-glass or glass ceramic fibers are known to have not only higher stiffness but also higher strength than typical HS-glass fibers. Their superior stiffness and strength, however, is often accompanied by a significant increase in their density, afact that requires special consideration oftheir specific strength and stiffness. In addition to alumina (see Chapter 4), beryllia [2) [16], ZnO [17], titania [4), lanthana [16), yttria [101 and CuO [10] (16) are known as modulus modifiers for silicate glass compositions (16). Lanthana, for example, can raise the fiber modulus of a typical HS-glass composition from 84-95 to100 GPa (Table V) , but because ofitshigher density, the specific modulus (37.7 Mm) remained unchanged. Zinc oxide, another effective modulus modifier, can raise the modulus to 104 GPa, and the specific modulus to 41.1 Mm. BeO yielded the first glass fiber (YM-31) that became officially known as a generic HM-fiber. It had a complex BeO-modified composition [2], but a simpler, statistically designed BeO-modified HM-glass fiber (Table V) had [15-16] a modulus of 112 GPa and a specific modulus of43 Mm. The effectiveness of BeO as a modulus builder is the result of the high field strength of the Be-2 ion and itsability tocoordinate four oxygen ions tightly toit [17]. The structure of BeO is ofthe wurtzite type. The only other wurtzite structure with oxygen is ZnO. In addition, ZnSi04 is isomorphous with Be2Si04• On a molar basis, the ZnO-modified and the original BeOmodified YM-31A compositions are identical [17], and on this basis ZnO is a less effective modulus builder than BeO (Table V) . The modulus of Y203-modified fibers (Table V) increases to 132 G [10]. At the same time, however, their density increases to 4.0 g/cc, their specific modulus drops to 30 Mm, their crystallization potential increases and their strength drops to 3.5 GPa [10]. Highly Y203modified HM-fibers possess a clearly observable nanocrystalline phase [10]. Upon heat treatment, large microcrystals «2 IJm) grow and significantly reduce the strength of these

141

Chapter 6

fibers. The value-in-use or cost considered performance of highly Y203-modified HM glass fibers is therefore lower than that oftypical HS-glass fibers (Table IV). Table V. High Modulus silicate glass fibers Modulus Modifier Composition, wt. % SiO, A~O,

CaO MgO La,O,

Zno

La,0,

ZnO

BeO

Y,O,

Y,O,

BeO/Y,o,

50.0 32.5

45.8

50.0 35.0

52.6 35.8

41.7 27.2

36.2 20.5

7.5

5.4

4.3

8.1

6.2

26.8

5.0 30.2

4.5 100 3.0 1.36 33.0

3.5 130 4.3 0.81 30.0

5.4 132 3.29 1.64 40.0

12.5 5.0

22.4

BeO

Y,O,

z-o, Li,o c-o,

Fe,O,

Structure Morphology References:

7.5

1.8 6.8 1.7 2.5 0.3

TiO,

Pristine Properties Strength, GPa Modulus, GPa Density, glee Sp. Strength, MIn Sp. Modulus, MIn

11.0 7.7

4.8 100 2.69 1.78 37.2

104 2.77 37.5

5.1 112 2.60 1.96 43.0

amorph. glass

amorph. glass

amorph. glass

(16) (18]

(17)

(15)

amorph. microcryst. glass glass cer. [10]

(10)

amorph. Glass [15] [18]

Y20:JBeO-modified HM-glass fibers (Table V) offers a better balance of mechanical properties than either Y203- or BeO-modified fibers. Their specific strength (1 .64 Mm) approaches that of the Y203-modified fiber (1.94 Mm) and their specific modulus (40 Mm) approaches that of the BeO-modified fiber (43 Mm). These fibers have a 30% higher specific stiffness than Sglass and would facilitate a 15% weight reduction of aerospace composites assuming the use of a 50% fiber volume fraction . Y20:JBeO-modified fibers are commercially unattractive because ofthe suspected toxicityof BeO.

6.2.3 Ultrahigh modulus glass-ceramic fibers By modifying the ternary Si02-Ab03-MgO or the quaternary Si02-Ab03-MgO-CaO eutectic compositions with appropriate oxide modifiers one can raise the fiber strength to 5.4 GPa (vs. 3.4 GPa for E-glass) and the fiber modulus to 132 GPa (vs. 72 GPa for E-glass). The increase in modulus iscaused byan increase ininternal order as evidenced bya change from an amorphous to a nanocrystalline structure. An increase in fiber modulus to 248 GPa can be achieved by inserting nitrogen into the oxide network [19-231, thereby creating a nitridemodified silicate, Si-AI-O-N, oroxynitride glass fibers. This increase is caused byan increase in surface tension as evidenced bymicrohardness. Thus, one mechanism seems to depend on increasing structural order, the other on crosslinking.

142

Chapter 6

(aJ Process and products Oxynitride fibers are formed by an adaptation of the conventional bushing process. Since nitrides would corrode precious metals in an oxidative environment, the bushings with up to 200 tips are therefore made from boron nitride-coated carbon or from molybdenum. The melts are formed under nitrogen at 1600-1750 ·C and refined at lower temperature. Fibers with diameters ranging from 12to 20 urn are continuously drawn from the melt -200·Cbelow the melt temperature, and mechanically wound at 1000-2000 m/min [24].

Silicon formation :

Si0 2 +Si1N"

~2Si+2SiO+2

N2

(1 ) (2)

Oxygen formation :

Si1N" · Si0 2 +12

Mo~4

Mo1Si+2 N 2 +0 2

(3) (4)

Silicon oxidation:

Si+0 2 ~Si02

(5)

The melt process may cause the reduction of silica yielding particulate silicon. Increasing numbers of silicon defects, when formed, impart a blue-gray (hazy) to dark-brown (opaque) appearance to the glass (Equations 1 and 2), and these defects may proportionately reduce the strength of the resulting fibers [23]. The formation of silicon defects can however be reversed byusing molybdenum (but not boron nitride) asthe bushing material (Equations 3 to 5), and by inserting an 8 hour refining cycle midway between the melt and fiber forming temperatures (as shown for the second entry inTable VI). A refined Si-AI-O-N glass is colorless and clear [23] and the resulting defect-free fibers [20] [24] are much stronger (>4.0 GPa) than those obtained from unrefined melts (2.0-3.0 GPa). Si-AI-O-N compositions with 15% nitrogen yield glass ceramic fibers with moduli ranging from 140 to 248 GPa, and microhardnesses ranging from 900 to 1200 kg'mm'2 [19]. The substitution of nitride for25% of a given oxide composition under the same conditions of synthesis no longer yields glasses. The respective melts will foam and crystallize [25]. Selected glass and/or glass ceramic Si-AI-O-N fibers are shown inTable VI. The first [23] is Sglass, and the second is a Mg-Si-AI-O-N composition that is similar tothe S-glass composition

143

Chapter 6

but modified to contain 7.5 wt.% SbN. [23). Since the latter had been properly melt refined before it was fiberized [24), it yielded high strength fibers. lis low nitrogen content produced a significant modulus increase beyond that of the S-glass control. The third example in Table VI, a Ca-Si-AI-O-N, has a still higher nitrogen content [19) and therefore a higher absolute and specificmodulus. lis low strength suggests that the melt was not or could not be adequately refined .

.-

26.0

a. '"

(!)

vi ::3 :; "0 0

E 0

~ til

24.0

I-

22.0

l-

20.0

-

18.0

l-

16.0

I-

14.0

I-

12.0

I-

10.0

I-

8.0

l-

.,

I

iii

6.0

,~

,...

_V'

.,..",,'

.....

,f.

• / A

i

I

I

20

1150

-

1100

-

1050

-

1000

'"vi

-

950

l:

-

900

-

10

1200

')I

E

~

l/)

~

700

-

600

-

500

-

400

Ql

"E

'"

s:

eQl '"

5

30

Nitrogencontent. weight % Figure 6. Modulus and microhardness of Sialon or oxynitride fibers versus nitrogen content. Redrawn from J. Kobayashi, M. Oota, K. Kada and H. Minakuchi, Oxynitride Glass and the Fiber Thereof, US Patent 4,957,883, September 18, 1990.

The remaining examples in Table VI are ultrahigh modulus (UHM) glass ceramic fibers. The commercial Ca-Mg-Si-AI-O-N development fiber [20) reflects proper melt refining and has a moderately high nitrogen content. Accordingly, it has a measured strength of 4.0 GPa and a measured modulus of 180 GPa, as well as a specific strength of 1.38 Mm and a specific modulus of 62.5 Mm [20]. The second example, a Y-Si-AI-O-N fiber, adds the predictable effect of yttria which increases measured modulus and density, but minimizes the specific modulus. The third example, a Ca-Mg-Si-AI-O-N fiber has the highest measured modulus (248.0GPa) and the highest specific modulus (75 Mm) ofany known oxide glass fiber.

144

Chapler6

In summary, nitrides offer only a modest crystallization potential but instead seem to act as modulus builders by crosslinking the oxide structure they modify. And the presence of nitrogen in the network structure, once introduced, may restrict the dimensions of crystals which can be formed . A selected oxynitride melt having a high nitrogen content may yield nanocrystalline fibers with an ultrahigh modulus, while a nitrogen-free oxide melt with the same cation ratio may already yield microcrystalline fibers with lower moduli. Thus, the modulus of oxynitride fibers can be increased to much higher levels than that of fibers from oxide melts. Crystallinity ofoxynitride glasses can be correlated with their nitrogen content by infrared methods [25]. HM-oxynitride glass fibers are x-ray amorphous. UHM-glass ceramic fibers [20] are nanocrystalline (as shown in Chapter 4). Table VI. Ultra-high modulus Si-Al-0-N glass fibers Fiber Modifiers Composition, wt. % SiO, Alp, CaO MgO Y,O, Si.,N. A~N,

N-content

Melting temp., °C Refining temp., °C Fiber forming T, °C Bushing tips Pristine properties Strength, GPa Modulus, GPa

Density, glee Sp . Strength, Mm Sp . Modulus, Mm

Structure Morphology References:

5-glass Control

Commercial Si-AI-0-N glass fibers yCaMgCa-Mg-

65.50 25.00

56.75 25.33

9.50

ID.42

0.00

7.51

32.38 50.78 16.84

Unspecified commercial 48SiO,' 43CaO· 5MgO· 4Al,o, based fiber 12.00

50.3 7.6

Si-Al-0-N Ca-Mg 11.1 2.2 62.6 0.7

25.1 23.4 17.0 4.37

13.5

0.00

2.54

10.00

1650

1750

1600

1790

1565 200

1720 1600 1500 200

1380

1560 1

1590

4.60 88-95 2.53 1.85 35-37

4.50 115.00 2.80 1.61 41.10

213.00 3.94

248.0 (3.3)

54.00

(75)

amorph. glass [23]

amorph. glass [23] [24]

137.0 2.8 48.9 amorph. glass (19]

4.00 180.00 2.89 1.38 62.30 nanocr. glass cer . [20] [23]

nanocr. glass cer. [22]

1

nanocr. glass cer. (19]

(b) Properties and applications Glass and glass ceramic oxynitride fibers can be produced with high strength by properly refining melts while they are formed, and with ultrahigh moduli by inserting nitrogen into a suitable oxide network structure. The highest measured modulus (248 GPa) that has been reported lies between those of standard modulus (8M) and intermediate modulus (1M) carbon fibers (230 and 303 GPa, respectively), and the highest specific modulus (75 Mm) lies midway between those of 8M carbon fibers (131 Mm) and E-glass fibers (27 Mm). Two types of applications are being pursued with oxynitride fibers.

Chapter 6

145

Sialons are reinforcing fibers for metal matrix composites. An aluminum alloy 6601 matrix reinforced with a development fiber (Table VI) had a strength of 4.0 GPa and a modulus of 180 GPa. The bending strength of the MMC was 25% higher than that of an alumina fiber reinforced control, and almost the same as that of a silicon carbide fiber reinforced control. A high value-in-use may result if this fiber were to cost less than the incumbents did. Oxynitride fibers have high alkali durabilityand may be useful as a diaphragm material in the electrolytic production of chlorine from aqueous NaOH [26] and as a reinforcement for cementicious composites [26]. 6.2.4 Fibers with high chemical stability This discussion of chemical stability refers to bare (uncoated) glass fibers, i.e., fibers having neither a specific acid or alkali resistant finish nor a secondary coating. Accordingly, the relationship between fiber composition and chemical stability in water, acids, and bases is complex. It depends on the interaction between (1) the chemical agent [27] to which the glass fiber surface is exposed, (2) the pH of the glass composition [33] in the fiber surface, and (3) the internal microstructure ofthe fiber [27].

(a) Chemical resistance ofglass fibers The first step in the attack of water on the bare surface of an alkali-free or near-alkali-free glass fiber is its adsorption. The adsorbed water molecules hydrolyze the siloxane bond by protonation ofthe oxygen atom, and yield a highly hydroxylated fiber surface. = Si - 0 - Si = ~ = Si - 0 ·· ·· · -Si = ~ =S-OH +HO-S

=H+ +OH -

(6)

With high alkali-glass fiber surfaces, the reaction of water represents an electrophilic attack by the addition of the proton (W) to the negatively charged oxygen atom of the =Si-O-M bond. It proceeds in the same fashion and results, practically speaking, in the ion exchange between W (or H30+) and either alkali ions and/or network modifying alkaline earth ions, and leads to the formation of SiOH groups [27]. The reaction products of water with a highly alkaline fiber surface are NaOH, Ca(OHh and hydrated sodium silicate. = S-OM +H + ----+=S-OH + M+

(7)

However, as soon as the supply of H+ions is exhausted, the corrosive attack ofH20 turns into a nucleophilic attack by alkali ions on the fiber surface. In other words, the reaction of a glass fiber with water turns into a reaction of a glass fiber with alkali. Siloxane bonds are broken and Si-O-Na groups are formed until the glass fiber is completely dissolved in the highly alkaline medium that initially consisted only ofwater. The reaction of alkaline media starts with a nucleophilic attack by hydroxyl ions on silicon atoms in the bare surface of a glass fiber (-OH' + =Si-O-S=) where it forms new bonds (=SiOH, and/or =Si-OM). Monovalent cations (e.g., sodium) are removed from the glass fiber, leaving behind a hydroxylated surface. Bivalent cations (e.g. calcium) remain attached to the glass surface and form a crystalline sheath growing in thickness. Such a sheath develops when bare silicate glass fibers (>50% Si02) including E-glass [27], AR-glass [30], basalt [29] and oxynitride fibers [26] or when bare aluminate glass fibers (>50% Ab03), e.g. calcium aluminate fibers [18] as discussed inChapter 4, are exposed toalkaline media.

146

Chapter 6

The crystalline sheath, mostly consisting of Ca(OH)2[26), increases the alkali resistance, but also limits the practical utility of the resulting fibers since it drastically reduces their strength. Zr02, SnO, La203, Ti02, Fe203 [10), Y203[26) and Na20 enhance the alkali resistance perhaps by delaying sheath formation, but rather large amounts of Na20 (>10%), ZrO (>15%) or Y203 (>30%) and combinations of Na20 (11%) and Zr02 (16%) are often employed. These fibers are not really alkali resistant, only more alkali resistant than E-glass. In the end, all lose their physical integrity and are destroyed as evidenced from the accelerated leach test shown for E-glass inTable VII. Table VII. Composition of borosilicate E-glass and its solutes after leaching (8 hours/95°C) Composition E-glass (%)

SiO, 53.8

AI,O, 14.9

Fe,O, 0.3

CaO 17.1

MgO 4.7

B,O, 8.7

Solute (%), in H,O in2nH,SO. in2nNaOH

0.2 1.7 25.6

0.2 14.6 7.3

0.3 0.2

0.2 16.6 8.4

0.1 4.6 1.3

0.7 8.6 6.3

While alkaline media are known to create a crystalline deposit or sheath on the surface of glass fibers, mineral acids selectively dissolve specific components of the glass, first of all the ions of network modifiers [10]. Silanol bonds, as a rule, are not broken, and Si~ is not dissolved. If however the amount of Si02is not sufficient to create a continuous network structure, cations can selectively dissolve inacid media. The addition ofZr, Ti, and Fe oxides, even to the high alkali oxide glasses shown in Table VII, substantially increases the acid resistance [10J, but in the end, the entire fiber, whether it is compositionally an E-glass or Aglass, isconverted into a porous high silica fiber (see Table VII and Chapter 6.4.4). In summary, the effect on the pH of the bare fiber surface and the effect of the interaction between a chemical agent and a bare fiber surface are predictable. Zr02 seems to increase both acid and base resistance. The effect of the internal microstructure [27J ofa fiber ishighly process dependent and not predictable without a thorough prior investigation of its microstructure. Importantly however, all fibers, except experimental single fibers, have a primary finish; some have an additional secondary coating. These modifications further reduce the predictability of their chemical resistance from their compositional make-up alone.

(b) Alkali resistant glass fibers CemFil, a commercial AR glass fiber [1-2) [30), an experimental AR glass fiber [28J, and an AR glass-ceramic basalt fiber [29J, commercial only in Russia, are shown in Table VIII. CemFil and the experimental AR fiber are highly Na20- and Zr02-modified glass fibers. Basalt fibers are derived from volcanic rock with high Fe 203+FeO levels. All have higher alkali resistance than E-glass. That of basalt fibers lies between E- and AR-glass. All require higher fiber forming temperatures than E-glass. Basalt fibers require an energy intensive, nonstandard process. Experimental AR fibers offer a 100°Clower fiber forming temperature than that ofCemFil and basalt fibers. In a test that simulates their suitability as a cement reinforcement [30J, bare AR-glass and bare E-glass fibers were immersed at25, 50 and 80°Cina solution (NaOH 0.88 gIl, KOH 3.45 gIl, Ca(OH)l 0.48 gIl, pH 12.5) that simulates the aqueous phase of Portland cement. Both

Chapler6

147

fibers lost strength between 24 and 96 hours, i.e., in the time frame during which the alkalinity of cement reaches its peak as it cures. At 25°C, the strength of bare E-glass dropped to 2/3 of its original value in 24 hours, and that of AR-glass in 96 hours. In the accelerated test at BO°C, the strength of E-glass dropped to 1/3 of its original value in 24 hours and that of ARglass in 96 hours (Figure 7). Thus, neither fiber is usable to structurally reinforce cement without having an effective secondary alkali resistant coating. The application of an alkali resistant finish or secondary coating is known to render even Eglass suitable for continued use as a durable reinforcement of cement structures, whether they are composite wraps for bridge columns, or net-like structures aimed at roadbed construction. But since even AR-glass loses strength in 70%) as well as silicate glass fiberscontaining very high combined levels of BaO and Ti02offer significantly higher dielectric constants, i.e. 13.0and 13.5, respectively. Table XI. Glass fibers wit h very high dielectric constants Glass fibers Composition, wt .% SiO, A~O,

B,O, PbO CaO MgO

srO

BaO TiO,

z-o,

Na,0 K,O Nb,O,

DieJ. constant, £r, @ 1 MHz Forming tem p ., TIog2.5, ·C Liquidus temp., T,, ·C ~T (F-L), ·C Fiber forming abilit y Referen ces:

E-glass cont rol

L-glass w/>70% PbO

BaD-TiO, silicate

54.3 14.0 6.6

26.0

40.0

55.0 2.5

47.16

7.5

9.0

7.03

7.5 15.0 23.0 7.0

6.0 15.0 7.8 1.7

7.03 14.05 13.95 3.27

3.0

7.51

1.5 72.0

22.1 0.6 0.5

Nb,0,-BaO-TiO, silicates

0.8 0.2

0.5

6.8

13.0

13.5

10.1

12.3

1299

850

1077

1199

1136

1063 +236 excellent Table I

650 +100 very poor [37J [38J

1214 -137 infeasible [37J [38J

1085 +114 very good [37J [38J

1089 +51 good [37J [38J

Lead (L) glass compositions possess a dielectric constant of 13.0, but glass fibers from lead glass compositions with such high PbO levels are difficult to form. They are very long melts which have a very high !1T (Tlog25 - Ti) between forming and liquidus temperatures and are therefore quite crystallization resistant. However PbO will evaporate violently during the melt process, thus affording highly non-uniform fiber compositions and frequent process discontinuities and fiber breaks [37-38). In addition, the use of PbO poses significant environmental concerns.

152

Chapter 6

The silicate glasses based on high levels of BaO and Ti02 have an equally high dielectric constant of 13.5, but fibers cannot be formed from their melts because the liquidus temperature is much higher than the preferred forming temperature. Crystallization would occur long before the melt reaches the preferred fiber forming viscosity of log 2.5 poise. In summary, glass fibers with very high levels ofPbO are sublimation prone and glass fibers with high combined levels BaO and Ti02are crystallization prone. In contrast, addition of 0.5to 15.0 mol% of Nb205toglass compositions having high levels of BaO and Ti02 slightly reduces the dielectric constant, and dramatically reduces the liquidus temperature, thereby reversing the ~T (TI09 25 - TL) from -137"C to +117"C (or 242°F). The ~T quoted here [37-38] describes the difference between Tl09 2 5 and TL, while the ~T customarily quoted in the technical literature refers to the difference between T109 3 and TL. This difference does however not affect the conclusions. High speed and high frequency information transmission is becoming increasingly important with the recent development of advanced information systems including mobile communication by car telephones and personal radios, as well as satellite broadcasting and cable television. As a result, there is an increasing demand for miniaturizing electronic devices and also microwave circuit elements such as dielectric resonators in conjunction with the electronic devices. Nb205 containing glass fibers are embedded for these applications in a resin such as polyethylene oxide that has a low dielectric tangent (tan 0) loss to obtain the desired high frequency performance of the resulting circuit boards for microwave applications [37-38]. In summary, microwave circuit elements can be made more compact when using a circuit board having a high dielectric constant. It acts toconcentrate the electromagnetic energy within the board and thereby minimizes the leakage ofelectromagnetic waves.

(d) Fibers with super- and semiconducting properties Superconducting glass fibers are obtained by incorporating a suitable ceramic material in the fiber core, yielding superconducting bicomponent sheath/core glass fibers (Chapter 6.3.3). Superconducting bicomponent metal ribbons are obtained by incorporating a suitable ceramic material inthe core ofcontinuous metal tubes. Superconducting single component fibers can be made drawing single superconducting ceramic preform rods bythe laser heated float zone orpedestal growth process (Chapter 4.4.2). Semiconducting glass fibers have been known for over 30 years [52] but have never attracted commercial interest because significant technical problems have never been solved. In principle, such fibers (e.g., CuO-CaO-Ab03-Si02) can be made by adding oxides of monovalent metals such as copper or silver to a suitable base glass and by subsequently reducing the glass fibers in various gaseous media, e.g., hydrogen. However these fibers are moisture sensitive, and prolonged storage leads to increased glass conductance (10). The electrical properties of glass fibers are best modified by (1) applying a permanent chemical coating to the fiber surface, (2) adding a suitable material to the binder or finish formulations, or (3) modifying the composite matrix that is being reinforced with a glass fiber. A semiconducting or conductive coating is applied by vacuum deposition, metallization from metal salts, decomposition of organometallic compounds or chemical metallization. Also, carbon black can be added tothe composite matrix.

Chapter 6

153

(e) Fibers with bone bioactive glass compositions Bioactive glasses have been developed since 1969 [67] and continuous bioactive glass fibers since 1983 [68-69]. Very recently (70), bone bioactive glass fibers were found tobond tobone tissue and help bone tissue growth when they are preferentially placed on the surface of a thermoplastic composite while carbon fibers are used tostiffen the core ofthe structure. A glass composition which was found particularly suitable for these applications (67) contained 52% Si02, 30% Na20, 15% CaO and 90%), the residual strength drops to about 80-85% of the low original strength. Upon exposure to temperatures between 300 and 500°C, surface crystallization occurs and residual tensile strength drops to about 35-50% of the original strength. 8ecause of their thin walls, hollow fibers are more fragile than solid fibers and need to be handled more carefully. They will not however crush under pressure inalaminating press. Hollow E-glass fibers can produce either composite parts with equal thickness and up to 25% lower part weight than E-glass, or composite parts with equal weight and up to 25% higher thickness. In equal weight filament wound cylinders, the hydrostatic collapse pressure is increased by 30% and the dielectric constant is reduced from 6.8 to 4.0. In equal thickness laminates, acoustic transmission is increased as the result of lower mass, dynamic fatigue is significantly increased, and thermal conductivity is reduced by up to 40% . In a hybrid laminate, partial substitution for graphite does not reduce the high specific modulus while removing up to one-half ofthe more expensive graphite fiber. Weight reduction is a powerful incentive in aircraft manufacturing. Every kilogram ofstructure that can be eliminated facilitates a commensurate reduction infuel consumption and therefore cost, ora commensurate increase inpayload atequal fuel cost. Hollow glass fibers qualify for use in non-load bearing interior aircraft applications such as sidewall and ceiling panels. The use of hollow glass fibers for such applications are more prevalent in Russia and CIS countries [53) than inthe United States and the Western world. (c) Hollow porous sheathlcore silicate glass fibers Hollow glass fibers with a porous sheath were made but are not commercially available. They were first formed on laboratory [55) orcommercial equipment [54) and then acid leached by a process otherwise used in the production of porous, solid glass fibers (see Chapter 6.4.4). For example, hollow porous glass fibers with >95% Si02were produced by acid leaching of hollow glass fibers including borosilicate E-glass [54). Acid leaching of hollow glass fibers based compositions having 35 to62% Si02, 1 to 11 %Ab03, 0 to54% 8203, 3 to9% Zr02 and I to 29% Na20 (but no CaO) gave silica-rich, porous hollow glass fibers with good alkali resistance. Acid leaching of hollow glass fibers containing 54.0% Si~, 22.4% CaO, 14.3% Ab03, 7.2% 8203and 1.0% Na20 (but no Zr02) gave porous, hollow glass fibers with low alkali resistance [54). One of the most challenging tasks is to produce porous, hollow glass fibers with controlled pore size as required for inorganic membranes having uniform mechanical properties including strength and stiffness, thermal and chemical stability, photochemical and biochemical durability, and superior resistance to compaction under high membrane pressures. The required control over the desired pore size in fibers has recently been demonstrated [55) with a glass composition primarily consisting of 57.2 mol%, 22.8% 8203, 9.2% CaO. This study serves as a model for the evaluation of hollow porous glass fibers in reverse osmosis, phase separations, salt extraction, and biochemical research. (d)

Hollow superconducting sheath/core glass fibers

High temperature superconductor materials are superconductive when cooled below their respective critical temperatures (Te). This temperature should be higher than 77°K so that

ChaplerG

159

they can yield functional performance in liquid nitrogen. Low temperature, metallic superconductor alloys perform only in liquid helium, and represent a far more costly process. Known high r, materials include Y-Ba-Cu-O (YBCO) and Bi-Sr-Ca-Cu-O (BISCO). High Tc superconducting fibers have been made by three processes: The first is the powder-in-tube or Taylor process that relies on encapsulating superconducting powders in a metal tube. The second is the laser heated float zone orlaser heated pedestal growth process and the third is the fiberglass process except that a different bushing isrequired. So far the most practical route is the powder-in-tube process that consists of continuously filling metal tubes with high Tc ceramic powders, reducing their diameter :n a metal drawing process, and pressing the resulting bicomponent sheath/core metal/ceramic wire into a ribbon shaped product. Wire drawing falls outside the scope of this book. The laser heated float zone or laser heated pedestal growth is a slow crystal growth process. It is useful for experimental exploration but not commercial production. Forming sheath/core fibers having a glass (rather than a metal) sheath surrounding the high Tc superconducting ceramic material goes back to1990 [44]. In the earliest process for making superconductive glass fibers [44·45], the raw material powder needed for Y-Ba-Cu-O type fibers was filled into a large, thick walled, closed end glass tube. The tube was first heated until the superconductor powder melted and was further heated until a solid concentric bicomponent sheath/core fiber could be downdrawn. One component, the sheath, was glass, and the second component, the core, was a superconducting material. Essentially, solid fibers were downdrawn from a preform having a molten core. Upon solidification, the fiber was proven to possess moderately effective superconducting properties. The most recent version of this process [46] seems to have overcome the deficiencies of earlier processes [44-45]. It yields a hollow, three component, superconducting glass fiber, having an outer glass sheath and an inner superconducting sheath, both surrounding a hollow core (see Figure 11 b). These fibers possess the requisite density and microstructure to effectively carry an electrical current while also possessing the flexibility and strength required for commercial uses. The first step in this process involves melting a superconducting composition, e.g., Bi-Sr-CaCu-O, and a leachable glass composition, e.g., 75% Si02, 20% B203, and 5% Na20 (by weight), in a double bushing (Figure 11b). The gas outlet of the double bushing is centrally located within the superconducting core portion. Concentric streams of molten glass and superconductor are discharged from the bushing, converge below the tips, and are drawn through tension applied from the winder. The resulting fiber is a hollow preform with an amorphous superconductor core but it does not yet possess superconducting properties. The as-produced fiber preform is unwound from the package and is first annealed at 700950°C, depending upon the superconducting system used to convert the core material from its amorphous state into a crystalline state capable of superconducting properties. The superconducting fiber that emerges from this heat treatment is then coated to prevent microcracking and moisture degradation. Alternatively, the fiber preform can first be pretreated at a temperature ranging from 400 to 650°C to promote a phase separation in its sheath into two separate phases. Phase separation converts the glass composition into a chemically active boron-rich phase and a

160

Chapter 6

chemically inert silica-rich phase. The boron-rich phase can be removed by chemically leaching in an acid bath, leaving behind a fiber preform with a microporous silica-rich surface. As before, however, the amorphous core ofthis fiber must still be converted into a crystalline, and therefore superconducting, core between 700 and 950°C. Hollow superconducting glass fibers made by this process can be 5 to10km long. They have high current carrying capacity (>10 4 amps/em'), high mechanical strength, substantially circular cross sections in these lengths, and diameters ranging from 10 to 100 IJm. Accordingly they can be easily wrapped about small cylinders to form coils for motors. (e) Solid side-by-side bicomponent glass fibers

A centrifuge process for making dual or bicomponent glass staple fibers has recently been described [56]. Two separate glass melts, each having a different composition and therefore viscosity, are supplied toeach round cross section tip in a typical multifilament bushing. The melt streams meet under the tips, fuse, and as they cool, yield side-by-side bicomponent glass fibers with essentially round cross sections. The resulting fibers develop an irregular crimp since a differential stress develops at the interface of the components. A related process [9) yields a mixture of single and dualglass insulation staple by simultaneously centrifuging single and dual melt streams. Backscattered electron images (BEl) show that the individual bicomponent fibers have cross sections ranging from round to oval and low, but widely variable, fiber diameters (3-10 IJm). Individual fibers can split at their dual glass interface either during manufacture, or in subsequent processing and handling, thus producing fine, low diameter chaff considered tobe undesirable in use. An individual cross section is shown in Figure 12. Component (A) has a different composition than component (8), and the interface between both components is clearly discernable.

Figure 12. Backscattered electron image (BEl) of a polished cross-section of a commercial bicomponent glass fiber (4000x). Courtesy ofMicron Inc., Analytical SelVices, Wilmington, DE.

Chapter 6

161

Successive secondary electron image (SEI) profiles can be used to determine differences in composition between composition (A) and composition (8). This isaccomplished by traversing a polished cross section of the fiber at a right angle to the interface between the two components, as illustrated for Ca and Mg (Figure 13). Accordingly, the Mg (or MgO level) in this sample is higher in component (A) than in component (8), and the Ca (or CaO) level is higher in component (8) than in component (A). In addition, the Si02and Na20 levels (not shown in Figure 13) were also higher in component (A) than in component (8) and the Ab03 levels were lower in component (A) than in component (8); only component (A) analyzed for potassium, and only component (8)exhibited boron.

9.00 E-02 7.20 E·02 c 0

~

Ca

5.40

~ E·02

~

E 3.60 Ql E E-02

...... . . .

~

Ql

tn

Mg

1.80 E·20 0.00 E-010.00

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

4.00

8.00 12.00 Distance, urn

16.00

20.00

Figure 13. secondary electron image (SEI) profiles of Ca and Mg - determined by traversing the cross section of the two fiber components at a right angle totheir interface, indicate that the MgO level is higher in component (A) than incomponent (B), and that the CaO level ishigher incomponent (B) than incomponent (A). Courtesy ofMicron Inc., Analytical Services, Wilmington, DE.

In summary, the individual components of the dual component fiber differ with regard to their Si02, MgO, CaO, Ab03, 8203, K20, and Na20 content. The differential stress between the two components during fiber formation produces a three-dimensional, crimped, non-straight fiber geometry. In thermal insulation batts, a main application, arrays of crimped bicomponent glass fibers offer higher bulk and loft than possible with a comparable array of straight fibers. Arrays of crimped fibers can trap more air in their dead spaces and therefore offer higher thermal insulation than comparable arrays ofstraight fibers having the same basis weight.

162

Chapter 6

Arrays of crimped staple fibers can be processed on conventional textile staple process equipment [58]. In the carding process, crimped staple fibers are mechanically opened, aligned and combed and converted into non-woven webs. In the needling orneedle punching process, arrays of crimped staple fibers are mechanically interlocked, densified and consolidated by repeatedly punching through the fiber batt with barbed needles. In the air lay process, crimped staple fibers are separated, aligned and consolidated in a non-woven web with a high velocity air flow. In these conventional textile processes, crimped bicomponent glass staple fibers aim to compete with commodity fibers, e.g., cotton and polyester staple, and premium high performance fibers, e.g., Nomex staple [58].

6.4 High temperature silica glass fibers Silica glass fibers are made by one of three generic processes. Ultrapure silica fibers (99.99 99.999% Si02) are formed either by downdrawing from preform rods (Chapter 4) or dry spinning from a sol-gel (Chapter 5). Pure silica fibers (99.5% Si02) are dry spun from viscous water glass solutions (Chapter 5) and high silica fibers are made by acid leaching ofselected silicate fibers (this chapter). Silica fibers, irrespective of process, are amorphous, including those called quartz fibers. Ultrapure and pure silica fibers retain useful strength up to 10001100°C. In contrast, high strength (HS) glass fibers such as S-glass (Chapters 6.2) retain useful strength to800-900°C. Leached high silica fibers are limited intheir use to
=

~n

2a '

Figure 5. Fiber structures and parameter definitions for chromatic dispersion studies. From Figure 4.2, L. B. Jeunhomme [6], Single-Mode Fiber Optics, Principles andApplications, Marcel Dekker, Inc., New York. NY(1983); with permission.

Waveguide dispersion may thus be controlled by appropriate index profile design through its influence on V, a parameter dependent on the specific index profile chosen, as well as b, and subsequently

v( d 2 (Vb)/dV 2 ). The total dispersion isthen = -A. d

D(..1)=D +D //I

wg

C

2n(..1)

d 2(Vb) _ n2t1 v d..1 2 d dV2

(13)

The total dispersion will be zero when

~d c

2n(..1) d..1 2

= _ n2t1 V d 2(Vb) d

dV 2

(14)

From knowledge of the dependence of V and b on the fiber core diameter, A, and the dimensions of the depressed index cladding (if applicable), one may design a fiber whose wavelength ofzero dispersion has been shifted toa desired value [3] [6]. Not only will the wavelength of zero dispersion be shifted, but at a local maximum, so introduced, the dispersion will vary little over a window of wavelengths. The dispersion behavior for actual fibers of varying designs is shown in Figure 6. Note that for seemingly similar index profiles, the behavior varies widely. This places strict requirements on the precision offiber manufacture.

Chapter?

177

Total dispersion is extremely sensitive to variations in the index profile of both the core and cladding parameters as illustrated by the behavior offibers denoted W2 and W3. Though the diameters of their cores differ by only 0.1 urn, and the diameter of the depressed inner cladding even less, their dispersion behavior shows significant difference. Not only are the wavelengths of zero dispersion different, but W3 exhibits almost no dispersion over a window approximately 50 nm wide, centered at about 1530 nm. Clearly, this configuration is more desirable for operation at1.5urn than the fiber designated W2. Intermodal dispersion inmultimode fibers causes each mode totravel ata different speed due to different group delays, caused by different path lengths between modes. For step index multimode fibers, the maximum spread in delay time between the fastest and slowest modes isgiven by M

= ~(n, -n] C

)(l_lJ

(15)

VIII

where Vm is an integer corresponding to the highest mode allowed for a given value of V, approximately Vm = 2 V/'" This spread in delay time will cause pulses consisting of many modes tobe broadened.

10

. E

:.
nt. The differential phase velocity is accompanied by a difference in local group velocities for the two polarization modes. This is responsible for broadening pulses as seen by a group delay time L'H per unit length between the modes. M L

=~ ({3" _ {3f) = linefj dco

c

_ wdlinefj cdoi

(23)

MIL is referred to as pulse mode dispersion (PMD). The length dependence for this dispersion islinear for short lengths and increases as the square root oflength for long spans.

Various strategies have been implemented for the elimination and reduction of PMD in fiber manufacture and cabling [11]. In addition to tailoring preform dimensions to improve geometry, a strategy of twisting the fiber during draw to introduce mode mixing has been successful in reducing PMD. For certain sensor and network uses, polarization-maintaining fibers are produced [12]. These use birefringences produced by asymmetric stress on the

Chapter 7

180

core either by deforming the core or by introducing stress-producing rods beside the core in the preform before draw.

--

/Ideal

-,

Elliptical cladding

Non symmetric stress

@

~:®

(1)8

He~l

(a)

Hd;l

8

Geometrical (2)(;;\

Stress

=0=~ ~ Lateral stress

Bend

Twist

(b)

Figure 7a. Anatomy ofa real fiber, b.Intrinsic and extrinsic mechanisms offiber birefringence. From C. P. Poole and S. Nagel, [10j, Polarization effects inlightwave systems. in Optical Fiber Telecommunications, IliA. pp. 115-161, I.P. Kaminow and T. L.Koch, Eds., Academic Press, San Diego, CA (1997).

7.3 Fabrication ofoptical fibers The earliest work on optical fibers began in the 1960s to produce multimode fiber with acceptably low loss. To then, typical optical quality glass exhibited losses on the order of 1000 dB/km. Losses 50% alumina, including YAG glass fibers, have recently been reported (Chapter 4.4). In contrast, polycrystalline alumina, YAG and zirconia fibers are accessible bya sol-gel route orbya slurry process. 8.2.1

General considerations

At high temperatures, a-Ab03, i.e., corundum or sapphire, is the stable form of alumina. Transition aluminas (Figure 1)are formed between 300 and 1100°Cthrough heat treatment of

Chapter 8

208

aluminum hydrates, for example, gibbsite (a-Ab03·3H 20) , bayerite (13-Ab03·3H 20), boehmite (a-Ab03·H20) and/or diaspore (13-Ab03·H 20) . The sequence of formation of transition aluminas depends on the starting material and is affected by crystallinity, grain size, heating rate, impurities and/or additives.

I

I

Gibbsite

~[

I

I

~

Eta

.1 'I

I

I

o

I

Kappa

I Delta

Gamma

Diaspore

I

I

I

Chi

Boehmite Bayerite

I

I

Alpha

I I Alpha

ITheta Alpha

Theta

Alpha alumina

I

I

I

I

I

I

200

400

600

800

1000

1200

Temperature, °C

Figure 1. Dehydration sequences ofalumina hydrates inair. Enclosed area indicates range ofstability and open area range oftransition [20]. Path (b) isfavored bymoisture, alkalinity and coarse particle size (100 IJm) and path (a) by fine crystal size « 10 IJm); reproduced with permission of the American Ceramic Society, PO Box 6136, Westerville, Ohio 43086-6136.

The low temperature transition aluminas including the chi, eta and gamma phases are poorly crystallized and contain residual water. The eta and gamma phases display a spinel structure with vacancies. The high temperature transition aluminas, comprising the kappa, delta and theta phases, are formed at 800-1000°C. Finally, the transformation of fine theta (or kappa) transition alumina at 1100·C into a-alumina is reconstructive and occurs with an increase in grain size (10). Transition aluminas can be stabilized by addition ofsmall amounts ofdifferent oxides, the most commonly used in alumina based fibers being silica, in order to shift the formation ofa-alumina toward higher temperatures (11) (12) [13]. Mullite (3Ab03·2 Si~) contains 71 .8wt.% Ab03 and 28.20 wt.% Si02, and displays a range of homogeneity assigned to a substitution of AI:l+ for Si 4+ in tetrahedral sites with formation of oxygen vacancies (Figure 2a) (14) (15). Assuming a generic formula with an atomic structure of AbVl [Ab-+<xSb"2'I'V 010-'. the value x = 0.25 corresponds to the nominal formula 3Ab03·2Si02 (with 71.80 wt.% Ab03) and the value x = 0.40 to the formula 2AbCh·Si~ (with 77.24 wt.% AbCh). The actual homogeneity range of mullile and the nature of its melting (congruent or incongruent) are still controversial areas.

209

Chapter 8

Wt. %

20

0

60

40

80

Liquid

Mullite 55 + Liq ...

2000

~

!i

e :>

1850 0

/

"

/

/

/

Mullite + Liq

/

Cor + Liq

/

18400

Cor + mullite 55

15950

1600

I I

5i02+ mullite 1400

/

5i02+ Liq

1800

c.

E ~

100

-{ -MullneS5 (a)

I 0

60

40

20

80

100

Mol, %

5i02

AI203

1500 .--------------,,~-------,

Mullite

SOOL-

o

'-10

-'--

(b) -'-_.....J

20

30

% 5i0 2

Figure2. The binary Ab03-Si02 system. (a) Equilibrium phase diagram by S. Aramaki and A. Roy, J. Amer. Ceram. Soc., 45, 229-242 (1962): reproduced with permission of The American Ceramic Society, Westerville, OH. (b) Non-equlibrium phase sequences from heating sol-gel silica aluminas [44): reproduced with permission of the Institute ofMaterials, London.

Crystallization of mullite from stoichiometric alumina-silica mixtures, referred to as mullitization, occurs at different temperatures depending on the scale at which the constituents are actually mixed [16]. For conventional powders (grain size ",,1 I-Im), the temperature for complete mullitization is 1600-1750°C, whereas that for alumina-silica gels is 1300-1450°C and even lower (1000-1100°C). Further, the formation of mullite from gels is considerably favored by addition of boron oxide. The addition of a second phase at grain boundaries impedes grain growth and creep in a-alumina. Examples are Nextel 720, a corundum/mullite fiber [17] [18] and PRD-166, a corundum/zirconia fiber [19] [20]. 8.2.2 Processing of alumina based fibers The processes used to prepare alumina based fibers by the solution route start with precursors which are either an aqueous solution of an aluminum salt or a solution of an organoaluminum compound in an organicsolvent. The level of viscosity required for spinning is achieved by properly controlling the degree of hydrolysis/polycondensation of the precursor

210

Chapter 8

orland by adding spinning aids such as polymers with high molecular weights. In the related slurry route, a fine orultrafine alumina powder is dispersed in the liquid precursor, in order to increase the equivalent oxide content and to lower fiber shrinkage during the dryinglfiring steps. (a) Polycrystafline alumina fibers Aqueous solutions of basic aluminum salts, especially basic aluminum chloride, are the most commonly used aluminum precursors. Depending on the synthesis conditions, different species are observed in the solutions [11) [21). The monomeric hexahydrated cation, [AI (H20)6j3+ isstable only inhighly acidic solution. Atan OH/AI molar ratio of 3, AI(OHh is formed and precipitated whereas atlower ratios, i.e., 2 to 2.5, polymeric species such as the complex pseudo-spherical cation [AI O. AI,2(OHh.(H20),2t, also referred to as the AlB cation, and larger polymeric species are formed. Aqueous solutions of aluminum chloride have relatively low viscosities up to a concentration threshold [11) [21). Above this threshold, the viscosity increases dramatically (Figure 3) and a glass-like solid is then formed. Concentrated AICh solutions undergo aging accompanied by a further viscosity increase. However, adding a high molecular weight organic polymer such as poly(ethylene oxide) or partially hydrolyzed poly(vinyl alcohol) is a much more effective way ofincreasing the viscosity.

60

3

4

<Jl

50

ai

a.. 2 ;E-

'iii <Jl

40

ai

a..

;E- 30

'w 0 Co)

s

<Jl

81

s

5

<Jl

0 24

26 28 30 Oxide content, %

32

20 10 0

10

14

18

22

2Q

30

Oxide content, %

Figure 3. Relationship of viscosity and concentration ofAI20J spinning dopes prepared from aqueous solution of basicaluminum chloride containing poly(vinyl alcohol) with PVAJAI20J mass ratios of(1)30/70 ; (2)25/75 ; (3)20180 ; (4) 18/82 and (5) 15/85 (shear rate = 40 s'), The inset shows the viscosity/concentration curve for an aqueous solution of basic aluminum chloride without addition of PVA (31); reproduced with permission of Huthig and Wepf Verlag.

Chapter 8

211

Continuous, polycrystalline a-alumina fibers, e.g., Fiber FP or Almax, are produced by the slurry process [22-26]. The spin dope iseither a dispersion ofa fine a-Ab03 powder (e.g., 0.5 urn mean size) [23-24] oran ultrafine y-transition alumina powder (0.02 urn mean grain size) [25] in an aqueous solution of a basic aluminum chloride, for example, the water soluble Ab(OH)sCI·2H20 salt containing various additives, including (a) high molecular weight polymers acting as spinning aids, (b) sintering aids and (c) a grain growth inhibitor such as MgO (as magnesium chloride). The dope isconcentrated by heating under vacuum and isthen dry spun through a spinneret. The green fibers are partially dehydrated in a drying column and are taken up on a collapsible reel [23-24] or further processed continuously on line [251. They are prefired at 600-800°C and subsequently sintered at 1350 - 1500°C in a furnace or/and a propane/oxygen flame. Fiber FP is 18-20 ~m in diameter whereas the Almax fiber obtained from ultrafine y-alumina powder displays a smaller diameter (",,10 pm). In both fibers, alumina is present as corundum. a-Alumina fibers can also be prepared, at least on a laboratory scale, via pure sol or solution routes from different precursors including aqueous colloidal sols [26-27J or organoaluminum polymers [28-30], as illustrated by the two following examples. (1) Colloidal sols can be prepared from aluminum chloride (or nitrate) by dissolving AICh (or AI(N03h-9H20) and aluminum pellets in water under reflux, followed by filtration [27]. The hydrolysis/polycondensation reaction is continued until the viscosity of the precursor is suitable for dry spinning. The green fibers are dried, prefired to 800°C and sintered at 1300°C in air. Crystallization of the gel occurs above 700°C. First transition aluminas are formed and then converted to a-alumina [27]. (2) Polymeric organoaluminum precursors can be prepared by the addition of glacial acetic acid to either pure ethyl 3-oxobutanoatodiisopropoxy-aluminum (EOPA), or to a mixture of EOPA and tri-sec-butoxyaluminum (SBA) without solvent and water [28-29]. The reaction mixture is refluxed at190°Cand isthought toproceed by the following equations, where EOB isthe ethyl 3-oxobutanoato didentate ligand, CH3COCHCOOC2Hs. AI(EOB) (0 ipr)] + A cO ff _ _ Ro f1 1-O} + 'r-on + AcOipr

(4a)

EOB AI(EOB)(O jP r)] + A1(0 '''BU)3+ AcOH - -

RoBl-~n +ROH + AcOR

(4b)

EOB

After the volatile species are vaporized under vacuum, a pale yellow polymeric precursor is obtained whose viscosity is time independent and could be controlled, without adding spinning aids, by the amount of added acetic acid, the EOPAISBA molar ratio and the temperature. The green fiber, calcined at500°C, is amorphous and undergoes crystallization within the temperature range 800-1300°Ctoa-alumina. (b)

Jransition alumina fibers

Transition alumina fibers are obtained by the sol-gel route, as long as the firing temperature is

212

Chapter 8

low enough to avoid the transition to a-alumina, i.e., 1000-1100°C for pure alumina (Figure 2b) [11-12) [31). Calcination yields a fiber that is initially amorphous and porous. The crystallization of this fiber occurs through two exothermic main transformations, the first being the formation of the transition aluminas starting with the n-phase and the second that of aalumina (Figure 1). Simultaneously, the fiber shrinks. If a shrinkage resistant transition alumina fiber is desired, sintering of the porous fiber should take place in a temperature range that is actually quite narrow for pure alumina. Thus, most commercially available transition alumina fibers are stabilized with oxide additives such as silica. The addition ofsilica (Figure 2b) considerably enlarges the apparent stability domain of transition aluminas, the formation ofa-alumina being shifted toalmost 1250°Cfor only 5 wt.% Si02 [34). Table II. Commercial silica-alumina fibers Fibers

Producers

Nextel 312 440 480 550 Altex

21
380 GPa) from mesophase pitch. Mesophase is a liquid crystal phase formed in the transformation of organic precursors to carbon. The chemical pretreatment of the pitch is complex and the potential cost advantage ofusing a pitch precursor isstill not fully realized [1-8). 9.1.2 Elemental carbon Elemental carbon exhibits two main crystalline modifications. Diamond is cubic and graphite is hexagonal. Fullerene, a new form ofcarbon, will not be discussed; fullerene related carbon nanotubes are discussed inChapter 3.

234

Chapter 9

In graphite, the atomic layers, referred to as graphene layers, are stacked in the c-direction. Within these layers, the C-C bond is extremely strong. Conversely, the graphene layers are only weakly bonded to one another. The two-dimensional character of the graphite crystal structure is responsible for the strong anisotropy, which is observed for most of the mechanical and physical properties ofthe single crystal ofgraphite. Since the layers in graphite are only weakly bonded along the c-axis, they can easily slide over one another orland undergo in-plane rotation. These motions lead to turbostractic carbon structures, which consist of roughly parallel and equidistant graphene layers rotated randomly. The terms: graphite material or graphite carbon should be used only if the interlayer spacing is less than 0.344 nm; otherwise the terms: carbon material ornon-graphitic carbon are preferred [3}. Most fibers fabricated from organic precursors should be referred to as carbon fibers, even when they exhibit rather high Young's moduli, and not as graphite fibers, with the exception ofsome fibers derived from mesophase pitch [3}. Graphitization is a solid state transformation of thermodynamically unstable non-graphitic carbon into graphite bythermal activation. Non-graphitizable carbons cannot be transformed into graphitic carbon by heat treatment up to 3000 K under atmospheric or lower pressure. Graphitizable carbons can be converted into graphite. Generally, non-graphitizable carbons are produced from non-fusing organic precursors, whereas most graphitizable carbons usually pass through a liquid crystal intermediate (the mesophase) during the carbonization process [9].

1200

Cr(GPa)

5r(llt5GPa") 1000

. e

11.

!i 's

~

5" = ~=

5.. =

800

5'2 = 5'3 =

98 2750 25000 ,16 ·33

800

= 0.722 sin'Z [44-45]. The Young's modulus of PAN based HM carbon fibers varies linearly with the reciprocal ofthe carbon layer mean size, La [35]. Values of some elastic characteristics for various mesopitch based and PAN based carbon fibers can be computed from those ofsingle fiber composites with a resin matrix [44].

1000 \\\\\\ I Pitch

'"

800

0.. (!)


So

(16c)

0"02

The data for recent Nicalon NL-200 fibers of a given gauge length (L=25 mm) are spread over a wide range from 1500 to 4000 MPa. The average is about 3000 MPa (Figure 12). In a first approximation, these experimental data can be fitted to a unimodal two parameter Weibull distribution. There are often a few data points, usually at low stress levels, which do not actually fall on the straight line predicted by equations of type 15a or 15b, suggesting that more than one flaw ispresent in the fiber.

••

0.90

•

0.60

l!! ::l

S (;

I

0.30

.~

0.10

•

D..

•

• • 1500

Nicalon NL200

2500

3500

Tensile strength. MPa

Figure 12. Weibull plot of tensile strength data for Si-C-O (Nicalon NL 202) fibers. with a unimodal Weibull distribution (80); reproduced with permission from VSP, Zeist, NL.

About 96% ofthe fracture origins in Nicalon NL200 fibers are surface flaws and only 4% are internal flaws. Since the internal flaws are minimal, the corresponding data can be treated as censored data and the tensile strength of the fiber can be statistically analyzed (Figure 12) with a unimodal Weibull distribution, where m = 4.5 and cro = 2670 MPa [80]. Data for an older Nicalon fiber. however, gave a better fit with a bimodal Weibull distribution corresponding to a family ofsurface flaws (m, = 3.64; crOl = 4.64 GPa) and a familyof internal flaws (rn, = 9.41 and cr02 = 5.08) [79]. More recently, the observation of two partly concurrent populations offlaws in Nicalon NL200 fibers has been attributed [82] to extrinsic and intrinsic flaws. Accordingly, a family ofextrinsic flaws (severe flaws at the fiber surface) is responsible for failure at the lowest stress levels

287

Chapter 10

(with m2 = 1,92), and a family of intrinsicflaws (located both atthe surface and in the volume) control failure athigher stress levels (with m, = 4,5), The identification of families of flaws, which are responsible for the failure of a batch of fiber specimens, is not straightforward. Before testing, the specimens must receive a soluble damping coating ofa wax ora polymer such as 1-3 polypropanediol, orthey must be tested in a liquid medium such as glycerol to absorb the shock wave energy due to specimen bursting atfailure. Under such precautionary conditions, the primary failure surfaces can be recovered and the origin ofthe fracture determined [79-80]. The surface flaws are mainly microvoids and microcracks induced by spinning and mechanical abrasion whereas the internal flaws are microvoids or inclusions originating from the PCS precursor fiber or formed during pyrolysis. The failure surface often shows the fracture origin-mirror-mist-hackle-crack branching morphology characteristic of brittle materials. The depth a and the width c of surface flaws in Nicalon NL200 fell in the ranges 0.09 < a < 0.6 IJm and 0.25 < c < 1.7IJm, respectively. The surface flaws were modeled as semi-elliptical cracks, allowing calculation ofthe fiber toughness, KIC, by the Griffith equation: I

O"R

K/c

= y' (J(Q c/

12

(17)

where Y isa geometric parameter depending on the shape ofthe defect (here, Y = 0.8331 , for alc = 0.3). K,C, calculated from failure stress (crR) and flaw depth (a) data, is of the order of 1.55 -1.9MPa m". The tensile strength of Si-C-O fibers decreases after exposure to elevated temperatures. When Nicalon NL 200 fibers are exposed for 1 hour to 1300°Cin argon (P = 100 kPa), their mean tensile strength and scale parameter, cr., decrease by 45% while their Weibull modulus remains unchanged [80-83]. Fibers exposed to more severe conditions (e.g.,for 5 hours in a vacuum at1500°C)are so weak that they cannot be tested. Finally, the fact that oxygen-free fibers maintain their tensile strength under similar conditions relates to the absence of silicon oxycarbide and its decomposition process. 10.7.2 Athigh temperatures The high temperature (HT) properties of SiC based fibers depend upon their pyrolysis temperature, their thermal stability, the time and the test atmosphere. When mechanical tests are performed at a test temperature that is higher than the pyrolysis temperature, some change occurs in the composition orland microstructure ofthe fibers which may strongly affect their mechanical behavior. (a) Tensile tests

Young's modulus and tensile strength of Si-C-O and Si-C fibers decrease when the test temperature is increased (Figure 13). The room temperature tensile strength of the Si-C fibers (Hi-Nicalon) is higher than that of Si-C-O fibers (Nicalon NLM 200), but the two fibers have almost the same strength, (1200 MPa), when tested at 1400°C. At any test temperature, the modulus ofthe Si-C fiber is higher than that ofthe Si-C-O fiber [31] [84]

288

Chapter 10

4 (a)

3 ca

Q.

e

~

15>


HC/ J

Hp

I

(2)

J

J

+4n CH~H2

n

Hydridopolysilazane (HPZ) can also beused as a precursor to produce Si-C-N(O) fibers. It is obtained by the exothermic reaction of trichlorosilane HSiCb with hexamethyldisilazane [21] [18]. Condensation yields a HPZ polymer with a (Si-N)n backbone; carbon is in the pendent methyl groups. HPZ polymers have a formula close to (SiH)39J(Me3Sih42(NH)373N226 [21]. Fibers must be drawn from the HPZ melt in inert atmosphere since HPZ is sensitive to moisture and oxygen. Chemical curing is achieved by exposing the green fibers to an argon stream containing trichlorosilane vapor ata temperature below the softening point: (3) In this step there is a decrease in carbon content. The cured fibers are pyrolyzed under flowing nitrogen at 1200°C. Their chemical composition after pyrolysis is about 58 wt. % Si, 10% C, 29% Nand 3% 0 [21-22]. The carbon content of the fibers depends on processing conditions. The empirical formula assigned tothe fibers is close to SiC0403Nl00300091 [22]. The ammonolysis of dimethyldichlorosilane offers another route to PSZ precursor fibers and ultimately silicon nitride fibers. Ina first step, Me2SiCh and MeHSiCh are mixed in nearly a 1:1 molar ratio inbenzene. The mixture is treated with NH3, NH4CI precipitates, and the solvent is removed. The ammonolysis product, a viscous polymer, is converted in a second step at 400°Cinto melt spinnable polymer, and is melt spun. The resulting PSZ precursor fiber has a (Si-N), backbone. Carbon is present in pendent methyl groups, and the empirical formula of the fiber is SiC 17No90H s7 [23]. (b) From polycarbosilazane (PCSl) fibers InPCSZ precursors, carbon atoms are in the polymer backbone. The C-Si-N bonds, yielding tetrahedral Si(C,N)4 units after pyrolysis, are already present. Most oligomers with carbon in the backbone [18] pyrolyze to yield ceramics with an undesirably high carbon content. In contrast, a copolymer route that combines a given silazane monomer with dimethyldichlorosilane (the starting material in the Yajima route) can yield melt spinnable

Chapter 11

301

PCSZ precursor melts and precursor fibers with a -Si-CH 2-NH-Si- backbone (via the Kumada rearrangement) having tailored Si/C/N ratios [8-11] [18] [24]. In the copolymer route, a polysilasilazane (PSSZ) oligomer is prepared by sodium copolycondensation of dimethyldichlorosilane (Me2SiCb) and 1,3-dichloro-1 ,3dimethyldisilazane (CIMeHSi-NH-SiHMeCI) ina 1:1 molar ratio inboiling toluene (Equation 4). Insolubles are removed by filtration, the product is concentrated and NH3is bubbled at O·C into the solution to break all residual Si-Cl bonds. Ammonium chloride is removed by filtration, solvent and low boiling oligomers are distilled offand a low molecular weight PSSZ is obtained ( M w = 2500). It is converted into soluble, fusible, high molecular weight PCSZ (M w =15 000; soft. point =170·C) by heating at 300-370·C. The PSSZlPCSZ conversion occurs with a viscosity increase and the evolution ofhydrogen. CH 3 x

CH3

I

I

+ (I-x) (CH) ] SiCI]

Cl-Si-N-Si-Cl I

H

I

H

I

H

Naj Refluxing toluene

f

C~CHj-

I 3 I CH si- N -Si

I

H

I

H

3

I

Si

I

H

(4)

+ 2 NaCI

I

x

CH 3

l -x

Continuous, 25 IJm diameter monofilament PCSZ precursor fibers, referred to as PCSZ-II, were melt spun [8-9]. Their empirical formula was close to SiC122No45H oOoOJ. They were either chemically cured by oxidation [9] or physically cured by irradiation with y-rays. Two successive weight losses occur during the pyrolysis in argon or nitrogen. The first, at 25450·C, is related to the volatilization of light oligomers and the second, at 450-950·C, was assigned to an evolution of hydrogen and CH.. The pyrolysis yields amorphous fibers whose chemical compositions remain constant within the 950-1400·C temperature range [9] [11]. The empirical formulas for fibers pyrolyzed at 1400·C, are SiCol8 N039 0061 for the fibers cured with oxygen, and SiCo.93No4600 05 for the fibers cured byy-radiation. The evolution of CH. during the pyrolysis step and the empirical formulas of pyrolyzed fibers suggest that carbon isonly partly present in the backbone of the polymer (assuming that CH4 isformed from pendent methyl groups). Finally, the NISi ratio in PCSZ based fibers, Le. 0.460.49, islower than that, 0.90-1 .00, inthe fibers derived from PSZ. 11.2.2 Structure and properties Si-C-N-O and Si-C-N(O) fibers are amorphous (Figure 1) but the molecular composition is still a matter of speculation. Generally speaking, they can be depicted as consisting of a continuum of tetrahedral units wherein silicon is surrounded by carbon, nitrogen and oxygen atoms; the latter are connected by their summits. Furthermore, some free carbon, which results from the pyrolysis of the pendent methyl groups, is usually present asBSUs.

302

Chapter 11

16000 C,Ar

jJ\

1500 C, Ar

-.

0

A

1400 C,Ar 0

16000 C, N2 I

20

I

~ I

I

I

60

40

I

I

80

29, degrees

Figure 1. XRD-paltems of Si-C-N-O fibers produced from PCSZ precursor cured byoxidation and pyrolyzed at various temperatures inargon ornitrogen [10]; reproduced with permission.

(a) Fiber Structure

Si-C-N(O) fibers derived from HPZ precursor fibers are nanoporous and heterogeneous with a skin/core structure. The composition changes from SiOxCy in the external porous surface to SiNxCr in the core. The molecular formula of this fiber is close to 4 mol.% Sieh, 81 mol.% SiNxCr (x = 1.02, Y= 0.23) and 15mol.% free C [22). The presence of complex tetrahedral units is supported by the 29Si NMR spectrum which shows a broad signal covering the chemical shift region expected for silicon oxycarbide, silicon oxynitride and silicon carbonitride units [21). The occurrence of free carbon, expected from the nature of the precursor, is supported by the C 1s XPS pattern [22). The structure of Si-C-N-O and Si-C-N(O) fibers derived from PCSZ is quite similar to that depicted for fibers derived from HPZ with, however, some noticeable differences. In Si-C-N-O fibers derived from PCSZ, the oxygen concentration is relatively high (e.g., 12-20%) and oxygen is distributed homogeneously over the entire cross section whereas it is concentrated in the skin offibers derived from HPZ [9-10). Thus, the tetrahedral units in the continuum are probably SiOxNyC, units rich in oxygen. In Si-C-N(O) fibers derived from PCSZ and cured by irradiation, the concentration of oxygen is very low and the continuum may consist mainly of SiNxCyunits. (b) Thermal stability

Si-C-N-O and Si-C-N(O) fibers remain amorphous up to 1300-1400°C (Figure 1). The N heteroatom in the fiber has a beneficial effect on the thermal stability of the amorphous

Chapter 11

303

tetrahedral continuum. The glassy Si-C-O phase in Nicalon fibers undergoes decomposition/crystallization at1100-1200·C. The core of the fiber derived from HPZ retains its amorphous state and elemental composition, even after long annealing in argon at 1300 or 1400·C. However, noticeable structural change occurs near the fiber surface where oxygen is concentrated, with formation ofcrystalline silica (cristobalite), SbN 20, I)-SiC and channels ofporosity [22). If it is assumed that silica is formed first by decomposition of the outer silicon oxycarbide skin, then SbN20 and SiC might result from slow interactions in the solid state between silica and the SiNxCy core. SiC could also result from decomposition ofthe core near the skin/core interface

0,---__ Organiclinorganic transition

•~

{

20



c:

:l

~

50

600

Si-C-N-O

800

1000 1200 1400 Pyrolysis temperature, °C

1600

Figure 3. Room temperature strength and modulus of y-ray cured Si-C-N fibers and oxygen cured Si-C-N-O fibers [11]: reproduced with permission.

Growth kinetics for the silica sheath obey parabolic laws (Equations 24a and 24b, Chapter 10). Thermal variations of the parabolic constants follow an Arrhenius law (Equation 25, Chapter 10). The apparent activation energy is 170 kJ/mol and the pre-exponential factor is Ko= 3 X 106 nm2 S·l for ~ [12]. Thus, the oxidation of Si-C-N-O fibers is rate controlled by diffusion. Compared with the oxidation kinetics of the PCS based Si-C-O fibers, those of PCSZ based Si-C-N-O fibers differ in important points. (1) The oxidation parabolic constant, Figure 4, is lower. (2) The activation energy for Si-C-N-O fibers, E.= 170 kJ/mole, is higher than that for comparable Si-C-O fibers, E. = 70 kJ/mol [26]. (3) The parabolic constant for SiC-N-O fibers ishigher atlow temperatures than those for bulk SbN4 and CVD-SbN4. However, the differences strongly decrease as temperature israised.

306

Chapter 11

HP-SiC

3

Bulk

Si3N4

in

'"Ec

-1

Si-C-N-O fiber

C\i ~

c

...J

-3

-5

-7

' - - _ - ' - - _ . . . L - _ - ' - _ - - ' - _ - . l . _ - '_ _"----I

0.5

0.6

0.7

0.8

0.9

Temperature,103K-l

Figure 4. Arrhenius plots ofthethermal variations ofthe kinetic parabolic constant ks, fortheoxidation of Si-C-Nfibers derived from PCSZ and related Sibased ceramics, according toref. (12); reproduced bypermission.

o

(e) Other properties

Si-C-N-O and Si-C-N(O) fibers have a low density (2.3-2.4 g/cm 3 for HPZ based fibers and 2.5 g/cm 3 for PCSZ based fibers). They also exhibit a low coefficient of thermal expansion. Its value is 3.0-3.9 x 10-6 " for the HPZ based fibers in the 350-1000·C temperature range. PCSZ based Si-C-N-O fibers exhibit low electrical conductivity with semiconducting behavior and alow apparent activation energy [10J. 0C

11.3 Si·N·O and Si·N fibers Continuous Si-N-O and Si-N fibers, almost free of carbon, can be prepared by pyrolysis of polycarbosilane precursor fibers in ammonia. The precursors have been discussed for fabricating Si-C-O and Si-C fibers (Chapter 10). 11.3.1

Processing

The pyrolysis of oxygen cured PCS fibers yields continuous Si-N-O fibers, whereas that of PCS based fibers, cured by irradiation under anaerobic conditions, affords continuous Si-N fibers [7J [27J. Their main advantage istheir extremely low electrical conductivity. (a) From Yajima type polycarbosilane (PCS) fibers

Chapter 11

307

When a Yajima type PCS is heated in an ammonia gas flow, it undergoes a large weight loss that is partly related to phenomena which have been already described for PCS pyrolysis in an inert atmosphere (Chapter 10). It is also related to nitridation of the Si-CH3, Si-CH2-Si and Si-H bonds which takes place between 400 and 800°C with a sharp decrease in the carbon content of the fiber. When the nitridation is complete, the fiber contains almost no carbon. 80th the pendent methyl groups and the methylene groups in the backbone have reacted with ammonia [28). The nitridation of PCS can be followed from the change occurring in the 29Si high resolution NMR signal [29]. Starting from a PCS fiber with a composition close to SiC193H 47IOool, the material is first cured by E-beam irradiation under anaerobic conditions. In a second step, the cured fiber is pyrolyzed in a gas flow of ammonia, yielding at1200°Can amorphous Si·N ceramic fiber with a composition close to SiN1.26CO.olHo.ls. When the pyrolysis temperature is raised to 1400°C, the fiber crystallizes. It is essentially free of carbon and oxygen and its empirical formula, SiN w orSbN m , isclose to that ofSbN 4 [7]. The composition of the Si-N-O fibers from the pyrolysis of oxygen cured PCS fibers in an ammonia gas flow depends on the curing conditions. Starting with the Yajima type PCS precursor, SiC1.93H 47IOool , Si-N-O fibers were obtained with compositions such as SiNll sOo.47 having an O/Si atomic ratio close to 0.4. These fibers remained amorphous up to 1500°C[7]

[27].

(b) From perhydropolysiJazane (PHPSZ) fibers PSZs without pendent carbon bearing groups, which are precursors to silicon nitride ceramics, can be obtained by ammonolysis of dichlorosilane, H2SiCb, or a stable complex, with a tertiary amine [18]. A stable H2SiCb/2-pyridine complex has been used toproduce Si-N fibers (Equations 5a, 5b). Pyridine acts as a catalyst in the condensation/crosslinking reaction, whereas other solvents, such as xylene, have the reverse effect [4-5]. The end product isa clear solution ofperhydropolysilazane (PHPSZ). A dry spinnable PHPSZ polymer (Mn =2000; Mw =12000) has an empirical composition close to SiN,.ooHo.950o.o3 and a structure consisting tentatively of rings linked by linear segments (Equation 5c). Green fibers are dry spun in an inert atmosphere from a PHPSZlxyiene solution. Vaporization of the solvent yields a material that does not melt or soften when the temperature is further raised and thus does not require a curing treatment.

308

Chapter 11

(5a) (5b)

+2Hz

-SiH-NH-

I

NH

I

-SiH-NH-

-N H, H/

.su:

H

':sr: H

S, k. .

Si

H H H .:. sf I

H

~ _~i_H 'N I

(5 c)

Si- H

I I N N, ./N -Si-NI:J........... / ' Si - N ' \ H/ Sl I -, Si I \ \ H ~ H H H/ \ H H H H

d

The green fibers are pyrolyzed in a flow of ammonia gas at 1000°C to complete the condensation/nitridation/crosslinking process and to increase the NISi ratio. The nitridation reaction occurs between 400 and 600°C. At600°C, the NISi ratio isclose tothat of ShN4 [30J. At 1000°C,the fibers consist of an amorphous hydrogenated silicon nitride. Finally, the fibers are densified by a heat treatment at 1400°C in nitrogen. The fibers have elliptical cross sections, and they are colorless, transparent, amorphous and almost free of carbon and oxygen. Their composition, SiN'2..0o.079Co.o'6, isclose tothat of ShN4 (or SiN1333).

(e) From other polysilazane fibers The curing of HPZ fibers byHSiCh could theoretically remove all pendent methyl groups if the reaction according to Equation 3, goes to completion. The fibers, produced bypyrolysis in a nitrogen gas flow, are almost carbon -free. Their actual composition is close to SiNuI03C040300091, typical of a silicon carbonitride fiber [22J. However, early fibers which had a low NISi ratio but much lower carbon content were found to correspond to a composition close to SiN1069C0 06900 064 [21 J. 11 .3.2 Structure and Properties Si-N-O and Si-N fibers are amorphous. The thermal stability of the amorphous continuum depends on the fiber composition, which is itself related tothe nature of the precursor and the pyrolysis temperature (Tp).

(a) Thermal stability PCS based Si-N-O fibers remain amorphous up to at least Tp = 1400°C whereas their Si-N counterparts undergo crystallization at 1300 < Tp < 1400°C [7J. For both Si-N-O and Si-N fibers, the crystalline phase which is formed is a-ShN4. PHPSZ based fibers with a composition close to SiN12l16 00.108 C0028 are amorphous when prepared at Tp = 1300°C. Their

Chapter 11

309

crystallization to a-SbN4 in nitrogen, obeys a first order kinetic law within the 1400-1500'C temperature range [6]:

da ldt=k(J -a)

(6)

where a is the weight fraction of crystalline a-SbN4in the fiber at time t; k is a kinetic constant. Temperature variations of k obey an Arrhenius law with an apparent activation energy of E. = 602 kJ/mol.

(b) Mechanical properties The mechanical properties of Si-N-O and Si-N fibers, measured in tensile loading at room temperature, depend on their processing conditions. For example, the modulus of PCS based Si-N-O fibers increases from 80 to 180 GPa as pyrolysis temperature is raised from 1000 to 1500'C. Their tensile strength shows a maximum close to 1.8 GPa when pyrolyzed between 1300 and 1400'C, and then decreases. The decrease isrelated to the onset offiber decomposition/crystallization, but, even after pyrolysis at 1500'C, the fibers still display a tensile strength of about 1 GPa [7]. 10 IJm diameter PHPSZ based fibers, which are prepared at 1300'C, have a composition close to SiN12200.'0C002, a Young's modulus of about 220 GPa and a tensile strength of 2.5 GPa (L = 25 mm) at room temperature. However, tensile strengths as high as 4 GPa are achieved when the fiber diameter is reduced to 6 IJm [5]. Annealing in nitrogen does not significantly decrease the RT strength of the fibers as long as the annealing temperature is maintained below the onset of the formation ofcrystalline a-SbN4.

(c) Other properties Si-N-O and Si-N fibers display a density ranging from 2.3 to 2.4g/cm3 and a low coefficient of thermal expansion (1.5 x 10-6 'C·l for fibers derived from PHPSZ). Furthermore, they display low electrical conductivity. The resistivity of the fibers is 5.0x 10'4Oem. Finally, Si·N·O and Si-N fibers exhibit good oxidation resistance, which is related to the formation of a silica surface in the passive oxidation regime. 11 .4 Si·B·Q·N, Si·B·N and Si·B·N·C fibers The insertion of boron into Si-N and Si-N-O shifts the onset of SbN4 crystallization to higher temperatures. This recognition has been used to produce experimental Si-B-O-N [15-16], SiS-N [17] and Si-S-N-C [17] [31] fibers. Pyrolysis in ammonia yields colorless and nearly carbon-free fibers, while pyrolysis in nitrogen or argon produces black, carbon containing fibers. The fibers are amorphous after pyrolysis at 1000'C; the thermal stability of the amorphous state is very high. For example, Si-S-N-C fibers with a composition close to SiSN 3C crystallize only at 1800-1900'C (17). 11.4.1

Processing

Two process routes are known. One starts with perhydropolysilazanes (PHPSZs), the other with trichlorosilylaminodichloroborane (TADB).

Chapter 11

310

(a) From perhydropolysilazane (PHPSZs) fibers Polyborosilazanes (PBSZs) are produced by reaction of PHPSZ with trimethyl borate (SIIB = 3) in pyridine, The solution is heated for 3 hours at 120·C. After cooling, o-xylene is added and the solvent is removed by vacuum distillation at50·C, yielding a white PBSZ powder (Mn =2100-2200) that is soluble in toluene and chloroform but does not melt < 500·C [16]. The PBSZ composition isclose to SiBo ,380o ,3iN1.o5Co ,31Hu~ . The chemical bonding between PHPSZ and trlmethyl borate is mainly via B-N bonds formed according to the following reaction:

t ~-NHt- + H

/I

(CHP)JB -

t~-N+I

H

+CHpH

(7)

/I

B(OCH J)2

SI-O bonds can also be formed as the result of a reaction of Si-H bonds in PHPSZ with CH30H, orwith the trimethyl borate itself:

tf-NHto I

+CH4

(8)

/I

B(OCH)2

PBSZ is a polymer consisting mainly of the structural units HSiN3, H2SIN2, H3SIN, HSION2, NB(OCH3h and NB(OCH3)N. PBSZ fibers are dry spun at 60·C from a filtered solution. Curing is not necessary since the material isinfusible after vaporization ofthe solvent. The green fibers can be pyrolyzed so as to either retain or to remove carbon. Nearly-carbon-free Si-B-O-N fibers are obtained by a two step pyrolysis. The green fiber is first pyrolyzedat 800·C in ammonia, yielding an amorphous fiber with a composition close to SiBo,.aOo,~NwCo ,olHo,ae that is nearly hydrogenfree. This fiber is then annealed in nitrogen at1500·C, yielding a composition that isclose to SIBo,.aOO,304Nl ,~CO ,OlHo ,~. If the pyrolysis is directly performed in nitrogen at 1500·C, very little carbon is removed, and the composition ofthe fiber isclose to SIBo,380o,38N1.1ICo,23Ho," [16]. Nearly oxygen-free PBSZs can be prepared by the reaction of PHPSZ with tris (dlmethylamino) borane in pyridine, in the presence of ammonia, by a procedure similar to that used for the synthesis of PBSZs from PHPSZ and trimethyl borate. The product is a clear solution of PBSZ in o-xylene or a white powder insoluble in organic solvents and infusible up to 500·C when the solvent is removed by distillation [31]. This PBSZ isa polymer consisting of HSIN3, H2SIN2, H3SiN and substituted borazine. Its empirical formula is close to SiBo,aeN2,38C1.380ooaH5.25. The pyrolysis of this PBSZ in a gas flow yields a residue which remains amorphous up to Tp i'= 1700·C and whose composition Is close to SiBo,aaNl,IaCo,330o,05HMI.

Chapter 11

311

(b) From trichlorosilylamino-dichloroborane (TADB)fibers

PBSZs, precursors to Si·B·N or SI-B-N-C oxygen-free fibers, can also be prepared by reacting trichlorosllylamino-dichloroborane (TAOB) with CI3Si-NH-BCb and methylamine, H3C-NH2 [17] [32]. The advantage ofTAOB lies in the fact that the Si-N·B bonds are already present in the main starting material. The low temperature reaction yields hexane and yields a soluble, thermosetting oligomer that can be further polymerized by heating at 160·C in a vacuum. Polymerization gives a yellow, glass·like PBSZ polymer (Mn =2500) that is fusible and soluble in organic solvents, but is sensitive to moisture and oxygen. Its composition, SiBN432C3le0007Hl1.11, corresponds to the formula SbB2(NHMNHCH3MNCH3) and its structure consists ofslx·membered Si3(NCH3)3 rings connected via HN-B- and N(CH3)·B- groups. Fibers are obtained from the PBSZ precursor either by melt spinning at 140·180·C or by dry spinning. No specific curing step is necessary. Pyrolysis at 1500·C yields amorphous fibers free ofoxygen. When the pyrolysis is performed in nitrogen or argon, the fibers are black and smooth, with a composition close to SIBN3C, and contain a significant amount of carbon. Conversely, when the pyrolysis is performed up to 1000·C in ammonia, the majority of the organic functionality Is removed at 800·C and the fibers are colorless; their composition is close to SbB3N7 [17] [32]. 11.4.2 Structures and properties Si-B·O-N, Si·B-N and Si-B-C-N fibers are amorphous; the thermal stability oftheir amorphous state is much greater than that observed for Si·N, Si·N·O and Si·C·N·O fibers. Thus, boron, alone or associated with oxygen, Impedes the crystallization ofthe pyrolytic residue In SbN4.

(a) Structure and thermal stability Amorphous SI-B-O-N ceramics are known to consist oftetrahedral SiN4.•O. units (with x =0, 1, 2, or 3) and trigonal BN3-yOy units (with y =0 or 1), where each unit is crosslinked through a three-dimensional network of Si·N·Si, Si-N·B, SI-O-Si, Si-O-B, B-N-B and B-O-B bonds [30]. SI·B-O-N fibers are mainly amorphous when pyrolyzed at 1700·C, but they crystallize at 1800·C In nitrogen (Figure 5). Crystallization, which is associated with a major (23%) weight loss and a decrease in the oxygen content, yields a mixture of c- and P-SbN4 but no crystalline h·BN. Its composition is close to Si3BN4. Oxygen-free Si-B-C·N ceramics resulting from the pyrolysis of polyborosilazanes behave similarly when annealed at high temperatures in nitrogen. When pyrolyzed at 1700·C, the pyrolytic residue ofPBSZ that resulted from the reaction ofPHPSZ with tris (dimethylamino) borane is amorphous, with each boron atom surrounded by three nitrogen atoms. Its composition is close to SbB3N3C. Crystallization, which occurs at about 1800·C is accompanied by weight loss; ityields amixture ofc- and p·SbN4 and P-SiC, but no crystalline h·BN [31]. The pyrolytlc residue ofPBSZ resulting from the reaction ofTAOB with methylamine, having a composition close to SiBN3C, is thermally more stable when annealed in nitrogen. Crystallization above 1900·C yields amixture of c- and p·SbN4 but no crystalline h-BN or B4C phase. At 1800·C, when annealing is performed in argon, crystallization yields Si3N4 and SIC, and is accompanied by a7wI. %loss ofnitrogen [17].

312

Chapter 11

1800° C

10

20

30

40

50

60

29, degrees

Figure 5. XRD patterns of Si-B-O-N fibers annealed in an atmosphere of nitrogen at increasing temperatures [16): reproduced with permission.

(b) Mechanical properties

The room temperature tensile strength of annealed Si-B-O-N fibers, measured at room temperature, is much higher than that of the Si-N fibers prepared from the same PHPSZ precursor (Figure 6). The tensile strength of SoN fibers decreases above an annealing temperature of 1300°C, the onset of crystallization, and the fibers are too brittle to be tested after pyrolysis at 1600°C. The tensile strength of the Si-B-O-N fibers is still >2 GPa after annealing at 1600°C, and the fibers are too brittle tobe tested after heat annealing at1800°C. The difference of about 300°C where the fibers exhibit the same residual tensile strength correlates with the increase in thermal stability of the amorphous state related to the addition ofboron [16].

11.5 Applications The fibers described in this chapter are still experimental materials. Si-N or Si-N-O fibers might find use in applications where good mechanical properties, good oxidation resistance and low electrical conductivity are required. Si-B-O-N or Si-B-C-N fibers, displaying high

Chapter 11

313

thermal stability, good mechanical properties at high temperatures and oxidation resistance, are potentially useful for reinforcing ceramic matrices. Chapter 12 deals with applications for carbon, ceramic oxide and silicon carbide fibers.

3

-o----'

oL--.l_ _....1..._ _..L-_--'I..lir_--'-_ _ 1700 1800 1600 1400 1500 1300

Temperature, · C

Figure 6. Tensile strength measured at room temperature after annealing at high temperatures for Si-N fibers derived from PHPSZ (ti) and Si-B-Q-N fibers derived from PBSZ (0). The fibers were first pyrolyzed under flowing ammonia at 1200·C and then annealed under flowing nitrogen at 1300-1800·C, except (e) Si-B-O-N fiber, derived from PBSZ, which was simply pyrolyzed to 1500·C under flowing nitrogen according to ref. (16); reproduced with permission.

REFERENCES (1) (2) [3] (4) (5) [6J

[7J

K. Okamura, Ceramic fibers from polymer precursors, Composites, 18[2], 107-120 (1987). J. lipowitz,J.A. Rabe and R. M. Salinger, Ceramic fibers derived from organosilicon polymers, in Handbook of Fiber Science and Technology: Vol. 11/. High Technology Fibers, Part C, M. Lewinand J. Preston, eds., Marcel Dekker, New York, 207-270 (1993). W. H. Atwell, Polymeric routes tosilicon carbide and silicon nitride fibers, in Advances inChemistry Series 224: Silicon-Based Polymer Science, A Comprehensive Resource, J. M. Zeigler and F. W. Gordon, M. Fearon, eds., The American Chern. Soc.,Wash., 593 (1990). T. Isoda, Surface of high purity silicon nitride fiber made from perhydropolysilazane, in Controlled Interphases ;n CompositeMaterials, H.lshida, ed., Elsevier Science Publishing, Amsterdam, 255-265 (1990). O. Funayama, M. Arai, Y. Tashiro, H. Aoki, T. Suzuki, K. Tamura, H. Kaya, H. Nishii and T. Isoda, Tensile strength of silicon nitride fibers produced from perhydropolysilazane, Nippon Seramikkusu Kyokai Gakujutsu Ronbushi, 98(1), 104-107 (1990). H. Matsuo, O. Funayama, T. Kato, H. Kaya and T. Isoda, Crystallization behavior of high purity amorphous silicon nitride fiber, J. Ceram. Soc. Japan, 102 [5], 409-413 (1994). K. Okamura, M. Sato and Y. Hasegawa, Silicon nitride fibers and silicon oxynitride fibers obtained by the nitridation ofpolycarbosilane, Ceramics Intern., 13,55-61 (1987).

314 (8) (9)

[10]

[llJ [12J [13J [14] (15) (16) [17J [18] [19] [20] [21] [22] [23] [24] [25] [26] (27) [28J [29J [30J [31] [32]

Chapter 11 D. Mocaer, R. Palller, R. Naslaln, C. Richard, J. P. Plllot, J. Dunogues, C. Gerardln and F. Taulelle, SI·C·N ceramics with a high microstructural stability elaborated from the pyrolysis of new polycerbosllazane precursors. Part I.The organic-Inorganic transition, J. Mater. ScI., 28, 2615·31 (1993). D. Mocaer, R. Palller, R. Naslaln, C. Richard, J.P. Plllot, J.Dunogues, O. Delverdler and M. Monthloux, SI-C·N ceramics with a high micro-structural stability elaborated from the pyrolysis of new polycarbosllazane precursors. Part II. Effect ofoxygen curing on properties ofex-PCSZ monofilaments, J. Mater. Sci., 28, 2632· 38 (1993). D. Mocaer, R. Palller, R. Naslaln, C. Richard, J. P. Plllot, J. Dunogues, O. Delverdler and M. Monthloux, SI·C·N ceramics with a high micro-structural stability elaborated from the pyrolysis of new polycarbosllazane precursors. Part III. Effect ofpyrolysis conditions on the natura and properties ofoxygen-curad derived monofilaments, J. Mater. Sci., 28, 2639-2653 (1993). D. Moceer, R. Palller, R. Naslaln, C. Richard, J. P. Plllot, J. Dunogues, C. Damez, M. Chambon and M. Lahaye, SI·C·N ceramics with a high microstructural stability elaborated from the pyrolysis of new polycerbosllazane precursors. Part IV. Oxygen·frae model monofilaments, J.Mater. ScI., 28, 3049-58 (1993). D. Moceer, G. Chollon, R. Palller, L. Flllpuzzi and R. Neslaln, SI-C-N ceramics with a high mlcrostructurel stability elaborated from the pyrolysis of new polycarbosllazane pracursors. Part V. Oxidation kinetics of model filaments, J.Mater. ScI., 28, 3059-68 (1993). T. Ishikawa, Recent developments ofthe SIC fiber Nlcalon and Its composites, Including properties ofthe SIC fiber HI·Nlcelon for Ultra-high temperature, Composites ScI. and Technology, 51,135·144 (1994). S. R. Rlceltlello, M. S. Hsu and T. S. Chen, Ceramic fibers from SI-B·C polymer pracursors, Sampe Quarterly, April Issue, 9·14 (1993). O. Funayama, T. Kato, Y. Tashiro and T. Isoda, Synthesis of a polyborosllazane and Its conversion Into Inorganic compounds, J. Amer. Ceram. Soc.,76 (3), 717·723 (1993). O. Funayama, H. Nakahara, A. Tezuka, T. IshII and T. Isoda, Development of SI-B-O·N fibras from polyborosllazane, J. Mater. Sci., 29, 2238-44 (1994). H. P. Baldus, G. Passing, D. Spom and A. Therauf, SI·B(N, C) a new ceramic material for high performance applications, Proc. In!. Conf. High Temperature Ceramic Matrix Composites, HT-CMC-2, A. G. Evans and R. Naslaln, eds.,Ceram.Trans., 58, 75·84, The Amer. Ceram. Soc" Westerville, USA (1995). M. Blrot, J. P. Plllot and J. Dunogues, Comprehensive chemistry of polycerbosllanes, polysllazanes and polycerbosllazanes, as precursors ofceramics, Chem. Reviews, 95, 1443-77 (1995). W. Verbeck, German Patent 2,218,960 (1973). B. G. Penn, J.G. Daniels, F. E. Ledbetter and J.M. Clemons, Preparation ofsilicon cerbide-slllcon nitride fibers by the pyrolysis ofpolycerbosllazane precursors: areview, Polymer Eng. ScI., 26 [17], 1191·94 (1986). G. E. Legrow, T. F. Lim, J. Llpowltz and R. S. Reaoch, Ceramics from hydrodopolysllazane, Amer. Ceram. Soc. Bull., 66 12J, 363-367 (1987). R. Bodet, N.Jla and R. E.Tressler, Microstructural Instability and the resultant strength ofSI-C-O (Nlcelon) and SI·N·C·O (HPZ) fibres, J. Europ.Ceram. Soc., 16,653-664 (1996). Y. C. Song, Y. Zhao, X. Feng and Y. Lu, Synthesis and pyrolysis ofpolysllazane as the precursor ofShNJSIC ceramic, J.Mater. ScI., 29, 5745·56 (1994). O. Delverdler, M. Monthloux, D. Moceer and R. Palller, Thermal behavior ofpo1¥mer-derlved ceramics. IV: SIC·N·O fibers from oxygen-curad polycerbosllazane, J. Europ. Ceram. Soc.,14, 313-325 (1994). R. E. Tressler and D. J.Pysher, Mechanlcel behavior ofhigh strangth ceramic fibers athigh temperatures, In Advanced Structurellnorganlc Composltas, P. Vlcenzlnl, ed., Elsevier, Amsterdam, 3·18 (1991). G. Chollon, Oxidation behavior ofceramic fibers from the SI·C·N·O system and SUbsystems, Key Engineering Materials, 164/165, 395-398 (1989). K. Okamura, M. Sato and Y. Hasegawa, SI-N-O fiber and SI·TI·C fiber obtained from polycarbosllane, In Proc. 5th In!. Conf. Composite Matarlals, W. C. Harrigan, J. Strife and A. K. Dhlngra, eds., The Metallurgical Society, Warrendale, PA, 535-544 (1985). S. Kamlmura, K. Watanabe, N. Kasal, T. Seguchl and K. Okamura, Silicon nitride fiber synthesis from polycerbosllane fiber by radiation curing and pyrolysis under ammonia, Ceram. Transactions, 58, 281·286 (1995). T. Takl, M. Inul, K. Okamura and M. Sato, A stUdy ofthe nitridation process ofpolycarbosllane by solld·state high resolution NMR, J.Mater. ScI. Letters, 8, 1119·24 (1989). O. Funayama, Y. Tashiro, A. Kamo, K. Okamura and T. Isoda, Conversion mechanism ofperhydropolysllazane Into silicon nltride·based ceramics, J. Mater. Sci., 29, 4883·88 (1994). O. Funayama, T. Aokl and T. Isoda, Synthesis and pyrolysis ofpolyborosllazane with low oxygen content, J. Ceram. Soc. Japan, 104 [5J, 355·360 (1996). H. P. Baldus, O. Wagner, and M. Jansen, Synthesis ofadvanced ceramics In the systems SI·B·N and SI·B·N·C employing novel precursor compounds, Mater. Res. Soc. Proc., Vol. 271, The Materials Research Society, Pittsburgh, PA, 821·826 (1992).

CHAPTER 12 APPLICATIONS OF CARBON AND CERAMIC FIBERS R. Naslain This chapter is an overview of the applications for carbon and ceramic fibers. These applications mainly involve composite materials where the fibers are embedded in a polymer or a metallic or ceramic matrix, depending on the service temperature or the specific properties which are required. 12.1 Fiber applications Carbon and ceramic fibers are usually embedded in a matrix. There are applications in the fields ofthermallnllulatlon, catalyst support, filters, porous electrodes, and battery separators where they are utilized as mats or felts, or with asmall amount ofabinder. Thermal insulation isby far the most important present application ofoxide fibers. Transition alumina fibers, e.g., eta-alumina fibers, are produced atIntermediate firing temperatures and are used as supports for catalysts and Insulation tiles such as those used for the space shuttle orbiter [1-2]. Carbon fiber felts are used as internal thermal insulation for vacuum furnaces at extremely high temperatures. Activated carbon fibers, which are obtained by partial oxidation ofselected carbon fibers, have extremely small pores and very high specific surface areas, ranging from 500 to 3000 m2/g. They are of great Interest In ultrafiltration as membranes for the treatment ofused waters and liquids [3-5]. Rechargeable lithium cells are being developed to overcome difficulties observed in charging the lithium electrode orlithium-Ion battery. These problems are caused by reaction between Li and electrolytic solutions, and by cell shorting due to Li dendrite growth. One way to solve these problems is to use a negative graphite-lithium electrode, and to intercalate Li into and de-intercalate LI from the carbon host material: with 0 s x S J

(1)

Carbon fibers derived from PAN or mesophase pitch appear to be promising electrode materials [6-9]. Lithium intercalation Is highly reversible and both the current densities and the cycle lives are reported to be compatible with power batteries and electric vehicle applications [8]. Zirconia fiber cloths, such as those made by the relic process (see Chapter 8), are used as separators in aerospace nickel-hydrogen and nickel-cadmium batteries [9]. These fibers display a high resistance to many corrosive media, Including hot potassium hydroxide. Zirconia fiber felts are also present in aerospace solid oxide fuel cells.

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12.2 Composite applications Polymer matrix composites (PMCs), also referred to as fiber reinforced plastics (FRPs), consist of strong and stiff fibers reinforcing, unidirectionally (10) or multidirectionally (nO), thermosetting orthermoplastic polymers. They have high specific strength (crR/p) or/and high specific stiffness (E /p), as well as high fatigue resistance; crR, E and p are tensile strength, Young's modulus and density. Thermoset polymers, including polyesters, vinylesters and epoxies, are resins which irreversibly crosslink during curing. Thermoplastic polymers are fusible, high molecular weight, long chain molecules, which melt upon heating and can give, in a reversible manner, either an amorphous (polycarbonates) or semicrystalline (nylon) solid when cooled toroom temperature. 12.2.1

Polymer matrixcomposites

Three reinforcements are mainly used in PMCs: glass fibers, aramid fibers and carbon fibers [11 -13J. Carbon fibers have the advantage of being simultaneously strong and stiff (3 ~ crR s 7 GPa; 250 ~ E s 950 GPa and 1.8 s p s 2.1 g/cm3) (as discussed in Chapter 9). Glass and aramid fibers are strong but lack stiffness (E = 75 and 130 GPa, respectively). Strong carbon fibers (E = 250 GPa; crR= 3 GPa and p = 1.8 g/cm 3) yield 1-0 epoxy composites with the stiffness ofsteel (200 GPa), ahigher strength and alower density. PMCs have a much higher fatigue resistance than metals. However, the cost ofcarbon fibers ishigh, and carbon fibers are therefore often used as a hybrid reinforcement with glass fibers to increase composite stiffness and take advantage of hybrid economics. PMCs are practically used from cryogenic temperatures to a maximum temperature of about 300°C. Silicon carbide (or nitride) fibers as well as oxide fibers with low electrical conductivities are used in PMCs for specific applications where good dielectric properties are required, e.g., in electronic devices oraerospace structures with alow radar signature. The fabrication of PMCs iswell-documented [14J. A prepreg route is often preferred in order to facilitate the handling of the plies and to avoid the evolution of large amounts of organic solvents. A prepreg is an intermediate product consisting of a layer of impregnated fibers in which the polymer matrix is only partially crosslinked and hence is still suitably soft, the solvent being partially removed. The full crosslinking of the matrix is achieved, after plies stacking, in an autoclave. Another frequently used technique isfilament (or tape) winding, in which an impregnated fiber tow (or tape) is wound on a mandrel, and the mandrel is removed after autoclave curing. This technique is used to fabricate axisymmetrical parts and complex shapes. In the pultrusion technique, impregnated fibers are pulled continuously through a heated die, yielding a 1-0 composite with a constant cross section and a high fiber volume fraction. More specific techniques, such as hot pressing, structural reaction injection molding (S-RIM), resin transfer molding (RTM) and others, have been designed mainly toreduce processing time [141. Space and aerospace were the first important fields for applications of composite materials, particularly for carbon fiber reinforced polymer matrix composites [12-13J [15J. Carbon/epoxies are main constituents of the rudder, elevators, stabilizers, outboard and inboard ailerons, flaps, landing gear doors, engine cowlings and even wing skins of advanced civilian aircraft. High performance PMCs reduce weight and improve performance and

317

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maneuverability of fighter aircraft. Composites make up 1% and 2% by weight of the F-15 and the F-16, while carbon/epoxies will make up about 50% of the outer surface of the F/A18, where the composites represent 10% ofthe total weight ofthe fighter plane. In the next fighter plane. carbon/epoxy components will be about 26% of the total weight, resulting in a 20-30% weight reduction for the aircraft. The use of PMCs in the structures of the stealth bomber contributes to lower its radar signature. More and more PMCs are also used in helicopters in the airframe/fuselage, and even in the most critical parts, such as the rotor blades themselves. Composite rotor blades exhibit much better fatigue resistance and much higher service life than their metal counterparts [15-17].

_ _

=

Figure 1.

Carboncomposite Aluminum Aramidcomposite SPFDBtitanium

Advanced composite materials inaircraft structures: Rafale fighter aircraft; courtesy Dassaut Aviation.

Carbon/epoxies are also extensively used for space applications where stiffness, dimensional stability and weight are key considerations. The shuttle orbiter 18 meter cargo bay doors are among the largest carbon/ epoxy structures ever built. The 15 meter long mechanical arm of the space shuttle orbiter can handle space payloads weighing 20 tons on earth. Most satellites and spacecraft have large antennas made of carbon/epoxy composites to achieve dimensional stability. Carbon fibers have a negative coefficient of thermal expansion (CTE) and epoxies yield composites which have no thermal expansion or contraction within the temperature range corresponding to the space environment. As a result, both the shape ofthe antenna surface and its precise alignment are properly maintained. Further, the space stations of the future will be built from structural beams fabricated from carbon fiber/polymer matrix composites, to obtain the required stiffness, and assembled in

318

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space. Finally, PMCs are used in strategic missiles (rocket motor cases), cruise missiles (tail fins and ailerons), and insatellite launchers [15] [18] [20]. The weight of automobiles could be significantly reduced and their fuel economy improved if most of the steel components could be replaced with much lighter PMCs. However, cost considerations so far have limited the use of high performance PMCs in automobiles. Metal replacement inautomobiles continues tobe dominated by glass fiber reinforced PMCs. Carbon fiber/polymer matrix composites are extensively used in sporting goods and recreational applications. This is the second most important field of applications for carbon fibers. It has been largely responsible for the rapid growth of carbon fiber production during the last two decades and hence for the decrease in fiber prices [12-13]. The properties justifying the use of PMCs in this field are the high specific strength and stiffness of the composites, coupled with rapid vibrational damping and fatigue resistance. Golf club shafts, tennis rackets, fishing rods, cross country skis and poles, alpine skis, and bows and arrows are well known examples of sporting goods fabricated with carbon fibers [21]. PMCs are also finding more and more use in marine applications because of their strength and stiffness as well as their corrosion resistance. Typical examples ofhighly loaded components are the masts, booms, and rigging of racing sail boats, the cross beams of oceanic catamarans and trimarans, as well as in recently designed hydrofoil sail boats (Figure 2).

Figure 2.Aquitaine Innovations, a60-foot oceanicracing sloop designed bythe Fino-Conq architecfs group, islargely buiK with carbon fiber reinforced composites, including hull, deck, mast, centerboard, rudder rods, and blades. The hull contains by itsen 2.5metric tons ofcarbon fibers and was fabricated according totechnology previously used for aerospace structures by Composites Aquitaine an Aerospatial

Sailplanes and modern racing bicycles [21-22] represent other applications for PMCs. Here, the high specific stiffness ofcarbon fiber/epoxy composites isa key factor in structural design and performance improvement. Carbon fiber PMCs have found a few niches in industry on

Chapter 12

319

the basis of their specific stiffness, low moment of inertia, dimension stability or fatigue and corrosion resistance. The high cost of materials is compensated by the improvement of component performance. The use of carbon fiber PMCs has enabled robotics engineers to design manipulators which display low inertia and small end effect vibrations, whereas high acceleration and precision were previously considered to be mutually exclusive [23-24]. Pressure vessels offer key opportunities for properly designed MMCs [12). They are filament wound over a permanent stainless steel or ductile aluminum alloy liner that acts as a mandrel and also prevents potential gas or fluid leakage under pressure. Such light tanks are used in emergency situations by firemen ordoctors (Figure 3). e Group subsidiary; courtesy Henri TibaultlDPPI.

Figure 3. Pressure vessels with a capacityof 3 to 9 liters and an internal pressure of 30 MPa, fabricated with carbon fiber/epoxy composites, are used byfiremen in extreme situations. The pressure vessels provide firemen an autonomy upto one hour and correspond to 50% weight saving; courtesy Composites Aquitaine. an Aerospatiale Group subsidiary.

High performance carbon fiber PMCs are also of interest in energy storage or energy conversion involVing components oflarge size (windmill generators) orspinning athigh speed (flywheels). Stored kinetic energy is proportional to the moment of inertia and tothe square of the spinning speed. For this application, lightweight carbon fiber reinforced composites rely on high rotational speed in preference to large mass [12). Carbon fiber PMCs are also used to fabricate large blades for wind turbine generators. Carbon fiber PMCs, consisting of light elements (H, C, 0, N), display high X-ray transparency. This property combined with high specific stiffness has been exploited in X-ray equipment for non-destructive testing and medical analysis. Rigid and lightweight tables for computerized

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X-ray tomography are fabricated by applying skins ofcarbon fiber polymer laminates toacore made offoam orhoneycomb (27). The use of SiC or alumina based ceramic fibers in organic polymer matrix composites remains limited. Their production cost ishigher and their performance isoften lower than that for carbon fibers. The exception has to do with their dielectric properties. The dielectric constant of epoxy composites reinforced with Nicalon fiber can be widely varied by merely changing the nature of the fiber itself. When high resistivity NL-400 grade fiber is used (p = 106 - W n cm), the dielectric constant of the composite is high (20-30 at 10 GHz), and with low resistivity N500 grade fiber (p = 0.5 - 5.0 n cm), the dielectric constant is low or zero. Composites with low orhigh microwave absorption are used in advanced military aircraft (28). 12.2.2 Metal matrix composites With a few exceptions, metal matrix composites (MMCs) are still in a development stage although it is well known that the presence of inorganic fibers in a metal matrix has a major effect on the properties of the metal. The reinforcement and manufacturing costs are inevitably too high for the potential markets, such as that ofautomotive engines [29-33). In an aluminum alloy, for example, the presence of inorganic fibers: (1) increases its stiffness, failure strength and fatigue resistance at room and high temperatures; (2) improves its wear resistance if the reinforcement is a hard material, such as alumina or silicon carbide; and (3) increases its dimensional stability in thermal cycling, the CTE ofthe fibers usually being lower than that of the matrix. In UHM carbon fiber/AI (or Mg) composites, the CTE of the composites can even be nil for a specific fiber volume fraction, the contraction of the fibers and the dilatation ofthe matrix compensating one another when the temperature israised. Most MMCs are systems out of thermodynamic equilibrium. Diffusion phenomena and chemical reactions occur at the fiber/matrix interfaces which degrade the mechanical properties when such systems are heated for short periods of time at high process temperatures or maintained for prolonged periods of time at intermediate service temperatures. This is the case with C/ftJ, SiC/AI and SiCITi composites [34-38]. Hence, the main advantage of MMCs, their potential for high temperature applications, is altered unless specific alloys and/or fiber coatings are used. For example, Al-Si alloys are less reactive toward SiC fibers than pure AI, and SiC/C filaments are usually coated with multilayered carbon coatings when used to reinforce titanium alloys. Under such conditions, reactions occur between coating and matrix and the fibers remain undamaged as long as there issome carbon in between acting as amechanical fuse [34-35)[39). Compared to PMCs, the fabrication of MMC components is more difficult owing to the high melting points of metal matrices (660°C for aluminum; 1083°C for copper, and 1668°C for titanium) and their high chemical reactivity for most carbon and ceramic fibers. MMCs can be fabricated bydifferent techniques (12)[40-45). In the field of MMCs, continuous carbon and ceramic fibers compete with other reinforcements on a cost/performance basis. Continuous, and hence costly, fibers can be used in applications where high specific performance, e.g., in space applications, is required. In other potential applications, e.g., in automotive engines, metal matrix composites are not cost effective.

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One of the oldest applications of MMCs in space is the use of S/AI tubes, serving as a frame for the mid-fuselage section ofspace shuttle orbiters. The key requirements are high specific stiffness, high specific strength, and high fatigue resistance. The reinforcement is a continuous, 140 IJm diameter boron/tungsten bicomponent fiber (Chapter 4), which is unidirectionally embedded in an aluminum alloy. Each structure, consisting of more than two hundred tubes and ranging from 60 to 180 cm in length, results in a 44% weight saving over the original aluminum design, or 140 kg per vehicle. Carbon fiber/AI (or Mg) composites have been fabricated into beams for large antennas and benches for optical instruments in satellites. The main requirements are specific stiffness, thermal conductivity and dimensional stability. UHM carbon fibers derived from mesophase pitch (Chapter 9) are well suited for these applications, owing totheir extremely high stiffness, high thermal conductivity and negative CTE [45]. MMCs are considered promising for the design of specific parts for the cars of the future, which will be more fuel efficient, safer and have lower emissions. Although reducing the vehicle weight is also an important target, this industry is extremely competitive and price sensitive. MMCs have been envisioned for drive shafts, brake rotors, cylinder liners, connecting rods and pistons [46-47]. These parts experience severe service conditions and require the use of materials with good stiffness and strength at high temperatures, excellent fatigue and wear resistance, and compatibility with hot hydrocarbons, oil or combustion gas. A 250-300·C peak temperature is observed at the top of a piston. MMCs meet these requirements and alumina as well as silicon carbide fiber reinforced piston rods have performed well in racing cars. Although the feasibility offibrous MMCs istechnically demonstrated, the cost ofthe fibers and of the manufacturing processes, e.g., squeeze casting, is still too high. Thus, particulate reinforcements are presently preferred. Potential applications exist for metallic (MMCs) and intermetallic matrix composites (IMCs) in next generation aerospace propulsion systems for advanced commercial subsonic aircraft and high speed civil transport. These applications require stiff, strong and light materials that operate athigher temperatures with long lifetimes [48]. One candidate is a titanium alloy that is reinforced with large diameter SiC/C filaments (see Chapter 4) and isfabricated by superplasticforming/ diffusion bonding. This MMC is suited to the fabrication of bladed compressor rings, shafts, ducks, fan components or structural rods for jet engines. Their use for parts submitted to still higher temperatures is limited by fiber/matrix reaction and environmental considerations. Titanium aluminide ThAI (or y-TiAI) matrices could permit an increase in the service temperature ofthe composites. Substantial improvements, e.g., better fiber coatings, matrix alloying and oxidation protection, are needed before IMC technology can be seriously considered for jet engine applications. An important property of MMCs, compared to their PMC counterparts, is their electrical and thermal behavior; they have been used for some time in innovative applications such as in electrical and electronic devices.

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Both carbon fiberllead and alumina/lead composites have been utilized for the fabrication of electrodes in lead acid batteries with a significant weight reduction. For example, reinforcing lead unidirectionally with 25 vol.% of a-alumina fibers increases the matrix stiffness by a factor of20 and reduces the electrical conductivity only slightly. Due tothe stiffening effect of the fibers, pure lead can be used instead of a complex alloy, thus reducing the cost of the matrix and improving its corrosion resistance [49). Due to their tailorable combinations of properties, MMCs may play a key role in advanced electronic packaging, e.g., in power semiconductor diodes. In this application, the silicon wafer cannot be directly soldered to the copper electrode, due to a CTE mismatch (a (Si) = 4.1 x 10-6; a (Cu) = 17 X 10-6 °C·'). Under service conditions, interface delamination would occur as the result of thermal stresses arising from temperature changes. In conventional power diodes, the problem is partly solved byusing an intermediate molybdenum ortungsten layer with a CTE between those ofsilicon and copper (a (Mo) = 5.0x 10-6 and a (W) = 4.6 X 10-6 °C·1) . Another concept is the use a carbon fiber reinforced copper electrode. Carbon fibers have negative CTEs with a CTE matching that ofSi [50). Packaging materials support and protect integrated circuits in electronic devices. They also playa key role in heat removal. The amount of heat to be removed increases rapidly as the density of the integrated circuit packaging is raised. A key requirement is a high thermal conductivity coupled with a low CTE (semiconductor chips and ceramic carriers display low CTEs compared to metals). For some applications, such as aircraft avionics, weight is an important parameter. Among conventional packaging materials, such as copper, aluminum, molybdenum, tungsten, Invar, Kovar, W/Cu, or Mo/Cu blends and laminates, none exhibits simultaneously a low CTE and a high specific thermal conductivity. Two different MMC families have been studied for integrated circuit packaging, namely SiC/AI and UHM carbon fiber/Cu composites. P120 UHM carbon fibers are unsurpassed in terms of thermal conductivity. They have a 50% higher thermal conductivity than copper and a negative CTE), but they are expensive, whereas poorer performers are inexpensive [51 -52). 12.2.3 Carbon and ceramic matrix composites Ceramic matrix composites (CMCs), in which carbon or ceramic fibers are embedded in a ceramic matrix, have been designed to overcome the intrinsic brittleness of monolithic ceramics with a view toward structural uses at extremely high service temperatures. The most commonly used are carbon (C/C) and SiC matrix composites (C/SiC and SiC/SiC). Ceramic matrix composites with a silica based glass or glass-ceramic matrices have also been studied [12) [53-56). Carbon fiber reinforced carbon (C/C) composites, which are fabricated by the polymer impregnation/pyrolysis process (PIP) or by chemical vapor infiltration (CVI), are commonly used in rocket engine nozzles, where the temperatures can exceed 3000°C, and divergents, such as the very large booster nozzles of the Ariane V satellite launcher (Figure 4). C/C composites are also utilized in the heat shields of space vehicles and in the nose cap and leading edge of space shuttle orbiters. These are the structural components which experience the highest temperature (1300 to 1500°C) during reentry into the earth's atmosphere. However, C/C composites must be protected against oxidation since many are reusable.

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In the US space shuttle orbiters, the C/C outer surface is first converted to SiC; then the material is impregnated with tetraethylorthosilicate which yields, after curing, a silica residue in the porosity; finally, a seal coating is applied to the surface [571. In the European project of space shuttle orbiter Hermes, several parts, including the very large wing box components, spar bars, panels and leading edge segments, were directly fabricated with C/SiC composites by isothermal/isobaric chemical vapor infiltration (I-CVI). Tests were successfully performed on a leading edge segment under conditions simulating reentry up to1550-1700·C [58]. Nickel based superalloys which have been successful in the development of gas turbine engines over the last 50 years have reached their limit. To operate the engines at higher temperatures, new materials are much needed. CMCs are thought to have enough potential to satisfy such conditions (SiC has a melting point of 2500·C and a density of 3.2 g/cml However, they are still very new and their behavior in long term exposures (the lifetime of an engine is of the order of 10,000 hours) to corrosive atmospheres under load is still partly unknown [59-60]. Feasibility and test rig evaluation have been conducted on a flame holder and an exhaust cone (operating temperature: 800-1100·C and 800-950·C, respectively) fabricated from Nicalon/SiC composites by I-CVI. Weight saving was 50% and 30%, respectively. Nicalon/SiC composite inner flaps (operating temperature: 850·C, weight saving: 60%) were mounted on a SNECMA M53-2 engine and flight tested on a Dassault Mirage 2000 fighter (Figure 5). C/SiC outer flaps (operating temperature less than 650·C, weight saving: 50%) have been flight tested in SNECMA M88-2 engines of the Dassault Rafale fighter and are now produced in series [59].

Figure 4. Nozzle of the booster rockets of the European satellite launcher Ariane V having anouter diameter of 2.8 m, an inner diameter of 90em, and anoverall weight of 6 metric tons. The nozzle is fabricated with C/C and carbon fiber/phenolic matrix composites; courtesy SEP division deSNECMA.

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Figure 5. C/SiC outer flaps, in-flight tested on a SNECMA-M53.2jetengine and a Dassault Mirage 2000 flQhter aircraft. The flaps were fabricated byCVI and represent a 50% weight savings; courtesy SEP division deSNECMA.

Recently, the lifetime of SiCISiC composites has been increased through the use of SiC based fibers, with much higher thermal stability such as the Hi-Nicalon fibers, and self-healing multilayered interphases and matrices. Hence, it is anticipated that most of the hot static components of gas turbine engines, including the combustor liner (operating temperature: 1400°Cand above) would be fabricated with SiC matrix composites [61]. Over one-half of the CIC composites in the world are used in aircraft brake systems (Figure 6). The use of CIC braking in racecars is cost effective, but in other land vehicles such as high speed trains it isnot [54].

Figure 6. C/C discs for the braking systems of civilian Airbus aircraft (left) and military Mirage 2000 fighter aircraft (right). The discs were fabricated according tothe I·CVI process; courtesy SEP division deSNECMA.

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Aircraft braking systems have a multiple disc rotorlstator configuration, the rotor discs being sandwiched between the stator discs, and the rotor driven by the wheel. The braking action is performed by pressing the discs together. The brake discs have two main functions: they provide the frictional torque which stops the vehicle and serve as heat sinks to absorb the large amount of heat generated during the braking action. Replacing steel with C/C composite in an aircraft braking system results in a 40% weight saving and a doubling of the service life [54] [62-63].

C/C braking systems are already extensively used in racing cars (such as Formula 1) with, however, a different configuration. Here, the braking action is performed by pressing pads against a ventilated disc. C/C composites fabricated by either the PIP process (impregnation of the fibers with e.g. a phenolic resin and pyrolysis) or the CVI process are actually used. Finally, as already mentioned, tests have been conducted in order to explore the potential of C/C braking systems in high speed trains [54]. C/C composites may be used in place of aluminum alloys in spark ignition (or diesel) engine applications, e.g. as pistons and valves, on the basis of their low CTE, high thermal conductivity, good tribological properties, and low density. The benefits ofusing C/C parts are reduced reciprocating weight (and hence reduced vibration) and higher operating temperatures. Further, the combined use of C/C pistons and C/C cylinder liners may eventually eliminate the need for piston rings, owing to the tribological properties of C/C composites. However, such applications suppose low raw material and manufacturing costs [64]. CMCs are also being evaluated for applications in a 100 kW ceramic gas turbine for automobiles, with a turbine inlet temperature of1350·C [65]. Heat management systems such as combustion systems, burners, heat exchangers, waste incinerators and steam generators are potential applications for CMCs. These systems produce or recover heat under severe temperatures and exhibit a higher resistance to corrosion . For example, metal heat exchangers that extract heat from combustion gases cannot be operated above 750·C and their durability is limited by corrosion. SiC/SiC composites could be operated at1000·C, increasing the efficiency of, e.g., a coal fired power generation plant [56].

C/C composites find increasing use in industry both at medium and high temperatures in place of asbestos or monolithic graphite. Although much more expensive than the materials they replace, C/C composites display higher mechanical properties, shock resistance and durability. C/C composites are also used as heating elements for furnaces operating at high temperatures with non-oxidizing atmospheres, and as charging stages or trays for carbonizing, carbiding, orheat treatment furnaces, inplace ofmonolithic graphite orrefractory metals (Figure 7). Their use increases part lifetime and facilitates the handling of the parts [53-54]. C/C composites are used for hot-pressing dies, molds for superplastic forming of Ti alloys at 900·C, and for components of oil-less compressors, pumps and devices used in materials processing.

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Figure 7.

C/C tray for heat treabnent furnaces; courtesy SEP division deSNECMA.

Figure 8.

C/C hipprosthetic device; courtesy SEP division deSNECMA.

Finally, carbon is known to have superior biocompatibility with bones, blood and soft tissues. As a result, carbon fiber tows or GIG composites are used as tendon replacements or as tooth implants. GIG composites with some residual porosity are used to manufacture bone prosthetic devices. Their structure can be tailored to match the mechanical properties of bones and their biocompatibility allows the in-growth ofnew tissue [13) [53-54).

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ACRONYMS 1D 3D ADFA BEl BMU BSU CAD CMC CR CSZ CTE CVD CVI DCCA DMAC DMF DRIFT DSC DTA ECR EDX EFG EMI EOPA EPMA

One-dimensional; unidirectional Three-dimensional Erbium doped fiber amplifier Backscattered electron image Basic microstructural unit Basic structural unit Computer aided design Ceramic matrix composite Corrosion resistant Cubic stabilized zirconia Coefficient of thermal expansion Chemical vapor deposition Chemical vapor infiltration Drying control chemical additive Dimethylacetamide Dimethylformamide Diffuse reflectance infrared spectroscopy Differential scanning calorimetry Differential thermal analysis High acid corrosion resistance Energy dispersive x-ray spectroscopy Edge-defined film fed growth Electromagnetic interference Ethyl 3-oxobutanodiisopropoxyaluminum Electron probe microanalysis ESCA Electron spectroscopy for chemical analysis FOG-M Military fiber optical glass HM High modulus HP High pressure HPZ Hydridopolysilazane HRTEM High resolution transmission electron microscopy HS High strength HT High temperature HTT Heat treatment temperature IC Integrated circuit 1M Intermediate modulus IMC Intermetallic matrix composite IMS Inviscid melt spinning ITA Itaconic acid LA Laurylamine LCVD Laser-assisted chemical vapor deposition LHFZ Laser heated float zone LHPG Laser heated pedestal growth L1GA X-ray lithography

332

LM Low modulus LMO Local molecular orientation LP Low pressure MA Methylacrylate MAS 2,4,S-tris(methylamino)borazine MCVD Modified chemical vapor deposition MMC Metal matrix composite MP Mesopitch MTS Methyltrichlorosilane MWNT Multi-walled carbon nanotube NMR Nuclear magnetic resonance OVD Outside vapor deposition PAN Polyacrylonitrile PSDPSO Polyborodiphenylsiloxane PSG Photonic band gap PSSZ Polyborosilazane PCS Polycarbosilane PCSZ Polycarbosilazane PCVD Plasma chemical vapor deposition PDMS Polydimethylsilane PEEK Polyetheretherketone PHPSZ Perhydropolysilazane PIP Polymer impregnation/pyrolysis PMS Polymethylsilane PSP Polysilapropylene PSSZ Polysilasilazane PSZ Partially stabilized zirconia PSZ Polysilazane RS Rapid solidification RT Room temperature SIC Sheath/core SSA Tri-sec-butoxyaluminum s-b-s Side-by-side SEI Secondary electron image SEM Scanning electron microscopy SLS SOlution-liquid-solid SM Standard modulus SNMS Sputtered neutral mass spectrometry STM Scanning tunneling microscopy TADS Trichlolosilylamino-dichloroborane TEM Transmission electron microscopy TEOS Tetraethoxysilane Si(OC2Hs)4 TGA Thermogravimetric analysis TSZM Traveling solvent zone melting TZP Tetragonal zirconiapolycrystals UHM Ultrahigh modulus VAD Vertical axial deposition VLS Vapor-liquid-solid

Appendix

333

Appendix

VS XPS

XRD YAG

Vapor-solid X-ray photoelectron spectroscopy X-ray diffraction Yttrium aluminum gam

GLOSSARY Acoustical wave velocity: The distance traversed bya periodic, orcyclic; acoustic signal per unit time. Ammonolysis: A process that involves heating a chloro compound with aqueous ammonia toform an amine. Amphibole fibers: Fibers of a class of variously colored hydrous silicates, consisting chiefly ofcalcium, magnesium, iron, aluminum and sodium. Annealing: Treatment of a material byheating to a predetermined temperature, holding for a certain time, and cooling to room temperature to improve physical properties, such as strength, orchemical resistance. Bandwidth: The number of discrete channels within a finite range of frequency, which defines the data transfer rate ofan electronic communications system. Basalt fibers: Fibers from igneous (volcenic] rock that is low in silica content, dark in color, and comparatively rich iniron and magnesium. Basic structural unit: A stack of2-4 graphene layers lying roughly parallel toeach other. Bicomponent fibers: Fibers consisting of two components. Commercially significant inorganic bicomponent fibers have either a concentric sheath/core ora side-by-side fiber structure. Bioactive glass: A glass material which bonds to, and facilitates growth ofliving tissue. Birefringence: An optical property in which a single ray of unpolarized light entering an anisotropic medium splits into two rays, each traveling in a different direction. One ray (called the extraordinary ray) is bent, or refracted, at an angle as it travels through the medium; the other ray (called the ordinary ray) passes through the medium unchanged. Calcining: Heating a material, such as an inorganic material, to a high temperature but without fusion in order to drive offvolatile matter or to effect changes, such as oxidation orpulverization. Carbonization: The pyrolyzation of organic material in an inert atmosphere at high temperatures, typically above 1000°C, all non-carbon elements being driven off in the process. Carbothermal reduction: Chemical reduction of a material by carbon at elevated temperature. Ceramic matrix composite: A composite consisting of a filler material, frequently fibers, embedded ina ceramic matrix. Chalcogenide: A binary compound of a chalcogen (Le., any of the elements oxygen, sulfur, selenium, and tellurium) with a more electropositive element orradical. Chemical vapor deposition: Deposition ofsolid material when gaseous reactants encounter ahot surface. Laser assisted CVD employs a hot laser-focus rather than a hot surface. Chemical vapor infiltration: A process whereby a reactive chemical species is generated in the vapor phase and allowed toreact with a solid substrate, thus modifying its chemistry. Coefficient of thermal expansion: The change in volume-per-unit volume produced by a one-degree rise intemperature. Compressive strength: The ability of a material to resist a force that tends to crush; the crushing load at specimen failure divided by the original cross sectional area of the specimen. Creep: The change in dimension of a plastic material under load over a period of time, not including the initial instantaneous elastic deformation.

336

Glossary

Crimp: The waviness of a fiber. Crimp can be either natural or can be mechanically or thermally induced. Debye-Scherrer pattern: A method of x-ray diffraction analysis widely used to identify polycrystalline materials. Dielectric constant: The ratio of the capacity of a condenser having a dielectric material between the plates to that of the same condenser when the dielectric is replaced by a vacuum; a measure ofthe electrical charge stored per unit volume atunit potential. Dielectric loss: A loss of energy eventually showing up as a rise in the temperature of a dielectric material placed inan alternating electric field. Diphasic gel: Agel consisting oftwo discrete, mutually incompatible phases. Discotic nematic liquid crystal: A liquid crystal formed from discotic (disk-shaped) molecules, which are more orless planar and which are stacked. Disinclinations: Rotation defects in the arrangement ofdiscotic (disk-shaped) molecules in a liquid crystal. Dispersion shifted fiber: An optical fiber engineered tocontrol the optical dispersion profile. Dry spinning: A process whereby a viscous solution is spun through spinneret orifices and the excess solvent isremoved in a drying column. E-beam curing: A process whereby green fibers are rendered infusible through irradiation with an electron beam. E-beam evaporation: Evaporation ofmaterial by electron beam irradiation. Eutectic mixture: The one mixture ofa set of substances able to dissolve in one another as liquids that, of all such mixtures, liquefies at the lowest temperature. On cooling, the components ofthis mixture separate simultaneously as an intimate mixture ofsolids. Evanescent field: That portion of the light being transmitted by an optical fiber which escapes into the sheath, outside the core. Fiber grating: A structure generated by periodically varying the refractive index of the core of an optical fiber. The period of the optical grating can be adjusted to reflect specific wavelengths. These are used in wavelength selective mirrors, resonant cavities and bandpass filters. Fragile melt: Molten material which isdifficult toprocess into fibers. Free radical: An especially reactive atom or group of atoms that has one or more unpaired electrons. Gauge length: The length over which deformation ofa sample ismeasured. Gel: A coherent mass consisting of a liquid in which particles too small to be seen in an ordinary optical microscope are either dispersed orarranged in a fine network throughout the mass. A gel may be elastic and jelly-like, as gelatin orfruit jelly, or solid and rigid, as silica gel. Glass transition temperature (T g): The temperature below which an amorphous material behaves like a solid rather than a viscous liquid. The glass transition occurs over a temperature range, whose location depends somewhat on the cooling rate and the method ofmeasurement. Graded index fiber: An optical fiber having a fixed refractive index in the sheath and a continuously variable refractive index in the core, Grain size: The size ofadiscrete crystallite orgrain within a solid material. Graphene layer: The structural unit in carbons resulting from the carbonization of an organic precursor; each layer consists ofa few fused aromatic rings. Graphitization: The solid state transformation of thermodynamically unstable non-graphitic carbon into graphite by thermal activation, usually attemperatures above 1800°C.

Glossary

337

Green fiber: A fiber precursor which must be heat-treated to obtain the final desired form. Group delay: The time delay between components ofa signal having different wavelengths. Incubation period: The time required for the onset ofnucleation in a crystallization process. Intercalation of carbon fibers: The crystal structure of graphite is of a kind that permits the formation of many compounds, called lamellar or intercalation compounds, by insertion ofmolecules orions between the graphitic layers. Internal flaws: Flaws within the volume ofamaterial, such as voids. Inviscid melt spinning: A process which allows low viscosity mollen material to be spun into fibers. The low viscosity jetischemically stabilized rather than rapidly solidified. Isotropic: Exhibiting properties, such as velocity oflight transmission orrefractive index, with the same values when measured along any direction in a material, i.e., along any of the principal axes ofacrystal. Laser heated floatzone process: In this process, a circumferential laser is placed around a preform rod to zone refine a segment of the material while simultaneously updrawing a single crystal fiber. Liquid crystal: An organic liquid whose physical properties resemble those of a crystal in the formation ofloosely ordered molecular arrays similar to a regular crystalline lattice and in the anisotropic refraction of light. Liquidus temperature: The lowest temperature atwhich a material will be completely liquid. Lithography: The process of producing patterns on semiconductor crystals for use as integrated circuits. Mat: A fibrous material for reinforced plastic use consisting of randomly oriented chopped filaments orswirled filaments with a binder. Melt spinning: Extrusion ofa molten polymer through a spinneret toform fibers. Mesophase pitch: See mesopitch. Mesophases: Liquid crystals, sometimes called mesophases, occupy the middle ground between crystalline solids and ordinary liquids with regard to symmetry, energy, and properties. Mesopitch: A carbonaceous solid, consisting primarily of a complex mixture of polycyclic aromatic compounds, which can form an anisotropic liquid crystal mesophase. Metal matrix composite: A composite consisting of a filler material, frequently fibers, embedded in a metal matrix. Microcoil: Acoiled structure of microscopic scale. Microfibers: Fibers generally 1-25 um diameter. Microhardness: The indentation hardness ofmicroscopic areas ofamaterial. Micropillar: A short inorganic fiber generally 1-25 urn diameter. Microsprings: Coiled microfibers. Microstructure: The structure ofa material on a microscopic scale. Mode coupling: Interaction of modes resulting in the transfer of light intensity between the guided (core) modes and the radiation (cladding) modes. Mode: Apath along which light may travel, e.g. in an optical fiber. Modulus: The ratio of stress to strain in a material over the range for which this value is constant. The type of modulus, which is measured, depends on the method of measurement, e.g., dynamic modulus, compressive modulus, elastic or tensile (Young's) modulus, shear modulus, torsion modulus, sonic modulus. Morphology: The size, shape and structure ofa material and its constituents. Mullitization: Crystallization ofmullite from stoichiometric alumina-silica mixtures.

338

Glossary

Nanocrystalline: Composed of crystals orcrystalline structural units which have nanometer dimensions. Nanofibers: Fibers generally 1-25 nm diameter. Nanopillar: A short inorganic fiber generally 1-25 nm diameter Nanotubes: Hollow nanofibers. Carbon nanotubes consist of one or more cylindrical structures orshells, each composed ofa single tubular graphitic layer. Nanowire: An electrically conductive nanofiber. Nucleation: The initial process that occurs in the formation of a crystal from a solution, a liquid, or a vapor, in which a small number of ions, atoms, or molecules become arranged in a pattern characteristic of a crystalline solid. This creates a site upon which additional particles are deposited as the crystal grows. Numerical aperture: Ameasure ofthe light gathering power ofan optical system. Optical amplifier: A photonic device which boosts a light signal in optical communications to remedy transmission loss. Optical dispersion: The change in refractive index for a transparent material as optical wavelength ischanged. Optical loss: Attenuation ofoptical light intensity during propagation through a fiber. Optoelectronic: Pertaining to electronic devices for emitting, modulating, transmitting, and sensing light. Photonic band gap: The energy gap between the energy minimum of the electron conduction band and the energy maximum of hole valence bands which occur at the same location in momentum space, allowing electrons and holes to recombine and radiate photons efficiently. Favorable band gaps occur with III-V compound semiconductors. Polycrystalline: Consisting ofcrystals variously oriented. Polycyclic aromatic hydrocarbons: Organic compounds consisting of fused aromatic ring moieties. Polymer matrix composite: A composite consisting of a filler material, frequently fibers, embedded in a polymer matrix. Polymorphism: The quality or state of being able to assume different forms such as the property ofcrystallizing in two ormore forms with distinct structure. Precursor fiber: A fiber from which another fiber will be formed . Preferred orientation: The anisotropic arrangement of crystals or polymer molecules; in fibers, thisis usually caused by process stresses during spinning ordrawing. Preform: A shaped form from which fibers are subsequently obtained. The use of a preform facilitates handling and control ofuniformity. Pyrolysis: Chemical change brought about by the action ofheat. Radiation modes: Modes which propagate in an optical fiber through the cladding (sheath). Rayleigh scattering: Scattering of light in a material because of microscopic fluctuations in refractive index, e.g., fluctuations in composition ordensity. Roving: A term used to designate a collection of bundles of continuous filaments either as untwisted strands or as twisted yarns. Glass rovings are predominantly used in filament winding. Sigmoidal curve: Acurve shaped like the letter S. Sintering: The welding together ofsmall particles by applying heat below the melting point. Sliver: An untwisted strand orrope oftextile fiber produced by a carding orcombing machine and ready for drawing, roving, orspinning.

Glossary

339

Softening point: The temperature above which a fiber will rapidly deform under its own weight. Sol-gel processing: A generic term covering process routes which differ from one another mainly by the nature of the starting chemicals, the most commonly used being sols or solutions of organometallic species. Gelation of a sol gives a viscous product which can be shaped as fibers. Specific modulus: Modulus divided by material density. Specific strength: Tensile strength divided bymaterial density. Spin orientation: Molecular orstructural orientation created inthe spinning process. Spinneret: A type of extrusion die, l.e., a metal plate with many tiny holes, through which a plastic melt isforced to make fine fibers and filaments which are hardened bycooling in air orwater orby chemical action. Staple fibers: Fibers of spinnable length manufactured directly or by cutting continuous filaments toshort lengths, usually one half totwo inches long. Step index fiber: An optical fiber having fixed refractive indices for the sheath and core components. Stoichiometric: Pertaining tothe quantitative relationship between constituents in a chemical substance. Strain rate: The rate ofdeformation as in a tensile test. Superconductor: A material which exhibits complete disappearance of electrical resistance especially atvery low temperatures. Supercool: To cool below the freezing point without solidification orcrystallization. Surface flaws: Flaws on the surface ofa material, such as surface scratches. Tensile strength: The maximum tensile load per unit area oforiginal cross section, within the gauge boundaries, sustained by the specimen during a tension test. Thermal conductivity: The ability ofa material to conduct heat; the physical constant for the quantity of heat that passes through a unit cube of a material in a unit time when the difference in temperature oftwo opposite faces isone degree. Thermophoresis: Transverse motion of particles in a gas stream caused by a thermal gradient perpendicular tothe direction ofthe gas stream. Thermoplastic polymer: A polymer capable of softening or fusing when heated and of hardening again when cooled. Thermosetting polymer: A polymer capable of becoming permanently rigid when heated or cured. Toughness: The energy required to break a material. This energy is equal to the energy under the stress-strain curve. High toughness leads tohigh damage resistance. Transmission loss: See optical loss. Turbostratic: A type of crystalline structure where the basal planes, such as graphene layers, have slipped sideways relative to each other, causing the spacing between planes to be greater than ideal. This structure is found in incompletely heat-treated carbon and boron nitride. Updrawing: A process whereby a rod is dipped into a pool of a viscous melt and then is raised vertically. The cooled melt adheres to the rod and a fiber is formed which cools and solidifies. Vapor-liquid-solid growth: A vapor toliquid tosolid phase transformation, the most common way to grow whiskers. Vapor-solid growth: A direct vapor to solid phase transformation, responsible for the catalyzed growth ofsome fibers.

340

Glossary

Volume resistivity: The ratio of the direct voltage applied to electrodes on opposite faces of a unit volume with unit separation to that portion of the current between them passing through a unit cross sectional area. Waveguide: Adevice (aduct, coaxial cable, orglass fiber) designed toconfine and direct the propagation ofelectromagnetic waves (Le., light). Weibull distribution function: A probability distribution function which can be used to describe the scatter in tensile strength data for a given lot of fibers having the same cross section area and test gauge length. Weibull modulus: A parameter which characterizes the width of a monomodal Weibull distribution. Wet spinning: A process whereby a viscous polymer solution isextruded through a spinneret into a spin bath. The spinneret is generally, but not always, placed in the spin bath, a coagulation bath in which solvent diffuses out of the extruded material and a nonsolvent diffuses into the extrudate. Whiskers: Filamentary single crystals X-ray amorphous: Non-crystalline, as determined by x-ray diffraction. Young's modulus: See modulus.

INDEX A Absorption, 173-174 Acid leaching, 128, 148, 158, 165--166 Acid-resistant glass fibers, 148-149 Acrylic polymers, 124 Advanced liquid phase processes, 26-29 Aerospace applications, 70-71 ,140,158,166, 315,316-318,322,323,324,325 A-glass, 149 Alkali-resistant (AR) glass fibers, 102, 146-148 Alkoxides, 206 Almax,211,219,222 Altex, 218 Alumina fibers, 51,53,64,70,88-89,207-225 properties, 90, 219-224 a-Alumina fibers, 211 ,220,224 Aluminate fibers, 4,70, 88-89 Aluminate glass fibers, 3, 98-99 Alumina-zirconium fibers, 215--216 Aluminum oxide, 53 Aluminum silicate fibers, 71, 85 Ammonia, 22, 5~0 Amoco P-series, 248 Amorphous fibers, 3,4,20, 62, 64,65, 87, 88, 92 from inviscid liquids, 103-113 Anisotropy, 88 Arc discharge, 13, 24-25 Architectural applications, 148 Asbestos, 3, 11, 39 Astroquartz (AO), 90 Automobiles, 318, 321, 324, 325

B Basal growth, 12 Basicstructural unit (BSU), 234-235 Balleries, 315 Bend stress relaxation, 291 Bicomponent fibers, 6, 41 , 85, 92, 97, 156-162 Bioactive materials, 153,326 Birefringence, 91 , 179-180 Bone biocompatibility, 153,326 Borazine fibers, 60, 69 Boron/aluminum materials, 71 Boron doping, 275--276 Boron/epoxy tapes, 70

Boron fibers, 3, 51 , 57, 61 , 64, 65-67 Boron-free fibers, 95,131-132,133-134,148, 214 Boron nitride fibers, 25, 37, 5~0 , 69 Boron oxide fibers, 5~0 , 69 Boron/tungsten (BIW) fibers, 21,56,57,61,63, 66-67,70-71 Borosilicate fibers, 130-131, 133, 165--166 Bragg gratings, 197, 198 Braking systems, 324-325 Buckyballs, 36, 37 Bushing process, 84, 85, 86, 88, 93,129,157 Butyl indium, 28

c CAD (computer-aided design), 73, 74 Calcination, 207 Calcium aluminate fibers,99,101-102,110-111 Carbon black, 23 Carbon fiber-reinforced carbon (C/C) composites, 320-326 Carbon fibers, 3, 11, 17-18, 20, 67-B8, 71 , 233-235 commercial applications, 37,41-43, 315 composaes, 41, 316-320, 322-326 infiltration, 59 processing of, 64, 235--245 structures and properties, 34-36, 245-261 Carbon ion bombardment, 25 Carbonization, 238-239, 245 Carbon nanotubes, 3, 11, 18, 20, 24, 25, 42-43, 60 endless, 60 modulus, 37, 38, 70, 72, 88-90 structures and properties, 36-39 Carbothermal reduction, 13,23-24,34 CC-glass, 149 CemFIL, 146, 147 Centrifuge processes, 11,28-29,129,160 Ceramic matrixcomposites (CMCs), 40,67-B8, 70,322-326 C-glass, 149 Chalcogenide glass fibers, 100-101 Chemical mixing processes, 23-24 Chemical resistance, 145--146 Chemical vapor deposition (CVD), 15,28,55

342 hot filament processes, 21-22, 55-56, 58, 65-68,69 metal catalyzed, 15-20,34,48,63,64,65 modified (MCVD), 185-190 plasma (PCVD), 22, 58, 190-191 See also Laser-assisted chemical vapor deposition Chemical vapor infiltration (CVI), 22-23, 59-60, 68-69,322 Coated carbon fibers, 261 Coefficient ofthermal expansion (CTE), 257-258,322 Commercial applications carbon fibers, 37,41-43,315 composites, 316-326 continuous fibers, 55-56, 65--Q8, 70-75, 129-132 glass fibers, 134-136, 147-148, 164-165 quartemary fibers, 102 silicon fibers, 39-40, 68 Compressive strength, 256 Computers, 149 Conical tips, 32, 33 Containerless laser melt process, 72, 85, 86, 107-108 Continuous fibers, 3, 4,6 dry spinning processes, 123-128 structures and properties, 60-70 vapor phase processes, 47--£0 See also Melt spinning processes Corundum fibers, 218-219 Creep,222-224,256-257,288-290 Crimped staple fibers, 161, 162 Cryogenic fibers, 3, 113 Cubic stabilized zirconia (CSZ), 225, 226 D

Damage resistance. See Mechanical toughness D-glass, 149-150 Diameter,4,15 Diamond, 21, 22,41, 58, 75 Dielectric constant, 149-152 Diodes, 39,42 Disclinations, 249 Discontinuous fibers, 3, 4 Dispersion, 174-179, 195 Double crucible technique, 97, 98, 157, 181-182 Downdrawing, 85, 86, 92 Draw towers, 191-192, 194 Dry spinning, 3,123-128,164-165

Index

E Edge-defined film-fed growth (EFG), 113-115, 118 E-glass, 84, 94-95,129 boron- and fluorine-free, 95,131-132, 133-134,148 borosilicate, 130-131, 133 corrosion-resistant (E-CR), 95, 132 hollow, 157-158 structures and properties, 83, 90, 93, 132-136, 138 Electrical and electronics applications, 140, 149, 152,315,321-322 Electrical conductivity, 258-259, 294 Electromagnetic shielding, 41, 150, 152 Endless fibers. See Continuous fibers Engines, 40,67--£8,71 ,322,323,324,325 Erbium doped fiber amplifier (EDFA), 195, 196 Etching, 29, 63 Evanescent field, 172, 173 Extrusion, 85, 244, 276 F Feedback control, 54-55 Fiber formation, 3, 103, 106, 220 Fiber FP, 70, 90, 211, 222 Fiberglass, 85,129,134-136 Fiber gratings, 196-197 Fiber optics. See Optical fibers Field-emitting structures, 30, 31, 40 Fluoride, 85, 97,101 Fluorophosphate glass fibers, 101 Four wave mixing (FWM), 196 Fracture toughness, 70. See also Mechanical toughness Fragile melts, 82, 84, 85, 87, 95-102 Fullerenes, 38, 60 Fumaces, 325

G Gelation, 206-207 General-purpose glass fibers, 94-95, 129-136 Germanium, 3,20, 21 , 34, 63, 64 Glass fibers, 3,124 chemical resistance, 145-149 commercial applications, 134-136, 147-148, 164-165 general-purpose, 94-95, 129-136 See also E-glass; Fiberglass; Inviscid melt spinning processes; Melt spinning

343

Index processes; Silica glass fibers; Silicate glass fibers Gold, 18,30 Graded index fibers, 172 Graphene layers, 234-235 Graphitization, 34-36, 234, 245 Green fibers. See Precursor fibers

H Hafnium carbide whiskers, 17 Heat management systems, 325 High-modulus (HM) fibers, 95,136,140-145, 148,235,247,250 High silica fibers, 165-166 High-strength (HS) fibers, 136-140, 148,235, 247,250 Hollow fibers, 58, 156-160 Hot filament chemical vapor deposition, 21-22, 55-56,58,65-68,69 Hybrid fiber forming processes, 100-101 Hydridopolysilazane (HPZ), 300, 302,303, 304

IMS-54,90 Indium,28 Industrial applications, 70, 149-150 Infrared (IR) absorption, 173-174 Infusible PCS fibers, 275 Insulation, 315 Inviscid melt spinning (IMS) processes, 84-86, 87-88,113-119 aluminate glass fibers, 90 amorphous fibers, 103, 105-107 metal fibers, 108-109 oxide glass fibers, 110-111 Ion bombardment, 25 Iron, 23 J

Jets. See liquid jet K Kashima Carbonic series, 248 Kumada rearrangment, 267, 269-270

L

Lanathana, 140 Laser ablation, 13,20-21 ,34 Laser-assisted chemical vapor deposhion (LCVD), 47-55,64, 65 automatic process control, 54-55 high pressure process (HP-LCVD), 53-54, 61-65,72,73 low pressure process (LP-LCVD), 49-53, 61, 62,65 rapidevaluation ofnew fibers, 71-73 rapid prototyping, 50, 53, 73-75 structure and property relationships, 61-65, 69 Laser-heated float zone (LHFZ) growth, 115-118, 159 Laser vaporization, 25, 60 Leaching, 165-166 Lead glass, 151 liquefied gases, 113 liquid jet, 103, 106 lifetime, 106, 107, 108 solidification, 111-113 liquid phase, 3-4, 26-29 liquid pitch melts, 11 liquidus temperature, 81 ,82-83, 93 lithography, 29, 30, 74 M Magnesium aluminosilicate, 95-97 Magnetron sputtering system, 58 Mats, 136 Mechanical prepolymer deposition, 74 Mechanical toughness, 70 Medical applications, 29,153,319-320 Melt assisted spinning, 237 Melt spinning processes, 81-86 generic, 85-86 and PCS, 269-272 structure and property relationship,87-92 viscous, 81-84 See also Fragile melts; Inviscid melt spinning processes; Strong melts Melt viscosity, 82 Mesophase pitch (Mesopitch, MP), 28-29, 233, 235,240-245,247-250 Metal catalyzed chemical vapor deposfion, 15-20,34,48,63,64,65 Metal matrixcomposites (MMCs),40, 320-322 Metal particle catalysts, 11 , 12, 13 Metglas, 105 Microcoils, 52,65 Microfilters,74

344 Micromechanical processess, 51, 53, 75 Micropillars, 11,29,30 Microporosity, 69 Microsolenoids, 52, 53 Microsprings, 17,51,58,69 Microstructures, 51-53, 54, 65, 74, 75 Microtexture, 60, 247-250 Microtubes, 58, 65 Microwaves, 152 Military applications, 140 Modified chemical vapor deposition (MCVO), 185-190 Modulus, 62, 70, 72, 88-90, 94 carbon nanotubes, 37, 38, 70, 72, 88-90 glass fibers, 136-137 single crystal fibers, 20 Molding applications, 41 Mullite, 208-209, 212-215, 216-218 N Nanopillars, 11 Nanotubes, 13, 24, 25. See also Carbon nanotubes Nanowires, 11, 13,20--21,26-27 Nextel fibers, 90, 218, 222, 224 Nicalon fibers, 68, 70, 71 , 72, 222, 279, 283, 286-287,290,291,292,320,323 Nickel-based alloys, 323 Niobium monocarbide (NbC),23 Nitridation, 22 Non-round cross sections, 154-155 Nylon, 154, 155

o Optical amplifiers, 179, 195-196 Optical fibers, 4,85, 92,169,198-199 devices, 194-198 drawing processes, 191-194 fabrication, 180--191 transmission principles, 169-180 Optical loss, 172-174 Outside vapor deposition (OVO), 183 Oxidation, 259-261, 291-293, 304-305 Oxide fibers, 110--111, 118-119 Oxygen-free SiC fibers, 272-275 Oxynitridefibers, 13,24,142-145 p

Packaging applications, 41, 71 , 322

Index Partially stabilized zirconia (PSZ), 225, 226 Perhydropolysilazane (PHPSZ), 307-308, 310 Phosphorescence, 116 Photonic band-gap (PBG) microstructures, 53, 75 Pitch, 235, 236, 239-245. See also Mesophase pitch Planar flow casting process, 103, 104 Plasma chemical vapor deposition (PCVO), 22, 58,190--191 Plasma processes. See Arc discharge; Carbon ion bombardment; Laser vaporization Polyacrylonitrile (PAN) fibers, 35, 36, 233,235, 236,237-239,247 Polyborodiphenylsiloxane (PBOPSO), 268 Polycarbonate substrates, 51 Polycarbosilane (PCS), 68, 266-268, 269-276, 294,307 Polycarbosilazane (PCSZ) precursors, 300--301, 302,303-304 Polycrystalline fibers, 62, 64, 65, 70, 84-85 alumina, 210--211 Polycyclic aromatichydrocarbons, 36 Polydimethylsilane (POMS), 267,268 Polymer impregnationpyrolysis (PIP), 322 Polymer matrix composhes (PMCs), 316-320 Polysilapropylene (PSP), 268-269 Polysilasilazane (PSSZ), 301 Polysilazane (PSZ) precursors, 299-300, 308 Polyzirconoxanes, 226-227 Porous hollow glass fibers, 156, 158 Potassium lithium niobate, 115 Powder-in-tube process, 118, 159 PRO-l66, 215-216, 222 Precursor (green) fibers, 3, 4,124 inchemical vapor infiltration, 59-60 PAN and pitch, 235-245 polycarbosilazane (PCSZ), 300--301, 302, 303-304 polysilazane (PSZ), 299-300, 308 Preforms,3,84,85,92,128, 163-164, 181, 191-193 Pressure vessels, 319 Prosthetics, 153, 326 Protective coatings, 192-193, 261 Pulse mode dispersion (PMD), 179 Pyrolytic processes, 23, 275-276, 322 Q

Ouaternary fibers, 84, 85, 90, 101-102 Quasi-stoichiometric SiC fibers, 266, 275-276 Quench rate, 103, 104

Index

R Radiation cured PCS fibers, 272-274 Radiation resistance, 150-151 Rapidevaluation processes, 71 -73 Rapid prototyping, 50, 53, 73-75 Rapid solidification (RS) process, 103-105 Rayleigh scattering, 173, 181 Rayleigh waves, 112 Reinforcing fibers, 66, 134 R-glass, 136-137, 138-139 Ribbons, 85, 103, 104-105, 155 Rice hull process, 13, 23, 34 Rockets, 73, 322 Rovings, 134

s Saphikon, 90 Sapphire fibers, 70, 71 , 72,114-115,118-119 Scanning devices, 32, 40 Scattering, 173 Selective laser sintering (SLS), 74 Sen-assembly processes, 11, 26-27, 60 Sen-catalyzation , 13 Sen-propagating high-temperature synthesis (SHS),24 Semiconductors, 34, 39, 40, 152 Sensor fibers, 73,74,108, 197-198 Service temperature, 67--Q8, 71,136 S-glass,83, 84, 90, 93,95, 137,149,157 Sharpened whiskers, 30-32, 33, 39-40 Sheath/core fibers, 21 , 22, 34, 58, 63,66-68, 69,97,156-160 Short fibers, 3, 6, 11 , 65. See also Amorphous fibers; Polycrystalline fibers SiAION fibers, 13,90,142-143,145 Side-by-side bicomponent fibers, 156, 160-162 Side growth, 12, 18,35,37,63,67 Silica glass fibers, 92-93 dry spinning, 123-128, 164-165 high-temperature, 162-166 Silicate glass fibers, 3,85, 88 fragile melts, 95-97 general-purpose, 129-136 non-roundcross section, 154-155 special-purpose, 136-153 structural,93-95, 129-162, 156-162 Silicon carbide/aluminum nitride composites, 71 Silicon carbide/carbon (SiC/C) fibers, 21 , 56, 63, 67--Q8,68, 70, 71, 72 Silicon carbide (SiC) fibers, 59, 64, 67--Q9, 70, 71 ,72,265-266 matrix composiles, 322, 323, 324

345 oxygen-free, 272-275 quasi-stoichiometric, 266, 275-276 structure and properties, 276, 279-280, 283-294 whiskers and nanowhiskers, 17,20,23-24, 34,40,70 Silicon carbide/titanium composites, 71 Silicon dioxide, 22 Silicon fibers, 11, 62, 63 Silicon nitride, 23, 63, 64, 69, 299 boride based fibers, 309-313 Si-C-N-O and Si-C-N fibers, 299-306, 312 Si-N-O and Si-N fibers, 306-309 whiskers, 17,20,24,63 Silicon oxide, 59 Silicon oxycarbide (Si-C-O) fibers, 70, 266-272, 276-279,281-294 passim Silicon whiskers and nanowhiskers, 3, 15-17, 18-19,20-21,39-40 structures and properties, 30-34 Silver nanowires, 11, 26-27 Single crystal fibers, 3,51,62,63,64,65,70,72 from inviscid melts, 86,113-119 melt processes, 86 See also Whiskers Sintering, 74,189-190,207 Sliver yarn, 124-125 Slurry processes, 70 Snell's law, 169-170 Soda-lime-silica, 181 Sodium borosilicate, 181 Sol-gel processing, 70,126-128 ceramicoxide fibers, 205-207 optical fibers, 193-194 Solid phase, 11, 12 advanced processes, 29-30 Solution-liquid-solid phase, 11 Solution-liquid-solid (SLS) phase, 28, 29 Solvent spinning. See Dry spinning Spacecraft, 317-318,321 ,322,323 Specialty fibers,6, 136-153 Spinorientation, 91 Sporting goods,70, 71,140,318,321 Steel,108 Step index fibers, 171-172 Stereolilhography, 74 Stiffness. See Modulus Strain sensors, 197 Strength, 66, 70, 91-92. See also Tensile strength Strong melts, 82-83, 87, 92-95 Structural fibers, 4, 93-95,129-162,156-162 Superconducting fibers, 42,116-118,152, 158-160 Supercooled melts, 86,97-99

346

Index

T Tellurite glass fibers, 98 Tensile strength, 91, 253-256, 284-285,

287-288

silicon nitride fibers, 312, 313 Ternary fibers, 88 Tetraethylorthosilicate (TEOS), 164, 193,323 Tetragonal zirconia polycrystals (TZP), 225-226 Thermal expansion, 257-258 Thermal stability silicon carbide fibers, 280-284 silicon nitride fibers, 302-304, 308-309, 311 Thermophoretic deposition, 189 Time-ta-failure, 224 Tip groW1h, 12,35,63,67 Titanium tetrabutoxide, 272 Tokawhiskers, 23 Transistors, 42 Transition alumina fibers, 211-212, 216, 222,

315

Transport properties, 258-259, 293-294 Traveling solvent zone melting (TSZM), 118 Tnchlorosilylamino-dichloroborane (TAOS), 311 Trimethylamine alane, 51 Tyranno, 272, 290

u Ultrahigh-modulus (UHM) fibers, 141-145,235 Ultrahigh temperature (UHT) fibers, 71-72, 162 Ultrapure silica fibers, 3,126-128, 162, 163-165 Uniformity. See Strength Updrawing, 85,86,97-100,102

v

Vapor-liquid-solid growth; Vapor-solid groW1h Vapor-solid (VS) growth, 3,13-14,24,34,63,64 Vertical axialdeposition (VAO), 183-185 Vibrational absorption, 173-174 Viscous melt spinning processes, 3,4,81-84

w Water glass solutions, 124-126 Waveguide physics, 169-172 Weibull distribution function, 253-254, 285 Weight-sensnive applications, 67,132,136,158,

321,323

Whiskers, 3,4, 11, 12 metal catalyzed, 15-20, 21 , 64 from organicsolvents, 27-29 structures and properties, 20, 30-34 See also Silicon carbide whiskers and nanowhiskers Wire drawing, 108, 109 Work-ta-break. See Mechanical toughness

x Xerogels, 207 X-ray lithography (L1GA), 74 X-ray transparency, 319-320 y

Yajima process, 267-269, 307 Young'smodulus, 37, 38, 250-251, 252-253 Ytlria-modified fibers, 84, 89, 90,141 Yttrium aluminum garnet (YAG) fibers, 3, 4, 72,

108,227-229

Vacuum microelectronics, 30, 40 Vapor-liquid-solid(VLS) groW1h, 3,12-13,15,

20.1 n , 23,24,25,29,34,63,64 Vapor phase, 3, 11-26,70,75. See also

Chemical vapor deposition; Chemical vapor infiltration; Laser vaporization;

z Zinc oxide modified fibers, 90, 140 Zirconia, 24,115,225-227,31