Carbon-carbon composites

CARBON-CARBON MATERIALS AND COMPOSITES Edited by John D. Buckley National Aeronautics and Space Administration Langley...

Author: John D. Buckley | Dan D. Edie

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CARBON-CARBON MATERIALS AND COMPOSITES

Edited by

John D. Buckley National Aeronautics and Space Administration Langley Research Center Hampton, Virginia

Dan D. Edie Clemson University Clemson, South Carolina

El

NOYES PUBLICATIONS Park Ridge, New Jersey, U.S.A.

Copyright 0 1993 by Noyes Publications No part of this book may be reproduced or utilized in any form or by any means, electronic or mechanical, including photocopying, recording or by any information storage and retrieval system, without permission in writing from the Publisher. Library of Congress Catalog Card Number: 92-35012 ISBN: 0-8155-1324-0 Printed in the United States Published in the United States of America by Noyes Publications Mill Road, Park Ridge, New Jersey 07656 10 9 8 7 6 5 4 3 2 1

Library of Congress Cataloging-in-Publication Data Carbon-carbon materials and composites / edited by John D. Buckley and Dan D. Edie. p. cm. Includes bibliographical references and index. ISBN 0-8155-1324-0 1. Carbon composites. I. Buckley, John D. II. Edie, Danny D. (Danny Dale), 1943- . TA418.9.C6C27 1993 620.1'93--dc20 92-35012 CIP

Contributors

John D. BucMey NASA Langley Research Center Hampton, Virginia

Frank K. KO Drexel University Philadelphia, Pennsylvania

Robert L. Burns Fiber Materials, Incorporated Biddeford, Maine

N. Murdie Southern Illinois University at Carbondale Carbondale, Illinois

Russell J. Diefendorf Clemson University Clemson, South Carolina

Louis Rubin The Aerospace Corporation El Segundo, California

J. Don Southern Illinois University at Carbondale Carbondale, Illinois

James E. Sheehan General Atomics San Diego, California

Dan D. Edie Clemson University Clemson, South Carolina

E.G. Stoner Clemson University Clemson, South Carolina

C.P. Ju Southern Illinois University at Carbondale Carbondale, Illinois

M.A. Wright Southern Illinois University at Carbondale Carbondale, Illinois

John J. Kibler Materials Science Corporation Blue Bell, Pennsylvania vii

Notice To the best of the Publisher's knowledge the information contained in this book is accurate; however, the Publisher assumes no responsibility nor liability for errors or any consequences arising from the use of the information contained herein. Final determination of the suitability of any information, procedure, or product for use contemplated by any user, and the manner of that use, is the sole responsibility of the user. The book is intended for informational purposes only. Carbon-carbon raw materials and processes could be potentially hazardous and due caution should always be exercised in the handling of materials and equipment. Expert advice should be obtained at all times when implementation is being considered. Mention of manufacturers' trademarks or trade names does not constitute endorsement by the authors, the U.S. government, or the Publisher.

...

Vlll

Preface

Carbon-carbon composites, which have been used extensively for missile applications, were a part of NASA's Apollo spacecraft heat shield system. The development of carbon-carbon materials began in 1958 and was nurtured under the U.S. Air Force space plane program, Dyna-Soar, and by numerous thermal protection systems developed by NASA for aerospace research. The purpose of this book is to present data and technology relating to the materials and structures developed for the production of carbon-carbon materials and composites. The text is composed of papers written by noted authors in their areas of expertise relating to the processes and production of these material systems and structures. The subject matter is arranged to lead the reader step by step through the materials processing, fabrication, structural analysis, and applications of typical carbon-carbon products. The information presented in the text is limited to data that can or has been published in the open literature including: fiber technology, matrix material, design of composite structures, manufacturing techniques, engineering mechanics, protective coatings, and structural applications using carbon-carbon materials and structures. The editors thank the authors for their contributions of time and effort in the development of this book. The use of trademarks or names of manufacturers in this book is for accurate reporting and does not constitute an official endorsement, either expressed or implied, of such products or manufacturers by the National Aeronautics and Space Administration. John D. Buckley NASA Langley Research Center Hampton, Virginia Dan D. Edie Clemson University Clemson, South Carolina

V

Contents and Subject Index

.

1 CARBON-CARBON OVERVIEW . . . . . . . . . . . . . . . . . . . . . . . . 1 John D. Buckley Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1 Carbon Fibers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4 Carbon Fibers in Carbon Matrix . . . . . . . . . . . . . . . . . . 5 Discontinuous Fiber Composites . . . . . . . . . . . . . . . . . . 5 Continuous Fiber Composites . . . . . . . . . . . . . . . . . . . . 6 Chemical Vapor Deposition . . . . . . . . . . . . . . . . . . . . . . 7 Carbonized Organic Composites . . . . . . . . . . . . . . . . . . 9 Mechanical Properties . . . . . . . . . . . . . . . . . . . . . . . . . 11 Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 14 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 14 Bibliography . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

.

2 CARBON FIBER MANUFACTURING . . . . . . . . . . . . . . . . . . . 19 D.D. Edie and R.J. Diefendorf Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20 Manufacture of PAN-Based Carbon Fibers . . . . . . . . . 20 Solution Spinning of PAN Precursor Fibers . . . . . . . . . 20 Melt-Assisted Spinning of PAN Precursor Fibers . . . . . 22 Heat Treatment of PAN Precursor Fibers . . . . . . . . . . 23 Oxidation of PAN Precursor Fibers . . . . . . . . . . . . . 23 Carbonization and Graphitization . . . . . . . . . . . . . . 24 Manufacture of Rayon-Based Carbon Fibers . . . . . . . . 26 Manufacture of Pitch-Based Carbon Fibers . . . . . . . . . 27 Mesophase Pitch . . . . . . . . . . . . . . . . . . . . . . . . . . . 28

ix

Contents and Subject Index

Melt-Spinning Mesophase Precursor Fibers . . . . . . . . 29 Heat Treatment of Mesophase Precursor Fibers . . . . . 31 Oxidation of Mesophase Precursor Fibers . . . . . . . . 31 Carbonization and Graphitization . . . . . . . . . . . . . . 31 Isotropic Pitch-Based Carbon Fibers . . . . . . . . . . . . . 33 Manufacture of Vapor-Grown Carbon Fibers . . . . . . . . 33 Mechanical Properties of Carbon Fibers . . . . . . . . . . . 35 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 37 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 37

3. EFFECT OF MICROSTRUCTURE AND SHAPE ON

41 CARBON FIBER PROPERTIES . . . . . . . . . . . . . . . . . . . . . . . D.D. Edie and E.G. Stoner Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 42 Carbon Fiber Processes . . . . . . . . . . . . . . . . . . . . . . . 43 Effect of Graphite Structure on Fiber Properties . . . . . 44 Brittle Failure Mechanism . . . . . . . . . . . . . . . . . . . . . . 45 Microstructure of Carbon Fibers . . . . . . . . . . . . . . . . . 47 Microstructure of PAN-Based Carbon Fibers . . . . . . . . 47 Microstructure of Pitch-Based Carbon Fibers . . . . . . . 48 Effect of Microstructure on Fiber Properties . . . . . . . . 53 Effect of Microstructure on Tensile Properties of Carbon Fibers . . . . . . . . . . . . . . . . . . . . . . . . . . . 54 PAN-Based Carbon Fibers . . . . . . . . . . . . . . . . . . 54 Pitch-Based Carbon Fibers . . . . . . . . . . . . . . . . . . 55 Effect of Microstructure on Compressive Properties of Carbon Fibers . . . . . . . . . . . . . . . . . . . . . . . . . . . 57 PAN-Based Carbon Fibers . . . . . . . . . . . . . . . . . . 59 Pitch-Based Carbon Fibers . . . . . . . . . . . . . . . . . . 59 Effect of Fiber Shape on Fiber and Composite Properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 61 Effect of Shape on Tensile Strength of Carbon 61 Fibers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . PAN-Based Carbon Fibers . . . . . . . . . . . . . . . . . . 61 Pitch-Based Carbon Fibers . . . . . . . . . . . . . . . . . . 61 Effect of Shape on Compressive Strength for 62 Carbon Fibers . . . . . . . . . . . . . . . . . . . . . . . . . . . . PAN-Based Carbon Fibers . . . . . . . . . . . . . . . . . . 64 Pitch-Based Carbon Fibers . . . . . . . . . . . . . . . . . . 66 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 66 67 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

X

Contents and Subject Index 4. TEXTILE PREFORMS FOR CARBON-CARBON

71 COMPOSITES . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Frank K KO Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 72 Classification of Preforms . . . . . . . . . . . . . . . . . . . . . . 72 Linear Fibrous Assemblies . . . . . . . . . . . . . . . . . . . . . 74 Fabric Preforms . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 75 Structural Geometry of 2-D Fabrics . . . . . . . . . . . . . . 78 Woven Fabrics . . . . . . . . . . . . . . . . . . . . . . . . . . . 78 Knitted Fabrics . . . . . . . . . . . . . . . . . . . . . . . . . . . 78 Braided Fabrics . . . . . . . . . . . . . . . . . . . . . . . . . . 80 Structural Geometry of 3-D Fabrics . . . . . . . . . . . . . . 82 Woven 3-D Fabrics . . . . . . . . . . . . . . . . . . . . . . . 82 Orthogonal Nonwoven Fabrics . . . . . . . . . . . . . . . . 84 Knitted 3-D Fabrics . . . . . . . . . . . . . . . . . . . . . . . 85 3-D Braided Fabrics . . . . . . . . . . . . . . . . . . . . . . . 89 Structure and Properties of Textile-Reinforced CCC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 89 Modeling of Textile Structural Composites . . . . . . . . . 93 Concluding Remarks . . . . . . . . . . . . . . . . . . . . . . . . . . 99 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 100 Bibliography . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 104

.

5 CARBON-CARBON MATRIX MATERIALS

. . . . . . . . . . . . . . 105

N. Murdie. C.P. Ju. J. Don and M.A. Wright Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Carbon Fibers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Fabrication Methods of CC Composites . . . . . . . . . . Liquid Phase Infiltration . . . . . . . . . . . . . . . . . . . . . . Pitch Matrices . . . . . . . . . . . . . . . . . . . . . . . . . . Thermoset Resin Matrices . . . . . . . . . . . . . . . . . . Gas Phase Infiltration Process . . . . . . . . . . . . . . . . . Isothermal Chemical Vapor Deposition . . . . . . . . . Thermal-Gradient Chemical Vapor Deposition . . . Differential Pressure Chemical Vapor Deposition . . . . . . . . . . . . . . . . . . . . . . . . . . . . Matrix Inhibition . . . . . . . . . . . . . . . . . . . . . . . . . . . . Microstructural Characterization Techniques . . . . . . . Optical Microscopy . . . . . . . . . . . . . . . . . . . . . . . . . X-Ray Diffraction . . . . . . . . . . . . . . . . . . . . . . . . . . Scanning Electron Microscopy . . . . . . . . . . . . . . . . . Transmission Electron Microscopy . . . . . . . . . . . . . . Microstructure of CC Matrices . . . . . . . . . . . . . . . . . . Pitch Matrix Composites . . . . . . . . . . . . . . . . . . . . .

106 108 111 112 112 115 118 119 119 120 121 123 123 124 125 126 127 127

xi

Contents and Subject Index

Resin Matrix Composites . . . . . . . . . . . . . . . . . . . . . CVI Matrix Composites . . . . . . . . . . . . . . . . . . . . . . Influence of Matrix on Composite Properties . . . . . . . General Background . . . . . . . . . . . . . . . . . . . . . . . . Elastic Modulus . . . . . . . . . . . . . . . . . . . . . . . . . . . Tensile Strength . . . . . . . . . . . . . . . . . . . . . . . . . . . Matrix Dominated Properties . . . . . . . . . . . . . . . . . . Two-Dimensional Reinforcements . . . . . . . . . . . . . . Three-Dimensional Reinforcements . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

.

OF MULTIDIRECTIONAL CARBON-CARBON COMPOSITE MATERIALS . . . . . . . . . . . . . . . . . . . . . . . . . . John J. Kibler Symbols . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Background . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Characteristics of CC Materials . . . . . . . . . . . . . . . . . Description of Model . . . . . . . . . . . . . . . . . . . . . . . . . Method of Analysis . . . . . . . . . . . . . . . . . . . . . . . . . . Multidirectional Composite Model . . . . . . . . . . . . . . . Degraded Properties Model . . . . . . . . . . . . . . . . . . . . Thin CC Composites . . . . . . . . . . . . . . . . . . . . . . . . . Property Predictions and Data Correlation . . . . . . . . . Effects of Degraded Properties . . . . . . . . . . . . . . . . . Concluding Remarks . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

135 142 149 149 150 151 154 156 156 158

6 MECHANICS

.

7 MANUFACTURING AND DESIGN OF CARBON-CARBON

xii

169 170 170 172 175 176 179 179 181 182 183 188 191 193

COMPOSITES . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 197 Robert L. Burns Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 198 Composite Design . . . . . . . . . . . . . . . . . . . . . . . . . . . 200 Discontinuous Reinforcement . . . . . . . . . . . . . . . . . . 200 Filament Wound CC . . . . . . . . . . . . . . . . . . . . . . . . 201 Laminates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 202 Through-the-Thickness Reinforced CC . . . . . . . . . . 203 Thick-Walled Constructions . . . . . . . . . . . . . . . . . . .204 Design Summary . . . . . . . . . . . . . . . . . . . . . . . . . . 209 Carbon-Carbon Composite Densification . . . . . . . . . 211 212 Laminates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thin-Walled 3-D Composites . . . . . . . . . . . . . . . . . 213 Thick-Walled 3-D Composites . . . . . . . . . . . . . . . . . 214 Composite Properties . . . . . . . . . . . . . . . . . . . . . . . 215

Contents and Subject Index

References

................................

221

8. HIGH-TEMPERATURE COATINGS ON CARBON FIBERS

AND CARBON-CARBON COMPOSITES . . . . . . . . . . . . . . . 223 James E. Sheehan Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 225 High-Temperature Coatings on Carbon Fibers . . . . . 226 Carbon Fibers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 226 Fiber Coating Methods . . . . . . . . . . . . . . . . . . . . . . 229 Chemical Vapor Deposition . . . . . . . . . . . . . . . . . 230 Physical Vapor Deposition . . . . . . . . . . . . . . . . . . 231 Electroplating . . . . . . . . . . . . . . . . . . . . . . . . . . . 234 Liquid Precursor Methods . . . . . . . . . . . . . . . . . .235 Liquid Metal Transfer . . . . . . . . . . . . . . . . . . . . . 236 Coated Fiber Properties . . . . . . . . . . . . . . . . . . . . . . 237 High-Temperature Coatings on CC Composites . . . . 239 Carbon-Carbon Composites . . . . . . . . . . . . . . . . . . 239 External Coatings . . . . . . . . . . . . . . . . . . . . . . . . . . 241 Space Shuttle Orbiter Thermal Protection . . . . . . . 241 Structural Applications . . . . . . . . . . . . . . . . . . . . 242 Performance Issues . . . . . . . . . . . . . . . . . . . . . . 245 Internal Coatings . . . . . . . . . . . . . . . . . . . . . . . . . . . 249 Liquid Precursor Method . . . . . . . . . . . . . . . . . . . 250 Matrix Chemical Modifications . . . . . . . . . . . . . . . 250 Chemical Vapor Deposition . . . . . . . . . . . . . . . . . 251 Advantages and Limitations . . . . . . . . . . . . . . . . . 252 Temperature Limitations . . . . . . . . . . . . . . . . . . . . . 253 Current Coatings . . . . . . . . . . . . . . . . . . . . . . . . 253 Ultrahigh-Temperature Coatings . . . . . . . . . . . . . 255 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 256

9. APPLICATIONS OF CARBON-CARBON

Louis Rubin References

. . . . . . . . . . . . . . . 267

................................

280

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Chapter 1 Carbon-Carbon Overview* John D. Buckleyt NASA Langley Research Center Hampton, Virginia

Introduction

1

Carbon Fibers 4 Carbon Fibers in Carbon Matrix

5

Discontinuous Fiber Composites 5 Continuous Fiber Composites 6 Chemical Vapor Deposition 7 Carbonized Organic Composites 9 Mechanical Properties Applications

12

Conclusions

14

References Bibliography

11

14

17

Introduction Carbon-carbon (CC) materials are a generic class of composites similar to the griiphite/cpoxy family of polymer matrix composites. These materials can be made in a wide variety of forms, from one-dimensional to n-dimensional, using unidirectional tows, tapes, or woven cloth (fig. 1). Because of their multiformity, their mechanical properties can be readily tailored. Carbon materials have high strength and stiffness potential as well as high thermal and chemical stability in *Similar version published in Ceramic Bulletin, vol. 67, no. 2, 1988, (QACerSL.

t Member, the American Ceramic Society. 1

CC Materaals and Composates

/?AL( ---

1-D

2-D

3-D

n-D

General properties of carbon-carbon composites Ultimate tensile strength >MPa 5 Thermal conductivity ~ 1 1 . W/(m.K) (>40 000 psi) Linear thermal expansion Modulus of elasticity >69 GPa =1.1 x 10-60/C (>lo7psi) Density <2990 kg/m3 Melting point >41OO0C Figure 1. Multiformity and general properties of carbon-fiber and carbonmatrix composites.

inert environments. These materials must, however, be protected with coatings and/or surface sealants when used in an oxidizing environment. The development of CC materials began in 1958 and was nurtured under the U.S. Air Force space plane program, Dyna-Soar, and NASA’s Apollo projects. It was not until the Space Shuttle Program that CC material systems were intensively researched. The criteria that led to the selection of CC composites as a thermal protection system were based on the following requirements: (1) maintenance of reproducible strength levels at 165OoC, ( 2 ) sufficient stiffness to resist flight loads and large thermal gradients, ( 3 ) low coefficient of thermal expansion to minimize induced thermal stresses, (4) oxidation resistance sufficient to limit strength reduction, ( 5 ) tolerance to impact damage, and (6) manufacturing processes within the state of the art. Carbon-carbon composites consist of a fibrous carbon substrate in a carbonaceous matrix. Although both constituents are the same element, this fact does not simplify composite behavior because the state of each constituent may range from carbon to graphite. Crystallographic carbon, namely graphite, consists of tightly bonded, hexagonally arranged carbon layers that are held together by weak van der Waals forces. The single crystal graphite structure is illustrated in figure 2 (ref. 1). The atoms within the layer plane or basal plane (a-b direction) have a covalent bond strength of ~ 5 2 kJ/mol 4 (ref. 2), while the bonding energy between basal planes ( e direction) is ~ 5 7kJ/mol (ref. 3). The result is a crystal that is remarkable in its anisotropy, being almost isotropic within the basal plane but with c direction properties that differ by orders of magnitude. On a larger scale, carbon,

2

Carbon- Carbon Overview in addition to its two well-defined allotropic forms (diamond and graphite), can take any number of quasicrystalline forms ranging continuously from turbostratic (amorphous, glassy carbon) to a highly crystalline graphite (fig. 3).

C

Ea

Reference directions

Figure 2. Tightly bonded, hexagonally arranged carbon layers (ref. 1) held together by weak van der Waals forces.

1do02

d0O2 2 3.440 8,

(a)

LcS508,

I-L,---i do02 = 3.354 A

(b)

Lc 2 300 8,

Figure 3. Comparison of ( a ) carbon turbostratic structure with ( b ) 3-0 graphite lattice (ref. 1).

The anisotropy of the graphite single crystal encompasses many structural forms of carbon. It ranges in the degree of preferred orientation of the crystallites and influences the porosity, among other variables. A broad range of properties is the result of this anisotropy, which is available in carbon material. In CC composites, this range of properties can extend to both constituents. Coupled with a variety of

3

CC Materials and Composites processing techniques that can be used in the fabrication of CC composites, great flexibility exists in the design of and the resultant properties to be obtained from CC composites. The wide range of properties of carbon materials can be shown when comparing the tensile moduli of commercially manufactured carbon fibers that range from 27.6 GPa (4 x lo6 psi) to 690 GPa (100 x lo6 psi). In fabrication, the fibers can be used in either continuous or discontinuous form. The directionality of the filaments can be varied ranging from unidirectional lay-ups to multidirectional weaves. The fiber volume used constitutes another variable. The higher the volume fraction of a specific high-strength fiber in a matrix, the greater the strength of the composite. The matrix can be formed via two basic approaches: (1) through the carbonization of an organic solid or liquid, such as a resin or pitch, and (2) through the chemical vapor deposition (CVD) of carbon from a hydrocarbon. A range of carbon structures can be obtained by either approach. Finally, heat treatment of the composite material at graphitization temperatures offers additional variability to the properties that can be obtained. Typically, there is an optimum graphitization temperature at which the highest strength can be obtained for a given composite composition of fiber and matrix (refs. 4 and 5).

Carbon Fibers The properties of carbon fibers can vary over a wide range depending on the organic precursor and processing conditions used. At present, graphite fibers are produced from three precursor materials: rayon, polyacrylonitrile (PAN), and petroleum pitch. Fibers having a low modulus (27.6 GPa (4 x lo6 psi)) are formed using a rayon precursor material that may be chemically pretreated by a sequence of heating steps. First, the fiber is heated to 3 400'C to allow cellulose to pyrolyze.$ Carbonization3 is completed more rapidly at > 1000'C. Upon completion of carbonization, the fiber is graphitized11 by heating to >2000'C; the fiber is now, for all practical purposes, 100 percent carbon. High-modulus carbon fibers from rayon precursors are obtained by the additional process of stretching the carbon fibers at the final heat treatment temperature. High-modulus (344 GPa (50 x lo6 psi)), high-strength (2.07 GPa (300 x lo3 psi)) carbon fibers are typically made from PAN or, in some cases, mesophase pitch precursors. These fibers are processed similarly in a three-stage operation (fig. 4, ref. 6). The PAN fibers are initially stretched from 500 percent to 1300 percent and then stabilized (crosslinked) in an oxygen atmosphere at 200'C to 280'C (under tension). Carbonization $Decomposition or chemical change by thermal conversion of organic materials to carbon and graphite. §Continued heating of organic material to >lOOO°C produced by pyrolysis.

to initiate ordering of the carbon structures

IlContinued heating of carbonized organic materials to the 2000'C 100-percent graphite-ordered crystal structure.

4

to 30OO0C range to produce a

Carbon-Carbon Overview of the fibers is conducted between 1000°C and 160OOC. Finally, graphitization is accomplished at >25OO0C. Mesophase pitch fibers undergo the same processing procedure as PAN fibers but do not require an expensive stretching process during heat treatment to maintain preferred alignment of crystallites (fig. 4, ref. 6). Control of fiber shape has resulted in improved fiber strength (4.1 GPa (600000 psi)), see ref. 7, when produced from melt-spun, mesophase petroleum pitch (fig. 5, ref. 7). Round fibers using the same method had a strength of 2.1 GPa (300 x lo3 psi), as shown in reference 4. Of the shapes studied, the c-shape and hollow fibers were found to be superior in strength to round solid and trilobal cross sections (refs. 4 and 7).

Spool

EPOXYsizing

Surface treatment

Figure 4. Carbon fiber production using PAN and pitch processes (ref. 6).

Carbon Fibers in Carbon Matrix Addition of a matrix to carbon fiber, either through the carbonization of an organic precursor or by the deposition of pyrolytic carbon, is conducted at 8OO0C to 150OOC. Subsequent heat treatment of the composite material may involve temperatures to 30OO0C.

Discontinuous Fiber Composites Fabrication of discontinuous fiber composites uses short carbon fibers combined with either a pyrolytic carbon or pyrolyzed organic matrix. This approach to CC composites generally does not have true fiber reinforcement as an objective. Rather, discontinuous fiber substrates have been used to: (1) increase fabrication capability of large-scale structures, ( 2 ) achieve a more nearly isotropic material, (3) increase the composite interlaminar tensile strength, and (4) along with continuous filament substrates, obtain a stronger composite by providing additional nucleation sites that serve to reduce composite porosity.

5

CC Materials and Composites

Cartridge housing

e Ited-pitch precursor

Melt-pressure indicat

Figure 5 . Melt spinning apparatus used to produce noncircular carbon fibers (ref. 7). The fabrication techniques most widely applied are a carbonized, rayon felt substrate with a pyrolytic carbon matrix, and short, chopped fibers in a pitchbased matrix. Felt is produced through the mechanical carding of viscous rayon fibers to produce a continuous web of fibers. The webs are folded one on top of another to produce a batt. The batts are then cut, stacked, and needled to produce the required felt. The rayon felt is subjected to a controlled carbonization cycle in an inert atmosphere or vacuum; the maximum temperature determines such factors as shrinkage, weight loss, and chemical composition of the felt. A maximum carbonization temperature of 12OO0C is a nominal standard the length of the carbonization cycle and rate of temperature rise are dictated by the thickness of the felt. Carbon content in the fibers is ~ 9 percent. 8 Carbon-carbon composites have also been fabricated from short carbon fibers using isotropic casting, flocking lay-up, spray lay-up, and pulp-molding techniques (fig. 6, refs. 8 to 10). The rationale for using these short fibers is to reduce composite properties of anisotropy, specifically, the effect that relatively long fibers used in other discontinuous fiber substrates produce fiber alignment during processing resulting in anisotropic composite properties (ref. 9).

Continuous Fiber Composites Continuous filament substrates use either the properties of high-strength filaments or achieve a high degree of preferred orientation on the macroscale of the matrix. The fabrication complexity involving continuous-filament substrates is

6

Carbon- Carbon Overvzew

Figure 6. Models of fiber arrangements for four short-fiber fabrication techniques: ( a )Pocking lay-up; (b) pulp molding; ( c ) isotropic casting; and (d) spray lay-up (ref. 5 ) . determined by two parameters: the directionality of the filaments and the amount of layer interlocking achieved in the substrate. The plies and filament winding of unidirectional tapes can be used to achieve a highly oriented substrate, usually with no interlocking between layers. Woven fabrics are used to form a two-dimensional laminate with no interlocking between layers. Helical filament winding, which is directional, results in continuous, adjacent layer interlocking. Multilayer locking is achieved through complex weaving patterns or yarn placement resulting in “multidirectional” substrates (fig. 7).

Chemical Vapor Deposition The deposition of carbon on the filament substructures just discussed is accomplished either by pyrolyzing an organic matrix or through CVD. The CVD of carbon from a hydrocarbon gas within a substrate is a complex process. Various techniques have been applied to infiltrate various fiber substrates including isothermal thermal gradient (ref. 1l), pressure gradient (ref. 12), and pressure pulsation (ref. 13). The first two have been the most extensively used. The isothermal technique is illustrated in figure 8. The substrate is radiantly heated by an inductively heated susceptor so that the gas and substrate are maintained at a uniform temperature. Infiltration is normally accomplished at llOO°C and at

7

CC Materials and Composites

Figure 7. Interlocking approaches of continuous filament substrates: (a) tape wrapped, shingle; ( b ) filament wound, helix; and (c) multidimensional. reduced pressures (6 kPa (50 torr)) with the flow rates primarily determined by the substrate surface area. This technique produces a crust on the outer surfaces of the substrate, thus requiring machining and multiple infiltration cycles. In the thermal gradient technique (fig. 9), the part to be infiltrated is supported by a mandrel that is inductively heated. Therefore, the hottest portion of the substrate is the inside surface, which is in direct contact with the mandrel. The outer surface of the low-density substrate is exposed to a cooler environment and results in a temperature gradient through the substrate thickness. Surface crusting is eliminated because the deposition rate is greater on the heated fibers near the mandrel, whereas the cooler outer fibers receive little or no deposit. Under proper infiltration conditions, the carbon is first deposited on the inside surface and, in a continuous process, progresses radially through the substrate as the densified substrate itself becomes inductively heated. Infiltration is normally accomplished at atmospheric pressure with a mandrel heated to = l 100°C (ref. 14).

Carbon-Carbon Overview

gas

Carrier gas

Original iiber substrate

Figure 8. Isothermal chemical vapor deposition to infiltrate fibrous carbon substrate.

Carbonized Organic Composites Carbonized organic composites have fabrication procedures that are similar to those of conventional fiber-reinforced, resin-laminating techniques. The starting material is usually a prepregged# fabric or yarn. These precursor materials are staged nominally at =lOO°C to achieve the desired degree of tack and flow of the resin. A laminate is then constructed and cured under pressure to compact the stack-impregnated fabric. Curing temperatures normally range from 125OC to 175OC with curing pressures on the order of 2.76 MPa (400 psi). The reinforced resin laminate is then postcured at 2OOOC to 275OC. As pyrolysis is initiated, shrinking occurs as the organic phase decomposes. Simultaneously, the release of vapors from pyrolysis expands the composite material. A slow release of these volatile by-products is required to minimize structural damage to the char. Finally, as higher temperatures are reached, thermal expansion of the carbon char itself occurs after pyrolysis is complete. After the initial carbonization, the material is then subjected to a series of reimpregnation and carbonization cycles until the desired density or the maximum density is achieved. The reimpregnation process is usually conducted under vacuum and pressure to aid in maximizing the pore filling. # A fabric impregnated with a matrix material in a tacky state.

9

CC Materials and Composites

Carbon substrate Induction coils Jacket Sleeve

Graphite susceptor

gas

gas

Figure 9. Thermal gradient chemical vapor deposition. If graphitization is desired, the high-temperature heat treatment may be used after each carbonization step or at the end of the reimpregnation and recarbonization cycles.

To summarize, a typical manufacturing cycle of a 2-D CC part is shown in figure 10. First, a woven graphite fabric that is preimpregnated with phenolic resin is laid-up as a phenolic-graphite laminate in a mold and is autoclave-cured. Once cured, the part is pyrolyzed to form a carbon matrix surrounding the graphite fibers. The part is then densified by multiple furfural alcohol reimpregnations and pyrolyzations. The resulting CC part then is ready for use in inert or vacuum environments. This process is very time-consuming. A single pyrolysis may take >70 hr in a low-temperature, inert-atmosphere furnace. Although CC materials can withstand temperatures >30OO0C in a vacuum or in an inert atmosphere, they oxidize and sublime when in an oxygen atmosphere at 60OoC. To allow for use of CC parts in an oxidizing atmosphere, they must be compounded with materials that produce oxidation-protective coatings through thermochemical reaction with oxygen at >20OO0C (ref. 15) or they must be coated and sealed to protect them (ref. 16). For applications such as the Space Shuttle

10

Cloth prepreg

Cut, lay-up, I d e b u l k & bag-

.-

25

-

High-strengthcarbon-carbon

k- 160 g$ cD=

g 2 - 1 2 0 6.;

(Qz

%!B -80 ~2 3

Carbon-Carbon Oueruaew turbine engine applications using CC composites include exhaust nozzle flaps and seals, augmenters, combustors, and acoustic panels.

Carbon matrix

Graphite fiber

0

Higher temperature performance without cooling

0

Low weight Potential low cost

Carbon-carbon microstructure

Nonstrategic materials

Figure 12. One-piece, bladed, carbon-carbon turbine rotor (ref 27).

Carbon-carbon material systems using coatings, TEOS, and additions to the basic CC recipe have improved the oxidation resistance of products made of CC composites by an order of magnitude. The ACC composites are being used in products such as the nozzle in the F-100 jet engine afterburner, turbine wheels operating at >40000 rpm, nonwetting crucibles for molten metals, nose caps and leading edges for missiles and for the Space Shuttle, wind-tunnel models, and racing car and commercial disk brakes (ref. 28). Pushing the state of the art in CC composites is the piston for internal combustion engines (refs. 27 and 29). The CC piston would perform the same way as any piston in a reciprocating internal combustion engine while reducing weight and increasing the mechanical and thermal efficiencies of the engine. The CC piston concept features a low piston-to-cylinder wall clearance; this clearance is so low, in fact, that piston rings and skirts are unnecessary. These advantages are made possible by the negligible coefficient of thermal expansion of this kind of

13

CC Materials and Composites CC (0.54 x lop6 cm/cm/OC (0.3 x lop6 in./in./°F)).## Carbon-carbon material maintains its strength at elevated temperatures allowing the piston to operate at higher temperatures and pressures than those of a comparable metal piston. The high emittance and low thermal conductivity of the CC piston should improve the thermal efficiency of the engine because less heat energy is lost to the piston and cooling system. The elimination of rings reduces friction, thus improving mechanical efficiency. Besides being lighter than conventional pistons, the CC piston can produce cascading effects that could reduce the weight of other reciprocating components such as the crankshaft, connecting rods, flywheels, and balances, thus improving specific engine performance (ref. 29).

Conclusions Carbon-carbon composites offer a unique combination of properties. In nonoxidizing environments, they retain room temperature mechanical properties at >2225OC. For applications in oxidizing environments, current coatings limit maximum use temperatures to M1600OC. High thermal conductivity and low thermal expansion of carbon-carbon composites make them excellent candidates for applications involving thermal shock. Because of the variety of fibers, weaving patterns, and lay-up procedures that can be used for carbon-carbon composites, their mechanical properties can be tailored over a wide range to fit the application. Continuing research on carbon-carbon materials in the United States emphasizes an understanding of material behavior. Of particular importance to both researchers and fabrication personnel are methods of improving matrix properties (particularly in-plane shear and out-of-plane tensile strengths) and improving oxidation-resistant coatings with higher use temperatures, longer lifetimes, and less costly fabrication methods.

References 1. Bokros, J. C.: Deposition, Structure, and Properties of Pyrolytic Carbon. Chemistry and Physics 6; Carbon-A Series of Advances, Volume 5 , Philip L. Walker, Jr., Marcel Dekker, Inc., 1969, pp. 1-118. 2. Kanter, Manuel A.: Diffusion of Carbon Atoms in Natural Graphite Crystals. Phys. Review, vol. 107, no. 3 , Aug. 1, 1957, pp. 655-663. ##Carbon-carbon composites can have a range of thermal expansion coefficients, depending on the processing techniques.

14

Carbon-Carbon Overview 3. Dienes, G. J.: Mechanism for Self-Diffusion in Graphite. J . Appl. Phys., vol. 23, no. 11, Nov. 1952, pp. 1194-1200. 4. Edie, D. D.; Fox, N. K.; Barnett, B. C.; and Fain, C. C.: Melt-Spun NonCircular Carbon Fibers. Carbon, vol. 24, no. 4, 1986, pp. 477-482. 5. Stoller, H. M.; Butler, B. L.; Theis, J. D.; and Lieberman, M. L.: Carbon Fiber Reinforced-Carbon Matrix Composites. Composites: State of the Art, J. W. Weeton and E. Scala, eds., Metallurgical SOC.of the American Inst. of Mining, Metallurgical and Petroleum Engineers, Inc., c. 1974, pp. 69-1 36. 6. Diefendorf, R. J.: CarbodGraphite Fibers. Engineered Materials Handbook. Volume 14omposites, ASM International, 1987, pp. 49-53. 7. Cogburn, John W.; Fain, C. C.; Edie, D. D.; and Leigh, H. D.: Processing C-Shape Pitch-Based Carbon Fibers. Metal Matrix, Carbon, and Ceramic Matrix Composites-1987, John D. Buckley, ed., NASA CP-2482, 1987, pp. 185-200. 8. Cook, J. L.; Lambdin, F.; and Trent, P. E.: Discontinuous Carbon/Carbon Composite Fabrication. Carbon Composite TechnologpWith Special Emphasis on CarbonlCarbon Systems, Proceedings of the 10th Annual Symposium of the New Mexico Section of ASME and University of New Mexico, Jan. 1970, pp. 143-171. 9. Lambdin, F.; Cook, J. L.; and Marrow, G. B.: Fiber-Reinforced Graphite Composite Fabrication and Evaluation. Doc. Y-1684, TID-4500 (Contract W-7405-eng-26), Nuclear Div., Union Carbide Corp., Sept. 4, 1969.

10. Lambdin, F.; and Cook, J. L.: Fabrication of Carbon-Carbon Composites Using Electrostatic Fiber Deposition (Flocking). Y-1786 (Contract No. W-7405-eng26), Y-12 Plant, Union Carbide Corp., June 1971. 11. Pierson, H. 0.: Development and Properties of Pyrolytic Carbon Felt Composites. Advanced Techniques for Material Investigation and Fabrication, Volume 14 of National Symposium and Exhibit, SOC. of Aerospace Material and Process Engineers, 1968, Paper 11-4B-2. 12. Kotlensky, W. V.; and Pappis, J.: Mechanical Properties of CVD Infiltrated Composites. Proceedings of 9th Biennial Conference on Carbon. Defense Ceramic Information Center, Compilers, 1969, pp. 76-80.

15

CC Materials and Composites 13. Beatty, R. T.; and Kipplinger, D. V.: Gas Pulse Impregnation of Graphite With Carbon. Nucl. Appl. & Technol., vol. 8, no. 6, June 1970, pp. 488495. 14. Theis, J. D., Jr.; Taylor, A. J.; Rayner, R. M.; and Frye, E. R.: Filament-Wound CarbonlCarbon Heatshield SC-I IFW-Y12-7, A Process History. SC-DR-70425, Sandia Labs., Dec. 1970. 15. Buckley, John D.: Static, Subsonic, and Supersonic Oxidation of JT Graphite Composites. NASA TN D-4231, 1967. 16. Strife, James R.; and Sheehan, James E.: Ceramic Coatings for Carbon-Carbon Composites. Ceramic Bulletin, vol. 67, no. 2, 1988, pp. 369-374. 17. Buch, J. D.: Graphite Crystals-A General Model for Diverse Carbon Forms. Metal Matrix, Carbon, and Ceramic Matrix Composites, John D. Buckley, ed., NASA CP-2357, 1984, pp. 119-135. 18. Rummler, D. R.; and Sawyer, J. W.: Properties and Potential of Advanced Carbon-Carbon for Space Structures. Metal Matrix, Carbon, and Ceramic Matrix Composites, John D. Buckley, ed., NASA CP-2357, 1984, pp. 149-170. 19. Ransone, Philip 0.;and Ohlhorst, Craig W.: Interlaminar Shear and Out-ofPlane Tensile Properties of Thin 3-D Carbon-Carbon. Metal Matrix, Carbon, and Ceramic Matrix Composites, John D. Buckley, ed., NASA CP-2357, 1984, pp. 137-148. 20. Webb, Richard D.: Oxidation-Resistant Carbon-Carbon Materials. Metal Matrix, Carbon, and Ceramic Matrix Composites-1985, John D. Buckley, ed., NASA CP-2406, 1985, pp. 149-162. 21. Gray, Paul E.; and Engle, Glen B.: Wettability of CarbonKarbon Composites and Carbon Fibers by Glass Sealants Used in Oxidation Inhibition. Metal Matrix, Carbon, and Ceramic Matrix Composites-1985, John D. Buckley, ed., NASA CP-2406, 1985, pp. 163-174. 22. Johnson, A. C.; and Finley, J. W.: CarbonKarbon Composites for Advanced Spacecraft. Metal Matrix, Carbon, and Ceramic Matrix Composites-I 985, John D. Buckley, ed., NASA CP-2406, 1985, pp. 175-190.

16

Carbon-Carbon Overview 23. Sawyer, J. W.; and Moses, P. L.: Effect of Holes and Impact Damage on Tensile Strength of Two-Dimensional Carbon-Carbon Composites. Metal Matrix, Carbon, and Ceramic Matrix Composites-1985, John D. Buckley, ed., NASA CP-2406, 1985, pp. 245-260. 24. Maahs, Howard G.; and Ransone, Philip 0.: Mechanical Property Evaluation of 2-D Carbon-Carbon Panels Fabricated From a Specialty-Weave Fabric. Metal Matrix, Carbon, and Ceramic Matrix Composites-1985, John D. Buckley, ed., NASA CP-2406, 1985, pp. 261-276. 25. Ohlhorst, Craig W.; and Ransone, Philip 0.: Effects of Thermal Cycling on Thermal Expansion and Mechanical Properties of Advanced Carbon-Carbon Composites. Metal Matrix, Carbon, and Ceramic Matrix Composites-1985, John D. Buckley, ed., NASA CP-2406, 1985, pp. 277-288. 26. Ransone, Philip 0.; and Maahs, Howard G.: Effect of Processing on Microstructure and Mechanical Properties of 3-D Carbon-Carbon. Metal Matrix, Carbon, and Ceramic Matrix Composites-1985, John D. Buckley, ed., NASA CP-2406, 1985, pp. 289-303. 27. Miller, T. J.; and Grimes, H. H.: Research on Ultra-High-Temperature Materials-Monolithic Ceramics, Ceramic Matrix Composites, and Carbon/Carbon Composites. Advanced Materials Technology, Charles P. Blankenship and Louis A. Teichman, compilers, NASA CP-2251, 1982, pp. 275-291. 28. Klein, J.: Carbon-Carbon Composites, Advanced Mater. & Process., vol. 130, no. 5, 1986, pp. 64-68. 29. Taylor, Allan H.: Carbon-Carbon Pistons for Internal Combustion Engines. NASA Tech Briefs, vol. 9, no. 4, Winter 1985, pp. 156-157.

Bibliography A. Becker, Paul R.: Leading-Edge Structural Material System of the Space Shuttle. American Ceram. SOC.Bull., vol. 60, no. 11, Nov. 1981, pp. 1210-1214. B. Rummler, Donald R.: Recent Advances in Carbon-Carbon Materials Systems. Advanced Materials Technology, Charles P. Blankenship and Louis A. Teichman, compilers, NASA CP-2251, 1982, pp. 293-312.

17

Chapter 2 Carbon Fiber Manufacturing D. D. Edie and R. J. Diefendorf Clemson University Clemson, South Carolina

Introduction 20 Manufacture of PAN-Based Carbon Fibers

20

Solution Spinning of PAN Precursor Fibers

20

Melt-Assisted Spinning of PAN Precursor Fibers Heat Treatment of PAN Precursor Fibers Oxidation of PAN Precursor Fibers

22

23

23

Carbonization and Graphitization 24 Manufacture of Rayon-Based Carbon Fibers Manufacture of Pitch-Based Carbon Fibers Mesophase Pitch

26 27

28

Melt-Spinning Mesophase Precursor Fibers

29

Heat Treatment of Mesophase Precursor Fibers Oxidation of Mesophase Precursor Fibers

31

31

Carbonization and Graphitization 3 1 Isotropic Pitch-Based Carbon Fibers

33

Manufacture of Vapor-Grown Carbon Fibers Mechanical Properties of Carbon Fibers

33

35

Summary 37 Acknowledgments 37 References 37

19

CC Materials and Composites

Introduction The high strength, superior stiffness, and light weight of carbon fibers have made them the dominant reinforcing fibers used in high-performance polymer matrix composites. However, the same fibers can also reinforce brittle materials, such as ceramics and carbon, thus creating a unique class of high-temperature composite materials. When properly protected from oxidation, these carbon fibercarbon matrix composites can withstand extended exposure to temperatures of up to 25OO0C, making them attractive for many aerospace applications. In addition, because of their improved friction performance and high wear resistance, carboncarbon (CC) materials are used in high-performance brakes of aircraft and racing cars. Using these CC brakes in passenger cars and trucks is currently being evaluated. At present, all commercial carbon fibers are produced by the thermal decomposition of various organic fiber precursors. The most popular precursor materials are fibers of polyacrylonitrile (PAN), cellulose (rayon), and pitch (ref. 1). A proposed alternate process, which produces a discontinuous, high-performance carbon fiber, is called vapor-growth. This chapter describes the similarities of these four fiber processes and discusses their differences.

Manufacture of PAN-Based Carbon Fibers Today, approximately 90 percent of all commercial carbon fibers are produced from a PAN precursor fiber. Normally, PAN is copolymerized with a small amount of another monomer, such as methylacrylate, to lower its glass transition temperature and control its oxidation rate. Figure 1 lists a few of the many monomers copolymerized with acrylonitrile to produce commercial PAN precursor fibers. Typically, the precursor fiber would contain 93 to 95 percent acrylonitrile units, with the remainder being one or more of these monomers. Because PAN decomposes below its melt temperature, it is normally extruded into filament form using various solution spinning techniques.

Solution Spinning of PAN Precursor Fibers Figure 2 shows the process schematic of a typical solution spinning process (ref. 2). In this process, the copolymer first is dissolved in a suitable solvent, such as dimethylacetamide, and loaded into a storage tank. Typically, the spinning solution is quite concentrated (from 15 to 30 percent polymer by weight). The solution is pumped through a die head, where it is filtered to remove impurities before being extruded through a spinnerette containing a large number of small (approximately 100-pm) capillary holes. In the process shown in figure 2, the solution immediately enters a coagulating bath as it exits the capillary. This is termed wet spinning.

20

-

Monomer Acrylic acid

Structure H\ c=c: H

,

C=O

H’

CH

Itaconic acid

H,

Y c=c:

H’

C=O CH2

C=O M

Methacrylic acid

H,

7

c=c:

C=O

CH3

H‘

H\ c=c: H

Methyl acrylate

C=O

H’

1

? w3

H\

Vinyl acetate

c=c:

H’

H

0

c=o w3

Acrylonitrile

H\ c=c: H H’

CN

C C Materials and Composites of the fiber center. The large concentration gradient across the fiber cross section makes the initial density of the center portion of a rapidly formed, wet-spun fiber much less than that of the fiber skin. As the solvent eventually diffuses out of the fiber center, the density of the inner portion of the fiber increases, causing the fiber skin to collapse, yielding a kidney bean (or dog-bone-shape) fiber. However, wet spinning can produce PAN precursor fibers with a circular cross section and a minimum of internal voids if the rate of solvent extraction is properly controlled.

In another variation of solution spinning, the polymer solution is extruded into a hot gas environment. In this case, the temperature and composition of the gas must be carefully monitored to control the rate of solvent evaporation and, thus, the structure of the fiber. This process variation, called dry spinning, produces an as-spun fiber with a dog-bone-shaped cross section. Often, both wet- and dry-spun fibers are washed after fiber formation to remove the final traces of solvent. Then the fibers are passed through one or two stages in which they are stretched to further align the polymer molecules parallel to the fiber axis. Finally, this fully drawn PAN precursor fiber is dried and packaged.

Melt-Assisted Spinning of PAN Precursor Fibers Solution-spun PAN fibers can be converted to carbon fibers with excellent mechanical properties. However, large amounts of solvent are required for solution spinning, and ultimately this solvent must be completely removed from the fiber and recovered. This process adds to the cost of solution-spun precursor fibers, and the trace impurities ultimately can limit the properties of the final carbon fiber. To overcome some of these problems, a melt-assisted process for producing PAN precursor fibers has been developed by BASF Structural Materials, Inc. (ref. 3 ) . In this process, the acrylonitrile copolymer is polymerized in an aqueous suspension. After polymerization, the copolymer is purified and dewatered before extrusion. The PAN copolymer then is pelletized and fed to an extruder. Excess water effectively plasticizes the polymer, allowing it to form a homogeneous melt well below its degradation temperature. Figure 3 shows a flow diagram of this novel fiber-forming process, termed melt-assisted spinning. In melt-assisted spinning, the plasticized PAN copolymer is extruded through a multiple-hole spinnerette directly into a steam-pressurized solidification zone. After passing through this steam environment, the fiber is stretched and dried in a series of steps similar to those found in solution spinning processes. The meltassisted process offers several advantages over conventional solution spinning, including completely eliminating the need for expensive solvents and reducing waste water treatment requirements. In addition, because the polymer content of the plasticized PAN is much higher than that of the solutions used in wet or dry spinning, coalescence during fiber formation is simplified. Thus, the cross-sectional structure of fibers formed by melt-assisted spinning should be more uniform.

22

Cata'ysts

* Recovered monomers

Slurry

-.

T G tz

Unreacted

Washing/ -dewatering-

Compounding/ pelletizing

Plasticating extrusion

CC Matemals and Composites

rn Polyacrylonitrile

C\lf/C\

'"1

F

II

/C\

'i

i

:,i

I'

/ \ /cnc'

7

,,//cc\,/c\~/E~N/c~~

'i

I

0 2 , -220-280°C

I 1

I

"\c/l,\c/f& H,

iI

I i H,

l l c / N \ c / N \ ~ N \ & N ~ c 4

l

l

l

6Al\c/E\p\ 1 I4

l

AC\ )a

l E./H%

I1

+ HCN + C02 + H20 Figure 4. Stabilization PAN precursor summarizing the most frequently observed functional groups (ref. 5 ) .

Tension control Exhaust gases

idized fiber wind-up

Tension control

Figure 5 . Continuous process for oxidizing PAN precursor fibers (ref. 6). Carbonization and Graphitization After being stabilized, the fiber is finally carbonized and sometimes graphitized by slowly heating it in an inert atmosphere to temperatures ranging from 1000°C to 2800°C. By definition, carbonization implies heat treatment at temperatures of 17OOOC or less, whereas graphitization means heat treating to higher temperatures (often approaching 30OO0C).During this final heat treatment, almost all noncarbon

24

Carbon Fiber Manufacturing elements are' driven from the fiber. In fact, the carbon content of the final fiber can range from 80 percent to in excess of 99 percent, depending upon the final carbonization temperature. Often, a carbon resistance furnace, similar to that shown in figure 6, is used for this process step. Thus, the inert atmosphere not only prevents oxygen from pitting the fiber at these high temperatures, but also it protects the carbon heating elements of the furnace from oxygen attack.

Tension control

Figure 6. Schematic of carbon resistance furnace used to continuously carbonize stabilized precursor fiber. Because gases such as CH4, H20, NH3, N2, HCN, C02, and CO are evolved at a rapid rate as the stabilized precursor fiber is heated to 1000"C, relatively slow heating rates are used initially (ref. 7). However, above 1000"C, only smaller molecules such as H2 and N2 are given off. Thus, carbonization often is conducted in two steps: precarbonization (heat treatment up to 1100°C) and carbonization (heat treatment at temperatures ranging from 1600°C to 1800°C). Even though the carbon content of PAN is 54 percent, carbon loss during the heat treatment steps makes the overall yield for converting PAN precursor fiber to carbon fiber approximately 40 to 45 percent (ref. 7). Both the final heat treatment temperature and the degree of molecular orientation of the molecules in the thermoset precursor fiber govern the modulus of the final carbon fiber product. As in all brittle materials, structural flaws limit the strength of the final carbon fiber. Thus, the purity in the precursor fiber, the final carbonization conditions, and even the void content of the precursor fiber can influence the strength of the final carbon fiber. After final heat treatment, most PAN-based carbon fibers are given a surface treatment to improve their bonding with polymeric matrix materials. Although surface treatment results in some roughening of the surface, its primary effect is to increase the concentration of oxygenated groups on the fiber surface. Various

25

CC Materials and Composites techniques can accomplish this: exposing the carbon fiber to gases (such as air or carbon dioxide) at elevated temperatures, submerging the fiber in sodium hypochlorite or nitric acid solutions, or electrolytically etching the fiber. The principal goal of this process step is to increase the interfacial bond strength between the fiber and the matrix material and, thus, improve the interlaminar shear strength of the composite. After being surface-treated, a small amount of size (approximately 1 percent by weight) is added to improve the wettability of the carbon fiber. Normally, this size is a low molecular weight form of the anticipated matrix polymer. In other words, epoxy-sized fiber is coated with a low molecular weight epoxy. It should be mentioned that the primary objective of this size is to improve the wettability of the fiber, not to improve its handleability. Therefore, even sized fibers can be difficult to handle during preform weaving and composite fabrication.

Manufacture of Rayon-Based Carbon Fibers Rayon (or cellulose) precursor fibers were pyrolyzed to form the first highstrength carbon fibers. However, currently less than 1 percent of all carbon fibers are produced this way. The molecular structure of cellulose, a naturally occurring polymer found in wood pulp and cotton, is shown in figure 7. A wet-spinning process produces these cellulose precursor fibers. To form the solution needed for wet spinning, raw cellulose is dissolved in a basic solution and then treated with CS2 to form cellulose xanthate. This soluble derivative of cellulose then is dissolved in NaOH and extruded through a spinnerette into a coagulation bath containing 10 to 15 percent sulfuric acid. As the cellulose xanthate enters the acidic bath, it is hydrolyzed, and cellulose filaments precipitate. The surface of these precipitated cellulose filaments is crenulated, a characteristic of wet-spun fibers. Unlike PAN precursor fiber, cellulose fiber does not need to be oxidized in order to render it infusible. Nevertheless, because oxidation significantly improves its carbon yield, the cellulose precursor fiber is oxidized by heating it in air to temperatures as high as 40OoC. Initially, as the fiber is heated, the physically absorbed water is desorbed. As heating continues, additional water is evolved because of the reaction of hydroxyl groups in the cellulose. Finally, as the cellulose begins to decompose, CO2, CO, and water are given off, and aromatization of the structure begins (refs. 7 and 8). Because the cellulose polymer decomposes as it is stabilized, prestretching or stabilizing under tension (useful for PAN precursor fibers) is ineffective (ref. 9). After being stabilized, cellulose precursor fibers are carbonized and graphitized in an inert atmosphere at temperatures similar to those used for PAN. However, because the cellulose molecules in the precursor fiber lose most of their axial

26

Carbon Fiber Manufacturing

gq0m r

HO Ho

1

:HO

:HO CH,OH

i

OH

1

%OH

in OH

OH

Cellulose

w--z?&: -350°C

...- 0

....

0

0 0

... 0

Figure 7.Molecular structure of cellulose and approximate structure during thermal degradation to carbon (ref. 8). orientation during pyrolysis, the fibers are strained at this high temperature to increase the preferred orientation and improve the final mechanical properties. The filaments are quite plastic at high temperatures and can be stretched as much as 150 percent. When stretched 100 percent during graphitization at 28OO0C, fibers with a modulus approaching 720 GPa can be formed (ref. 9). However, if the same fibers are graphitized (but not stretched) at the same temperature, they attain a modulus of only 72 GPa. The overall yield for converting the cellulose precursor fiber to carbon fiber ranges from 10 to 30 percent, compared with 40 to 50 percent for the PAN precursor. This low yield is the direct result of the low carbon content of cellulose (44 percent) and the extensive decomposition that occurs during stabilization. This low conversion, especially when coupled with the expense of the stretchgraphitization, accounts for the high cost of rayon-based carbon fibers.

Manufacture of Pitch-Based Carbon Fibers Mesophase pitch-based carbon fibers are an attractive precursor candidate for carbon fibers because of the high availability of low-cost raw pitch. There are several reasons why the mesophase pitch process should produce a lower cost,

27

C C Materials and Composites high-performance fiber (estimated as $16/kg). First, the starting material (petroleum or coal-tar pitch) costs 40 to SO percent less than the monomers used to form PAN. Second, because pitch-based carbon fiber begins with a structure closer to graphite than PAN does, less energy is required to convert it to graphite. Because of this, lower carbonization temperatures and/or shorter carbonization times are required in the production of pitch-based carbon fibers. Third, mesophase precursor fiber contains a smaller percentage of nitrogen, hydrogen, and other noncarbon elements than PAN precursor fiber and, therefore, less material is driven off during carbonization. Because of this, the percent yield (in kilogiams of carbon fiber per kilogram of precursor fiber) is approximately 75 percent for mesophase pitch precursor fiber compared with only 40 to 45 percent for PAN precursor fiber.

Mesophase Pitch Mesophase pitch can be produced by the thermal or catalytic polymerization of a suitable petroleum or coal-tar pitch. When a highly aromatic pitch, such as a decant oil pitch, is heated to temperatures of 4OO0C to 4SOoC for approximately 40 hr, 45 to 65 percent of it will transform from an isotropic material to an optically anisotropic fluid phase-a mesophase or liquid crystal (ref. 10). A free-radical mechanism is believed to be responsible for polymerization of the carbonaceous material (ref. 11). Another method is to use solvents such as benzene, heptane, and toluene to first extract a portion of the isotropic pitch. The solvent insoluble portion can be converted to an anisotropic pitch by heating to between 23OoC and 4OO0C for less than 10 min (ref. 12). The anisotropic, or oriented, phase is composed of stacked, polynuclear aromatic hydrocarbon molecules. These molecules tend to be disc-shaped with an average molecular weight of approximately 1000 (although the molecular weight can vary considerably). The molecular structure of the mesophase produced from coal-tar pitch is characterized by higher aromaticity, whereas the structure of the petroleum-derived mesophase is more open because of its higher content of aliphatic side chains (ref. 13). Figure 8 shows the structure of a typical polynuclear aromatic hydrocarbon in mesophase. Initially, small spheres of mesophase form in the isotropic pitch when heated for an adequate time at a sufficiently high temperature. Upon further heating, the concentration of mesophase spheres increases and causes the spheres to collide and coalesce, creating a mosaiclike, nematic liquid-crystal structure (ref. 14). Mesophase products that have a high average molecular weight and no side groups or small molecular components to cause disordering often decompose before becoming fluid enough to flow. Because of this, the mesophase used to melt spin fibers is normally a mixture of high molecular weight molecules that still have a small number of side groups. Therefore, commercial mesophase precursors have certain characteristics of both mixtures and solutions: they soften over a range of temperatures and orient under an applied stress.

28

Carbon Fiber Manufacturing

0 yf&mCH3

a% CH2

CH2

\

\

/

/

C/H Molecular = 1.50weight = 1178

’ ’

Harorn Haliph = 1 -30 carorn

Caliph = 6.15

/

CH

/

CH \ /HZ

/

CH3

/

\

\

\

\

I

Figure 8. Typical polynuclear aromatic hydrocarbon in mesophase (ref.13). (Molecular weight given in glmol.) Numerous studies have concluded that, in general, the flow behavior of mesophase pitches is shear thinning at low shear rates but approaches Newtonian at high shear rates (refs. 13 and 14). In addition, it is the large response time for changes in flow rate which indicates that mesophase pitch is somewhat viscoelastic (ref. 14). However, the most unusual characteristic of mesophase pitch is the extreme temperature dependency of its viscosity (ref. 15). Even though mesophase pitch can be formed into fibers by conventional melt-spinning techniques, extremely precise temperature control is required (ref. 16).

Melt-Spinning Mesophase Precursor Fibers Figure 9 shows a schematic of a process for melt-spinning mesophase precursor fibers. Normally, the extruder screw consists of three zones: solid feed, melting, and pumping. The initial zone transports the solid mesophase feed to the melting zone, where it is heated to a temperature at which its viscosity is approximately 200 Pa.s (the optimum viscosity depends on the exact composition of the mesophase being extruded). Then, the pumping zone of the extruder forces this molten precursor into the top of a die head. The die head often contains a filter for removing solid impurities from the precursor. Finally, the molten mesophase exits 29

CC Materials and Composites through a multiple-hole spinnerette attached to the bottom of the die head. An initial orientation develops as the liquid crystalline precursor flows through the small capillaries in the spinnerette. As the precursor exits the capillaries, it is simultaneously cooled by the quench air and drawn before wind-up, yielding a precursor fiber with a high degree of molecular orientation. Because the aromatic, sheet-like molecules are already oriented in the direction of the fiber axis, additional drawing of the as-spun fiber is unnecessary. Melting zone

Filter

Quench air

Fiber wind-up device

Figure 9. Schematic of process for melt-spinning mesophase precursor ,fibers.

Because of the extreme temperature dependency of mesophase, fiber diameters can vary widely if the spinnerette temperature is not accurately controlled. Edie and Dunham (ref. 15) showed that only a f3.5OC variation in temperature across the face of the spinnerette can result in a f 1 5 percent variation in the diameter of the as-spun fibers (ref. 16). Their analysis also indicated that, even when process conditions during spinning are controlled, the tensile stress in the mesophase filament is almost one-half of its ultimate strength.

30

Carbon Fiber Munufucturing Heat Treatment of Mesophase Precursor Fibers Oxidation of Mesophase Precursor Fibers After spinning, the pitch-based fiber must be thermoset, in a manner similar to the PAN process, to render it infusible. The exact temperature and time required depend on the chemical composition and diameter of the mesophase fiber. The temperature must be below the softening point of the mesophase to minimize any fiber-to-fiber sticking. However, higher temperatures increase the rate of the stabilization reactions, decreasing the time required for this step. Commercially, the temperature selected for stabilization is a compromise between minimizing the required time for this process step and maximizing the mechanical properties of the final carbon fiber. Typically, the as-spun mesophase fibers are heated to temperatures of approximately 3OO0C for a period ranging from 30 min to 2 hr to be adequately stabilized for final heat treatment. Because the as-spun mesophase fiber already possesses a high degree of molecular orientation, applying tension during stabilization is unnecessary. The low tensile strength of the mesophase fiber, both before and after stabilization, makes fiber handling during this step extremely difficult. Even though the final carbon fiber exhibited a tensile strength of 2.1 GPa, Mochida et al. (ref. 17) tested mesophase fibers before carbonization and found their tensile strength to be only 0.04 GPa (less than 2 percent of its final strength after carbonization). This lack of fiber strength restricts the design of the oxidation ovens used for mesophase fibers. Obviously, fiber handling must be minimized to avoid fiber breakage. Figure 10 shows an apparatus for oxidizing the as-spun fiber without removing it from the spool used for spinning. Air, heated to the proper oxidation temperature, is forced through the porous wind-up spool and then passes through the fiber bundle. Because the oxidation reaction is exothermic, this flow geometry is important for heat as well as mass transfer. Designs such as this minimize damage to the as-spun fiber by completely eliminating fiber handling during oxidation. Commercially, many other processes are used to oxidize mesophase fibers. Although the designs vary considerably, all attempt to minimize handling of the fragile, uncarbonized mesophase fiber (ref. 16).

Carbonization and Graphitization After thermosetting, mesophase fibers (like PAN and cellulose precursor fibers) are either carbonized or graphitized in an inert atmosphere to develop their final properties. However, when mesophase fibers are carbonized the principal gases that evolve are CH4 and H2. Like PAN precursor fibers, most of these gases are evolved below 1000°C. Thus, normally the stabilized mesophase fibers also are precarbonized for a few minutes at 9OO0C to 1000°C. After precarbonization, they are either carbonized or graphitized at the desired temperature. Here again,

31

CC Matemals and Composates Exhaust gas

Oxidizing spool of fiber Inlet air

Air heating zone

Figure 10. On-the-spool oxidation of mesophase precursor fibers (ref. 16). hydrogen is the principal gas evolved above 1000°C. While still not extremely strong, after oxidation the mesophase fiber can be handled if the tow is sufficiently large. Thus, the ovens used to carbonize and graphitize mesophase fibers (fig. 11) are similar to those used to process PAN precursor fibers. The process used to surface-treat and size mesophase pitch-based carbon fibers is similar to that used for PAN-based fibers. However, because pitch-based fibers are less reactive to surface oxidation, more severe reaction conditions are used during surface treatment.

Power connection Water cooling coils

Oxidized fiber

Carbonized fib

Figure 11. Schematic of carbon resistance furnace used to carbonize mesophase precursor fibers.

32

Carbon Fiber Manufacturing Isotropic Pitch-Based Carbon Fibers Currently, a variety of carbon fibers are produced from isotropic pitch. However, unless these fibers undergo an expensive and difficult final stretchgraphitization step, their modulus is an order of magnitude less than that of mesophase pitch-based carbon fibers. Although they are useful for applications such as filtration, asbestos replacement, and static dissipation, the poor mechanical properties of isotropic pitch-based carbon fibers limit them to nonstructural applications. Therefore, the manufacturing of isotropic pitch-based carbon fibers has not been included in this chapter, and the reader is referred to Edie (ref. 16) for a detailed discussion of this process.

Manufacture of Vapor-Grown Carbon Fibers This process is able to produce only short, discontinuous lengths of carbon fiber. Nevertheless, these short fibers may be attractive for applications such as CC brake pads. Since some believe that the vapor-growth process may produce the first low-cost, discontinuous, high-performance, carbon reinforcing fiber, it is not surprising that several companies currently are conducting pilot-scale studies to evaluate its potential. Although only recently developed as a continuous process for producing highperformance reinforcing fibers, this technique was one of the first used to produce carbon filaments. Hughes and Chambers first detailed the vapor-growth process in an 1889 patent (ref. 18). In this patent they showed that small carbon filaments could be grown in a hot iron crucible under an atmosphere of methane and hydrogen-the vapor-growth process. Vapor growth is a dendritic type of growth from a catalyst particle. As Hughes and Chambers first discovered, metallic particles, normally containing iron as the primary constituent, tend to catalyze the growth of thin, partially graphitic fibers, when exposed to a hydrocarbon atmosphere at temperatures of approximately 1000°C. If the carburizing potential of the gas is low, a fraction of the fibers can be grown to macroscopic length, while still retaining the diameter of the catalyst particle. Because the most effective catalyst particles have a diameter of only 15 nm, these initial fibers are extremely thin. However, if the carbon potential of the gas is raised to a sufficient level, pyrocarbon can be deposited on the surface, permitting the filament diameter to increase to that of conventional carbon fibers (approximately 10 pm). This step is critical because the small diameter of the initial filaments makes them a potential carcinogen. Because pyrocarbon deposits with the basal planes parallel to the fiber surface, the fiber is highly oriented and has a high modulus. Figure 12 shows a schematic of the various stages of this catalyst-induced growth of carbon fibers (ref. 19). Tibbetts and his coworkers at General Motors Corporation used this growth technique to produce filaments with lengths up to 30 cm in an atmosphere of

33

-

CC Muteriu1.s and Composites

0 0 0

---

Catalyst particles

Filaments

( a ) Fiber growth stage.

Fibers

(b)Fiber thickening stage.

Figure 12. Schematic showing fiber growth and thickening stages during vapor-growth process. methane and hydrogen (ref. 19). Several metals (including nickel, cobalt, ironnickel powder, and Fe(N03)~)have been used as catalysts (refs. 19-21). Even though the feed stock (methane and hydrogen) was inexpensive and process temperatures of only 1000°C were used, the batch nature of this process used for early studies made it uneconomical for commercialization.

To overcome the low productivity of the batch process, Koyama and Endo (ref. 22) recently patented a continuous method for producing vapor-grown fibers. In their process, the catalyst particles are either incorporated in the feedstock or produced in the reactor by the decomposition of an organometallic. A simple schematic of this vapor-grown carbon fiber process is shown in figure 13. The catalyst and hydrocarbon feed are introduced at the top of the heated reactor, and

34

Carbon Fiber Manufacturzng short fibers are continuously withdrawn from the bottom. Fiber lengthening and thickening can be continuously controlled by adjusting the carbon potential of the gas within the reactor. The technique, which could be considered fluidized catalytic growth, allows carbon filaments with varying length, diameter, and physical properties to be continuously produced. If the thickening step can be accurately controlled and the projected process costs are correct, this continuous process could well replace isotropic and mesophase pitch-based, as well as PAN-based carbon fibers, in composite applications where chopped or short-fiber reinforcement is adequate.

Hydrocarbon feed (CH4, H2)

Heater

Figure 13. Schematic of continuous process for producing vapor-grown carbon jibers.

Mechanical Properties of Carbon Fibers As the next chapter will explain, the structure of carbon fibers, to a large extent, controls their tensile strength and modulus. Because of this, the manufacturers

35

CC Materzals and Composates of both PAN-based and pitch-based carbon fibers are attempting to develop new methods that can modify this structure during the fiber formation or heat treatment steps. Currently, as figure 14* shows, PAN-based carbon fibers exhibit higher tensile strengths, but lower moduli than mesophase pitch-based carbon fibers. However, the new varieties of mesophase pitch-based fiberst recently introduced by du Pont and Nippon Steel (denoted in fig. 14 as “Improved mesophase pitch”) exhibit significantly improved tensile strengths. The reported moduli for vaporgrown fibers are comparable to carbonized PAN fibers, but their tensile strengths are slightly lower. As expected, isotropic pitch fibers exhibit the lowest strengths and moduli of all carbon fibers. In the manufacture of both PAN-based and mesophase pitch-based carbon fibers, increasing the final heat treatment temperature improves the degree of preferred orientation within the fiber and, thus, the fiber modulus. Because of this, the various grades of fiber available from a particular manufacturer are normally the result of changes in this temperature.

I-

0

100

200

300

400

500

600

700

800

900 1000

Fiber modulus, GPa

Figure 14. Tensile strength and modulus of various types of carbon fibers (ref. 19 and footnotes *, t, and 1). G means the final heat treatment temperature is above 20OO0C,and C indicates it is below 2000°C.

The more perfect graphitic structure of mesophase pitch-based carbon fibers, compared with PAN-based carbon fibers, accounts for its higher thermal

* Bacon, R:

Amoco Performance Products, Incorporated, personal communication, 1989.

t Sato, K.: Nippon Steel Corporation, Tokyo, Japan, personal communication, 1989.

*

Ross, R.: E. I. du Pont de Nemours and Company, Incorporated, Chattanooga, Tennessee, personal communication, 1989.

36

Carbon Fiber Manufacturing conductivity. In fact, mesophase pitch-based fibers recently developed by Amoco Performance Products, Incorporated, exhibit a thermal conductivity that is three times that of copper.

Summary The precursor fibers used to produce current commercial carbon fibers are produced by melt, melt-assisted, and solution spinning. Melt spinning normally is the preferred fiber formation process because it eliminates the problems of solvent recovery and produces a purer precursor fiber. However, conventional melt spinning cannot be used for polymers, such as PAN and cellulose, which degrade below their melt temperatures. Nevertheless, melt-assisted spinning, a new process, permits PAN to be spun as pseudo-melt. The PAN-based and rayon-based precursor fibers are thermoset, carbonized, and graphitized with similar equipment and at similar conditions. However, their low tensile strength makes pitch-based carbon fibers much more difficult to handle before final heat treatment; therefore, special oxidation ovens are often used for this product. Short, vapor-grown carbon fibers represent the latest entry to the high-performance fiber field. If health issues can be adequately addressed, discontinuous filaments could become a low-cost reinforcement for composites.

Acknowledgments The authors thank G. P. Daumit of BASF Structural Materials, Incorporated, and Elsevier Science Publishers B. V. for permission to reproduce figure 3. Thanks are also given to Kluwer Academic Publishers for permission to reproduce figures 8, 10, and 12. Finally, the authors thank R. Bacon for generously providing the original drawing of figure 14.

References 1. Diefendorf, R. J.; and Tokarsky, E.: High-Performance Carbon Fibers. Polymer Eng. & Sci., vol. 15, no. 3, Mar. 1975, pp. 15G159. 2. Ram, Michael J.; and Riggs, John P.: Process for the Production of Acrylic Filaments. U.S. Patent 3,657,409, Apr. 1972. 3. Daumit, Gene P.; and KO, Yoon S.: A Unique Approach to Carbon Fiber Precursor Development. High Tech-The Way Into the Nineties, Klaus Brunsch, Hans-Dieter Golden, and Claus-Michael Herkert, eds., Elsevier Science Publ. Co., Inc., 1986, pp. 201-213.

4. Delmonte, John: Technology of Carbon and Graphite Fiber Composites. Van Nostrand Reinhold Co., c.1981.

5. Clarke, A. J.; and Bailey, J. E.: Oxidation of Acrylic Fibres for Carbon Fibre Formation. Nature, vol. 243, no, 5402, May 18, 1973, pp. 146-154.

37

C C Materials and Composites 6. Thome, D. J.: Manufacture of Carbon Fibre From PAN. Strong Fibres, W. Watt and B. V. Perov, eds., Elsevier Science Publ. Co., Inc., 1985, pp. 475-494. 7. Riggs, Dennis M.; Shuford, Richard J.; and Lewis, Robert W.: Graphite Fibers and Composites. Handbook of Composites, George Lubin, ed., Van Nostrand Reinhold Co., c.1982, pp. 196-271. 8. Fitzer, E.: Carbon Fibers: Present State and Future Expectations. Carbon Fibers, Filaments and Composites, J. Figueiredo, C. A. Bemardo, R. T. K. Baker, and K. J. Hiittenger, eds., Kluwer Academic Publ., 1990, pp. 3-41. 9. Bacon, Roger: Carbon Fibers From Rayon Precursors. Chemistry and Physics of Carbon-A Series of Advances, vol. 9, P. L. Walker, Jr., and Peter A. Thrower, eds., Marcel Dekker, Inc., 1973, pp. 1-102. 10. Singer, Leonard Sidney: High Modulus, High Strength Carbon Fibers Produced From Mesophase Pitch. U.S. Patent 4,005,183, Jan. 1977.

11. Singer, L. S.; and Lewis, I. C.: ESR Study of the Kinetics of Carbonization. Carbon, vol. 16, no. 6, 1978, pp. 417423. 12. Diefendorf, Russell J.; and Riggs, Dennis M.: Forming Optically Anisotropic Pitches. U.S. Patent 4,208,267, June 1980. 13. Fitzer, E.; Kompalik, D.; and Mayer, B.: Influence of Additives on Pyrolysis of Mesophase Pitch. Carbon ‘86Proceedings of the International Conference on Carbon, Deutschen Keramishchen Gesellschaft, Bad Honnef (Baden-Baden, Federal Republic of Germany), 1986, p. 842. 14. Nazem, F. F.: Flow of Molten Mesophase Pitch. Carbon, vol. 20, no. 4, 1982, pp. 345-354. 15. Edie, D. D.; and Dunham, M. G.: Melt Spinning Pitch-Based Carbon Fibers. Carbon, vol. 27, no. 5, 1989, pp. 647455. 16. Edie, D. D.: Pitch and Mesophase Fibers. Carbon Fibers, Filaments and Composites, J. Figueiredo, C . A. Bernardo, R. T. K. Baker, and K. J. Hiittenger, eds., Kluwer Academic Publ., 1990, pp. 647455. 17. Mochida, Isao; Toshima, Hiroshi; Korai, Yozo; and Naito, Tsutomu: Modification of Mesophase Pitch by Blending. Part 2-Modification of Mesophase Pitch Fibre Precursor With Thermoresisting Polyphenyleneoxide (PPO). J . Mater. Sci., vol. 23, no. 2, Feb. 1988, pp. 678-686.

38

Carbon Fiber Manufacturing 18. Hughes, T. V.; and Chambers, C. R.: Manufacture of Carbon Filaments. U S . Patent 405,480,1889. 19. Tibbetts, G. G.: Vapor-Grown Carbon Fibers. Carbon Fibers, Filaments and Composites, J. Figueiredo, C. A. Bernardo, R. T. K. Baker, and K. J. Hiittenger, eds., Kluwer Academic Publ., 1990, pp. 79-94. 20. Baker, R. T. K.: Electron Microscope Studies of Catalytic Growth of Carbon Fibers. Carbon Fibers, Filaments and Composites, J. Figueiredo, C. A. Bemardo, R. T. K. Baker, and K. J. Hiittenger, eds., Kluwer Academic Publ., 1990, pp. 405439. 21. Endo, M.; and Komaki, K.: Formation of Vapor-Grown Carbon Fibers by Seeding Method of Metal Ultra-Fine Particles. Extended Abstracts of 16th Biennial Conference on Carbon, American Carbon Society, San Diego, CA, p. 523, 1983. 22. Koyama, T.; and Endo, M. T.: Method for Manufacturing Carbon Fibers by a Vapor Phase Process. Japanese Patent 1982-58,966, 1983.

39

Chapter 3 Effect of Microstructure and Shape on Carbon Fiber Properties D. D. Edie and E. G. Stoner Clemson University Clemson, South Carolina

Introduction 42 Carbon Fiber Processes 43 Effect of Graphite Structure on Fiber Properties 44 Brittle Failure Mechanism 45 Microstructure of Carbon Fibers 47 Microstructure of PAN-Based Carbon Fibers 47 Microstructure of Pitch-Based Carbon Fibers 48 Effect of Microstructure on Fiber Properties 53 Effect of Microstructure on Tensile Properties of Carbon Fibers

54

PAN-Based Carbon Fibers 54 Pitch-Based Carbon Fibers 55 Effect of Microstructure on Compressive Properties of Carbon Fibers PAN-Based Carbon Fibers 59 Pitch-Based Carbon Fibers 59 Effect of Fiber Shape on Fiber and Composite Properties Effect of Shape on Tensile Strength of Carbon Fibers

57

61 61

PAN-Based Carbon Fibers 61 Pitch-Based Carbon Fibers 61 Effect of Shape on Compressive Strength for Carbon Fibers

62

PAN-Based Carbon Fibers 64 Pitch-Based Carbon Fibers 66 Summary 66 Acknowledgments 66 References 67

41

C C Materials and Composites

Introduction In carbon-carbon (CC) composites, carbon fibers reinforce the brittle carbon matrix material. Because these reinforcing fibers determine, to a large extent, the strength and stiffness of this composite material, optimizing the properties of CC composites requires a thorough understanding of both the properties and peculiarities of carbon fibers. As in other classes of fiber-reinforced composite materials, the required fiber properties depend on the particular composite application. In certain propulsion applications of CC composites, high fiber strength is critical and stiffness is less important. In these applications, polyacrylonitrile (PAN) based carbon reinforcing fibers are the logical choice. In other structural applications, the CC composite experiences both tensile and compressive loading. This loading makes fiber compression properties critical, and, again, PAN-based fibers are preferred. In applications in which interlaminar strengthening is necessary, fiber strength and stiffness are not as critical as the ability to weave a complex fabric preform from the carbon reinforcing fibers. Here the low modulus of rayon-based carbon fiber, combined with its medium strength, results in the moderate strain to failure necessary for processing on weaving equipment. However, in large space structures, low thermal expansion or extremely high stiffness may be critical. Pitchbased carbon fibers are uniquely suited for these applications. Finally, the strength of the bond between the fiber and the matrix also can be critical to composite performance. A strong interfacial bond may make the composite more resistant to interlaminar shear. On the other hand, because the carbon matrix material is brittle, a weak interfacial bond can, as in ceramic composites, serve as a toughening mechanism. Thus, in CC composites, fiber-matrix bonding often is a compromise between toughening the composite and improving its interlaminar properties. To meet these different requirements, three types of carbon fibers presently are used in CC composites: these are rayon-based, PAN-based, and pitch-based fibers. Each has its own particular strengths and weaknesses. To better understand both the limitations and the potential of these different carbon fibers, this chapter briefly reviews the processes used to produce them and the ultimate properties that they can attain. The mechanism of brittle fracture then is discussed to better understand the practical limits for carbon fiber properties. Finally, the effect that processing has on the carbon fiber structure and the relationship between this structure and the physical properties of the fiber will be explained. This explanation will show the potential for future increases in physical properties for each type of carbon fiber. Rayon-based fibers were the first carbon fibers used in CC composites. However, because of their higher strength and stiffness, PAN-based and pitch-based carbon fibers are used to reinforce the vast majority of current CC structures. This chapter will concentrate on PAN-based and pitch-based carbon fibers, the

42

Effect of Macrostructure and Shape on Carbon Fiber Properties two principal reinforcements used in current CC composites, and the two most promising candidates for the CC applications of tomorrow.

Carbon Fiber Processes Carbon fibers are produced by the thermal decomposition of various organic fiber precursors, such as PAN, rayon, and mesophase pitch. As mentioned in chapter 2, rayon and PAN are normally extruded into filaments using solutionspinning techniques. These precursor fibers then are thermoset by oxidizing them in air for extended periods. For PAN precursor fiber, this results in the formation of an infusible cyclized network of hexagonal, carbon-nitrogen rings. For rayon precursor fiber, some depolymerization occurs, and a similar aromatization of the remaining structure begins. The stabilized rayon and PAN fibers finally are carbonized by heating them in an inert atmosphere to approximately lSOO°C (ref. 1). Subsequent graphitization at temperatures ranging from 2SOO0C to 30OO0C can improve the mechanical properties of both fibers. The mechanical properties of rayon are enhanced considerably if the rayon is stretched during this graphitization. If the thermoset rayon precursor fiber is graphitized without stretching at 28OO0C, it develops a tensile modulus of only 69 GPa. The same fiber can develop a modulus of 690 GPa if it is stretched 100 percent during graphitization at the same temperature (ref. 2). However, this stress-graphitization step is extremely expensive, and the resulting high-modulus fiber is brittle, making it difficult to handle. The starting material used to produce pitch-based carbon fibers is usually a highly aromatic coal or petroleum pitch. For the production of high-performance fibers, the pitch first must be converted to an oriented liquid-crystalline material-a mesophase. Heating the raw pitch to temperatures ranging from 4OO0C to 41OoC for periods of up to 40 hr (ref. 3) usually accomplishes this conversion. Solvent extraction (ref. 4) is another method for preparing mesophase, which involves first separating the high-molecular weight fraction of the pitch. The extraction step removes the smaller disordering molecules and concentrates the higher molecular weight portion of the pitch in the insoluble portion. This insoluble fraction then can be polymerized to a 100-percent anisotropic phase by heating it to between 23OoC and 4OO0C for only 10 min. This liquid-crystalline product, or mesophase, is then melt-spun into a precursor fiber. The viscosity of commercial mesophase precursors is extremely temperaturedependent, making precise temperature control during melt-spinning vital. Because mesophase is actually a lyotropic liquid-crystalline solution, it becomes highly oriented during extrusion and subsequent drawdown (ref. 5). Therefore, unlike PAN and rayon, as-spun pitch precursor fiber has an extremely high degree of molecular orientation.

43

CC Materials and Composites Similar to PAN and rayon, the melt-spun mesophase pitch fibers must be thermoset to render them infusible and then carbonized to develop their final properties. Because mesophase pitch fibers develop an extremely high degree of molecular orientation during fiber formation, the object of subsequent heat treatment is not to develop, but instead to preserve, the molecular orientation. Because of this, the conditions used in the thermosetting and carbonization steps are usually mild compared with those used for PAN and rayon precursor fibers. Typically, temperatures of approximately 3OO0C and exposure to air for a period ranging from 30 min to 2 hr are required to fully thermoset mesophase fibers. Subsequent heat treatment in an inert atmosphere, such as nitrogen, allows the thermoset mesophase precursor fiber to develop its final properties. Because the precursor fiber is already highly oriented, mesophase pitch precursor fibers develop tensile moduli of approximately 690 GPa when heat-treated at 28OOOC for a few minutes, even when no tension is applied (ref. 1).

Effect of Graphite Structure on Fiber Properties Carbon fibers are composed of 99.9 percent pure carbon, most of which is arranged into graphite crystallites. These graphite crystallites organize into layer planes, and the mechanical properties of carbon fibers are a result of this structure. This layer-plane structure and the corresponding lattice dimensions of graphite, as well as its approximate orientation within carbon fibers, are shown in figure 1 (ref. 6). Within these layers, chemical bonds link the carbon atoms in the crystallographic a-direction. These anisotropic, nonpolar, cT-bonds, created by the sp2 hybridization of the electron orbitals in the carbon layers, make the graphite structure extremely strong in the crystallographic a-direction (ref. 7). This makes the theoretical tensile modulus and the ultimate tensile strength of graphite extremely high, approximately 1060 GPa (ref. 8) and 106 GPa (refs. 9 and lo), respectively, for a load applied along this crystallographic axis. On the other hand, very weak van der Waals bonding exists between atoms in adjoining planes. These weak bonds make the mechanical properties of the graphite crystal in this direction, the crystallographic c-direction, quite low. As a result, the theoretical tensile modulus for a load applied in this direction, normal to the crystallographic a-axis, is only 35 GPa (ref. 1). The structure of the graphite crystallites in a carbon fiber is not perfect, as shown in figure 1. The layer planes of carbon are slightly offset, creating what is termed turbostratic graphite. This offset also results in a slight increase in the interlayer spacing (3.37 A to 3.45 A), compared with the 3.35 A found in a perfect crystal. Even though the interlayer spacing is not perfect, the orientation of these graphite layers is more or less parallel to the fiber axis, and it is the strong covalent bonding within these layer planes that accounts for the longitudinal properties in carbon fibers. Thus, as one would expect, the layer planes in high-modulus carbon fibers are almost perfectly aligned with the fiber axis.

44

Effect of Microstructure and Shape on Carbon Fiber Properties

,a

?

?

-

Fiber axis

m -w

7

a-direction

3.35 A for true graphite (3.37 A t0 3.40 A for mesophase fibers)

-

2.46 A

4

n oi t-";;--

'

b

.42 A c-direction

Figure I . Structure of graphite and approximate orientation in carbon fibers (ref. 6). Today commercial pitch-based carbon fiber is available with a tensile modulus that is 85 percent of the theoretical maximum for the perfect graphite crystal. By comparison, the tensile strength of this fiber is only 5 percent of that theoretically possible. Although increasing the layer plane alignment directly improves the tensile modulus of carbon fibers, flaws limit the tensile strength of these brittle fibers. A brief review of the interaction between crystallite size and brittle fracture follows to help understand how this limitation relates to the microstructure of the carbon fiber.

Brittle Failure Mechanism Because flaws limit the tensile strength, considerable effort has been made to minimize contamination in both PAN-based and pitch-based carbon fibers (refs. 11 and 12). Yet even when flaw-inducing particles are minimized, structural flaws within carbon fibers still limit the tensile strength. Johnson (ref. 13) recently presented an excellent explanation of these structural limitations of carbon fibers and of the interaction between microstructure and the flaw sensitivity of these fibers, based on the mechanism for tensile failure proposed by Reynolds and Sharp (see ref. 13). As mentioned previously, the crystallites and carbon layer planes are not aligned perfectly in carbon fibers, and misoriented crystallites are relatively common. When a stress is applied parallel to the fiber axis, these crystallites align until movement is restricted by a disclination in the structure (as shown in fig. 2). If the stress is sufficient, a basal-plane rupture of the misoriented crystallite can relieve the stress within the fiber. When the size of this ruptured crystallite (perpendicular to the fiber axis) is larger than the critical flaw size, a catastrophic failure occurs, breaking the entire fiber.

45

CC Materials and Composites

‘

1

Figure 2. Reynolds and Sharp mechanism for tensile failure (ref 13). Misoriented crystallite linking two crystallites parallel to fiber axis (left), tensile stress causing basal plane rupture in L, direction (center), and catastrophic failure occurring if crystallite size is greater than j a w size (right). Even if the ruptured crystallite is smaller than the critical flaw size, catastrophic failure can still occur if the crystallites surrounding the disclination are continuous enough to allow a crack to propagate into neighboring crystallites. Johnson (ref. 13) indicates that this failure mode, created by extended three-dimensional (3-D) crystalline order, explains the difference in flaw sensitivity between PANbased carbon fibers and the more highly graphitic, pitch-based carbon fibers. The strong covalent bonding within the layer planes thus causes the tensile modulus of carbon fiber to increase as the crystallite orientation parallel to the fiber axis increases. However, when a stress is applied parallel to the fiber axis, misoriented crystallites must rupture to relieve the stress. As the size of the crystallites increases, it becomes more likely that some will exceed the critical flaw size. Because of this, the final balance in the physical properties of carbon fibers depends on the size, the extent of 3-D order, and the orientation of the graphite crystallites.

46

Effect of Microstructure and Shape o n Carbon Fiber Properties

Microstructure of Carbon Fibers Because the precursor fiber for PAN-based carbon fibers is formed by precipitating a linear polymer from solution, whereas the precursor fiber for pitch-based carbon fibers is formed by melt-spinning a cyclic organic liquid crystal, it is not surprising that the arrangement of the graphite crystallites (or microstructure) in these carbon fibers differs considerably. As explained in the previous section, the physical properties of carbon fibers are a direct result of their microstructure. Therefore, the microstructures of both PAN-based and pitch-based carbon fibers must be understood before the potential, as well as the limitations, of both types of carbon fiber can be appreciated.

Microstructure of PAN-Based Carbon Fibers Early work by Diefendorf and Tokarsky (ref. 14) indicated that the macrostructure of PAN-based carbon fibers, similar to that of many other synthetic precursor fibers, is fibrillar. Although the exact nature of this microstructure has been refined in more recent studies, the ribbon-like undulations (characteristic of a fibrillar structure) are an excellent visualization of the microstructure of PAN-based carbon fibers. Diefendorf and Tokarsky (ref. 14) also showed that the amplitude of the undulation in the fibrillar structure was highest in the center and lowest near the surface of PAN-based carbon fibers. This indicates that the modulus of a PANbased carbon fiber varies throughout the cross section of the fiber. Figure 3 shows a model of this microstructure. Johnson (ref. 13) studied various types of PAN-based carbon fibers in an effort to elucidate the relationship between the microstructure of these fibers and their properties. Using wide-angle X-ray diffraction, he found that the layer planes of PAN-based carbon fibers have no regular 3-D order. In the skin region of the fiber, small-angle X-ray diffraction and transmission electron microscope (TEM) analysis of longitudinal and transverse fiber sections revealed needle-shaped voids between crystallites and layer planes that essentially were parallel to the surface. However, in the core region, Johnson found that the layer planes were folded extensively, often through angles of 180'. He proposed that misoriented crystallites interlink with other oriented and misoriented crystallites, accounting for the voids, shown in figure 4. Based on these results, Johnson developed the 3-D schematic representation of the microstructure of PAN-based carbon fiber shown in figure 5. Endo (ref. 15) also conducted a crystallographic analysis of a PAN-based carbon fiber, Torayca M46, manufactured by the Toray Company. X-ray diffraction studies revealed no separation of the 100 and 101 peaks and the absence of a 112 peak. He also concluded that these fibers have little 3-D order. Endo found that the average crystallite thickness L, was approximately 6.2 nm and the average layer spacing (doo2) was 0.3434 nm.

47

CC Materaals and Composates

A

Figure 3. Schematic 3-0 structural model of Fortajl.5-Y PAN-based, 345 GPa tensile modulus carbon fiber (ref. 14). Based on these studies, PAN-based carbon fibers appear to contain extensively folded and interlinked turbostratic layers of carbon with an interlayer spacing considerably larger than that of true graphite. They show a low degree of graphitization, and the turbostratic layers are not highly oriented with the fiber axis.

Microstructure of Pitch-Based Carbon Fibers Unlike PAN-based carbon fibers, pitch-based fibers show a variety of microstructures. Varying the spinning conditions of the liquid-crystalline precursor produces these various microstructures (shown schematically in fig. 6), which are apparent upon microscopically examining the fiber cross section. Commercial fibers usually exhibit either radial, flat-layer, or random microstructures, and these three formations appear to be the preferred microstructures of mesophase pitch. Endo (ref. 15) showed that Carbonic pitch-based fibers have a radial-folded microstructure; other investigators have produced pitch-based fibers with line-origin, onion skin, and quasi-onion microstructures in laboratory experiments. Each type of microstructure, except for random microstructure, is viewed as a collection of large flat plates extending down the fiber axis, arranged in some geometric order

48

Effect of Microstructure and Shape on Carbon Fiber Properties

Figure 4 . Interlinked structure and resulting void (ref. 13).

Figure 5 . Schematic of microstructure of PAN-based carbon fiber (ref. 13).

49

CC Materials and Composates across the fiber cross section. In the random microstructure, the plates are relatively small and have no long-range geometric ordering across the fiber. Nevertheless, they are still oriented almost parallel to the fiber axis.

Flat layer

Radial folded

Line origin

@ v @L&

\\@!>y11 Q Q ,

Radial

Onion skin

Random

Quasi-onion

Figure 6. Microstructures of pitch-bused carbon fibers. In his study, Endo (ref. 15) compared several grades of commercial pitch-based fibers, Thornel PlOO and P120 (Amoco Performance Products, Incorporated) and Carbonic HM50, HM60, and HM80 (Kashima Oil Company), with the Torayca PAN-based fibers. Table 1 lists the tensile strength, Young's modulus, failure strain, interlayer spacing, and crystallite thickness of each fiber.

The X-ray diffraction results showed that only the Thornel fibers exhibited separation of the 100 and 101 peaks and the appearance of a 112 peak. This indicated that only the Thornel fibers had a high degree of 3-D ordering. Endo also found that, as the strength of Carbonic fibers increased, the X-ray diffraction profile began to resemble that of the Torayca PAN-based fibers. In addition, the layer planes in the Carbonic fibers appeared to be oriented within 15' of the fiber axis, while the layer planes in the Thornel fibers were oriented almost perfectly with the fiber axis. This misorientation in the Carbonic fibers was more typical of PAN-based fibers. Further TEM studies indicated that the layer planes of the Carbonic fibers were a folded structure, while the layer planes of the Thornel fibers exhibited no folds. These findings led Endo (ref. 15) to conclude that (1) Thornel fibers consist of straight, well-oriented plates with a

50

Fiber name

Tensile strength, GPa

Tensile modulus, GPa

Failure strain, percent

Interlayer spacing, do02 ? nm

Crystallite thickness,

2.2 2.4

690 830

0.3 0.3

0.3392 0.3378

24 28

2.8 3.0 3.5

490 590 790

0.6 0.5

0.3423 0.3416

13 15

0.4

0.3399

18

2.5

450

0.5

0.3434

6.2

Lc2

nm

Pitch-based fiber Thornel"

PlOO P120

Carbonicb HM50 HM60 HM80 Pan-based fiber Torayca'

M46

Run number 1

Stimng rate, rlmin 0.0

Interlayer spacing, do02 nm 0.3377

Crystallite thickness, nm 28

Microstructure Radial

3

Lc

1

2

9.3

0.3391

20

Quasi-onion

5

29.5

0.3392

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6

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0.3392

19

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*

Effect of Microstructure and Shape o n Carbon Fiber Properties

“ : I

Stirrer

---

‘ Capillary Pitch fiber

il D ! l a

Figure 8. Stirrers used to vary microstructure of pitch-based carbon fibers (refs. 16 and 17).

Hamada also measured the transverse magnetoresistance A p / p of these fibers. As the magnetic field increases, a highly graphitized fiber shows a positive and increasing A p / p , while a poorly graphitized or turbostratic fiber has a negative and decreasing A p / p . From these studies, he found that the radial microstructure showed a high degree of graphitization and that microstructures produced with stirring were turbostratic. The layer planes in Hamada’s fibers with radial microstructures (refs. 16 and 17), similar to those in high-modulus Thornel fibers, appear to be more perfectly oriented to the fiber axis and less folded than the layer planes in the Carbonic fibers Endo studied (ref. 15). However, the most significant result of Hamada’s work was the demonstration that interrupting the preferred flow pattern during mesophase extrusion modifies the fiber microstructure, resulting in a smaller average graphite crystal size, a larger layer plane spacing, and a lower degree of 3-D order or graphitization.

Effect of Microstructure on Fiber Properties Currently, neither PAN nor pitch precursor fibers develop balanced physical properties upon heat treatment. While PAN-based fibers dominate the market in high tensile strength applications, their tensile modulus is much lower than that of pitch-based carbon fibers. Conversely, current pitch-based fibers are the

53

CC Materials and Composites predominanr fiber used in high stiffness applications, but their strength normally is less than that of PAN-based carbon fibers. An analysis of the microstructure in terms of brittle fracture explains why the tensile properties of PAN-based and pitch-based carbon fibers differ and the reason for their imbalance in properties. Because carbon fibers fail by brittle fracture, this analysis also will provide insight as to how much, and how readily, the physical properties of these fibers can be modified.

Effect of Microstructure on Tensile Properties of Carbon Fibers Assuming that the brittle fracture mechanism proposed by Reynolds and Sharp (see Johnson, ref. 13) applies, the results of microstructural studies might explain the differences in strength and modulus of pitch-based and PAN-based carbon fibers. Given the strength and modulus of any carbon fiber, the Griffith relationship (ref. 18) can determine the critical crack size.

o2 - 'Ya_ 2E TC

_

~

(1)

where o is the ultimate tensile strength, E is the modulus, ya is the apparent surface energy, and C is the critical crack length. The most generally accepted value of Y~ is 4.2 J/m2; although higher values of this constant have been reported, they serve only to increase the critical crack size and, therefore, are not of interest in the limiting case.

PAN-Based Carbon Fibers As was previously mentioned, Endo (ref. 15) found the ultimate tensile strength of Torayca M46 to be 2.4 GPa, the Young's modulus to be 450 GPa, and the average crystallite thickness to be 6.2 nm. According to equation (l), the critical crack length for this fiber would be 209 nm. Johnson (ref. 13) notes that the large difference between critical crack length and crystallite thickness is typical of PAN-based carbon fibers, except in regions of enhanced crystallization surrounding inclusions. Because commercial producers use a high-purity PAN solution for spinning the precursor fibers to minimize such inclusions, the chance of failure by exceeding this critical size in a single crystal is also minimized. Thus, if failure is to occur by the Reynolds and Sharp failure mechanism, the crack must propagate outside the initiating crystallite into neighboring crystallites. The large interlayer spacing and lack of 3-D order in PAN-based carbon fibers lower the probability that such propagation will occur, and the folding nature of the crystallites will probably also hinder crack growth. Thus, the lack of inherent orientation of the precursor, which hinders the development of graphitic structures, actually serves as a crack-stopping mechanism, increasing the final fiber strength.

54

Effect of Microstructure and Shape on Carbon Fiber Properties The Reynolds and Sharp failure mechanism may also explain the higher failure strains of PAN-based fibers. The combined effect of many small crystallites, which can fail and relieve the applied stress without causing a catastrophic failure, could yield this increased elongation. Even the relatively low-tensile modulus of PAN-based carbon fibers is a direct result of its microstructure. Recall that the high modulus of carbon fibers is a direct result of strong bonds in the layer planes oriented parallel to the fiber axis. In PAN-based fibers, the layer planes are less oriented with the fiber axis than they are in pitch-based fibers. In addition, the low degree of graphitizability of PAN-based fiber implies that these planes are small. A lower modulus for PAN-based carbon fibers, therefore, would be predicted from crystallographic analysis and the microstructural model proposed by Johnson (ref. 13). Because this folded, turbostratic microstructure is created during initial fiber formation and is characteristic of many solution-spun polymers, major improvements in the modulus or graphitizability of PAN-based carbon fibers are unlikely unless a totally different spinning technique is used.

Pitch-Based Carbon Fibers Using the data from Endo’s study of Thornel and Carbonic pitch-based carbon fibers (ref. 15), the prediction from equation (1) shows that their critical crack lengths will be approximately 130 nm. Although this is smaller than the critical crack length estimated for the Torayca PAN-based fiber studied by Endo, it is still a factor of 10 larger than the average crystallite thickness. However, Endo found that high-modulus Thornel fibers had a high degree of 3-D order, and thus, a continuous medium for crack propagation. He also postulated that the flat-layered structure of Thornel fibers will propagate cracks more easily than the folded crystallites found in PAN-based fibers. If the crack must propagate outside the initiating crystallite for a fiber to fail, these observations would explain why the high modulus Thornel fibers showed inferior tensile strengths. On the other hand, Endo found the modulus of Thornel P120 to be nearly 83 percent of that theoretically possible. This extremely high modulus is a direct result of the nearly perfect orientation of the closely spaced layer planes found in the Thornel pitch-based fibers. Endo’s study (ref. 15) revealed that, although both the Carbonic and the Thornel pitch-based carbon fibers are melt-spun from a mesophase pitch precursor, they are quite different. The low degree of graphitization and the presence of crystallite folding appear to be responsible for the increased strength of Carbonic fibers. Endo also found that, as the X-ray diffraction results of these fibers become more similar to those of PAN-based fibers, the strength increases. Carbonic fibers also have significantly higher failure strains than Thornel fibers, a phenomenon that also might be explained as the cumulative effect of numerous noncatastrophic failures in misoriented crystallites.

55

CC Materials and Composites Based on the similarities between Carbonic and PAN-based fibers, one would expect the Carbonic pitch-based fiber to have a lower modulus than the Thornel pitch-based fiber. However, while Carbonic HM80 is 46 percent stronger than Thornel P120, its modulus is only 5 percent lower. Evidently, a high degree of graphitization is not necessary to develop a high modulus. Instead, high preferred orientation, which is characteristic of a melt-spun fiber using a liquid-crystalline precursor, is largely responsible for the resulting fiber modulus. Hamada et al. (refs. 16 and 17) also found that a lower degree of 3-D order leads to increased strength in pitch-based carbon fibers. He found that fibers with a nonradial, turbostratic microstructure,produced by disturbing the flow profile of the mesophase prior to extrusion, were significantly stronger than those with a radially oriented microstructure and the same modulus. Interestingly, the pitch-based fibers recently introduced by Nippon Steel Corporation and E. I. du Pont de Nemours and Company* have a random microstructure. This microstructure probably results in a lower degree of 3-D order within the fiber. Although their moduli are similar to those of the Amoco Performance Products, Incorporated, Thornel pitch-based carbon fibers, these new fibers show considerably higher tensile strengths (see fig. 9, refs. 19 and 20). Although these fibers have a random microstructure when viewed perpendicular to the fiber axis, the graphite layer planes still have excellent alignment along the fiber axis. This orientation perpendicular to the fiber axis usually is determined during the flow through the extrusion capillary. Nazem (ref. 21) demonstrated one method of creating a random orientation during fiber formation with mesophase. By extruding mesophase through a round extrusion capillary that had a small lengthto-diameter ratio and contained a porous media, he prevented a stable, parabolic flow profile from developing during extrusion. The resulting fiber had a totally random microstructure perpendicular to the fiber cross section. The melt-spun fiber still developed the excellent modulus that is characteristic of pitch fibers because the drawdown of the molten mesophase after having been extruded through the capillary still oriented the mesophase molecules parallel to the fiber axis. Edie et al. (ref. 22) also showed that changing the extrusion capillary can readily change the microstructure of the melt-spun mesophase fibers. In this study noncircular extrusion capillaries were used to melt-spin a mesophase precursor into noncircular, pitch-based carbon fibers. These fibers had a line-origin microstructure that emanated from the center lines of the lobes of the noncircular fibers. Even when the noncircular fibers were cooled slowly (permitting them to collapse into a circular shape before solidification), they retained a lobal microstructure, confirming that the microstructure was developed during flow through the extrusion capillary.

*Sato, K., Nippon Steel Corporation, Tokyo, Japan, personal communication, 1989.

56

6-

a

a (3 4 5 0

c

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N

c u)

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z=

2-

i

p

p

3;

o

, ,o <

T T Amoco

Nippon Steel

I

I

I

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I

CC Materials and Composites strength. Even though the mechanism for fiber failure in compression is not fully understood, it is evident that, for a given tensile modulus, round PAN-based carbon fibers have a higher compressive strength than round pitch-based carbon fibers. It is agreed, therefore, that microstructure has a major influence on the compression properties of carbon fibers.

(d

a

PAN-based carbon fiber v)

v)

0

b

Mesophase carbon fiber

0 LL 0-

1

I

200

300

I

I

I

400 500 600 Fiber tensile modulus, GPa

I

700

800

Figure 10. Compressive strength as function of tensile modulus for carbon fibers (ref. 23).

Kumar and Helminiak (ref. 24) have shown that, for both pitch-based and PAN-based carbon fibers, the compression strength decreases as the thickness of the graphite crystallites increases. Additionally, increases in the void content within the fiber or in the interlayer spacing correlate with increased compression strength. Therefore, it appears that a carbon fiber with an extended highly graphitic structure is likely to be weak in compression. On the other hand, fibers with microstructures that impede the development of an extended graphitic structure should exhibit higher compressive strengths.

58

Effect of Microstructure and Shape on Carbon Fiber Properties PAN-Based Carbon Fibers Johnson has used a single-filament recoil test to study the compression failure of PAN-based and pitch-based carbon fibers (ref. 25). In these tests, he found that the first response of high-strength PAN-based and low-modulus pitch-based fibers to a compression load was buckling and the formation of kink bands. As deformation progressed, a tensile crack formed on the tension side of the buckled fiber, and the kink bands propagated inward. Finally, the tensile crack and the kink bands met, and failure occurred. This failure sequence is shown in figure 1l(a).

( a ) High-strength PAN-based and low-modulus pitch-based fiber.

(h)High-modulus pitch-based fiber.

Figure 11. Compressive failure of low-modulus and high-modulus carbon fibers (ref 2.5).

Pitch-Based Carbon Fibers When Johnson studied high-modulus pitch-based carbon fibers using the same technique, he observed a markedly different failure mode (ref. 25). Under initial compression, these fibers developed kink bands across the whole fiber by simple shear deformation. Then, as shown in figure 1l(b), the fracture simply propagated along the kink plane. These observations have led Johnson to conclude that the low compressive strengths of high-modulus pitch-based carbon fibers are caused by

59

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&

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ppon Steel

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Effect of Microstructure and Shape on Carbon Fiber Properties

Effect of Fiber Shape on Fiber and Composite Properties Changing the shape of the fiber as well as the microstructure can change the mechanical properties of a fiber. Often the effect of these two variables is difficult to separate. However, fibers with noncircular cross sections have long been used in the synthetic fiber industry to improve the wetting characteristics and increase the buckling resistance of polymeric fibers. During 1989, Owens-Coming Fiberglas Corporation introduced a variety of glass fiber with a trilobal cross section, claiming the shape resulted in a significant improvement in fiber stiffness as well as tensile strength (ref. ?6). This new noncircular fiber is, coincidentally, targeted for composite appl cations.

Effect of Shape on Tensile Strength of Carbon Fibers PAN-Based Carbon Fibers Carbon fibers, similar to other fibers, also appear to display different properties when circular fibers are compared with noncircular fibers. Because they are formed by precipitation, the cross-sectional shape of solution-spun PAN fibers tends to be either circular or dogbone. Figure 13 shows a comparison of the balance between modulus and strength for these two shapes of PAN-based carbon fibers, as determined by Diefendorf and Tokarsky (ref. 14). This same trend also has been found in other studies (refs. 27 and 28). As the modulus of a dogbone fiber increases, so does its strength. For increases in modulus above 250 GPa, however, the strength of a round fiber actually drops (ref. 29). While this decrease may be related to flaw-inducing impurities in the fiber (ref. 12), differences in the flaw sensitivity of the fiber, or differences in residual stresses (ref. 14), the evidence remains that the balance of properties differs between round and dogbone PANbased carbon fibers.

Pitch-Based Carbon Fibers Because pitch-based carbon fibers are melt-spun, they can be extruded into a variety of cross-sectional shapes merely by changing the shape of the extrusion die. Edie et al. (ref. 22) extruded a solvent-extracted mesophase through the trilobal spinnerette capillary shown in figure 14(a) to produce fibers with a trilobal cross section. When heat-treated at 19OO0C, the tensile strength of these noncircular fibers was 39 percent higher than round fibers of equal cross section that were produced for comparison. Because the microstructure of the trilobal fibers (fig. 14(b)) differed from the radial microstructure seen in the round fibers, it is difficult to tell if the increased strength is the result of the change in fiber shape or the change in microstructure. The investigators noted that this increase in strength could be caused by several factors, including the possibility that the

61

600 -

500

-

400

-

a

a

a

g 3 -0

E

2 .-

Dogbone

300 -

v)

c

I-" 200

100

I

I

I

I

1

Effect of Microstructure and Shape o n Carbon Fiber Properties

127 pm

(a) Noncircular extrusion capillary.

(b)SEM of trilobal fiber (effective diameter is 12.6 pm).

Figure 14. Noncircular capillary and resulting trilobal fiber (ref. 22).

(a)Noncircular extrusion capillary.

(b) SEM of trilobal fiber (effective diameter is 25 pm).

Figure 15. Noncircular capillary and resulting trilobal fiber (ref. 30). it seems likely that changing the microstructure will increase the compressive properties of pitch-based carbon fibers. However, recent studies indicate that changes in the macroscopic-buckling characteristics of carbon fibers may result in additional increases in their compressive properties.

63

CC Materials and Composites PAN-Based Carbon Fibers Recently, BASF Structural Materials, Incorporated, began producing PANbased carbon fibers using a new melt-spun precursor technology (ref. 31). In this process, the acrylonitrile is polymerized in an aqueous suspension, eliminating the need for an organic solvent. After the PAN is purified and dewatered, it is compounded, pelletized, and fed to an extruder. When melted in the extruder, the pellets, plasticized with excess water, form a homogeneous melt that can be extruded into fiber form. During extrusion the excess water flashes off, allowing the PAN precursor fiber to solidify. This process has several advantages compared with standard solution-spinning technology, including eliminating the need for solvent recovery and decoupling the polymerization step from the spinning step. Additional advantages are that, because large amounts of residual solvent do not have to diffuse out of the fiber during solidification, a more radially uniform structure should result. Also, since the fibers are essentially melt-spun, the fiber shape can be controlled. BASF Structural Materials, Incorporated, has already used this melt-spun precursor technology to produce experimental quantities of noncircular, PAN-based carbon fibers with ribbon and tetralobal cross sections (see fig. 16). These varieties are being developed for composite applications in which improved compressive and interlaminar shear strength is critical (ref. 31).

( a ) Ribbon-shaped fiber.

(17)

1ctr.cllohd ,Jihcr..

Figure 16. Noncircular PAN-based carbon fibers produced with melt-spun technology (ref. 31). 64

10 -

5 $a

a, E$

zg

78.5 percent hollow fibers

8 -

f2 "E 6 -

D m .-E

%a 01

5s J=U

5: .o c ,o

4 -

35 percent hollow fibers

CC Materials and Composites Pitch-Based Carbon Fibers Diwan (ref. 33) investigated the effect of noncircular pitch-based carbon fibers on the performance of polymeric composites. A solvent-extracted mesophase was melt-spun into trilobal and hexalobal carbon fibers. After thermosetting and carbonization, these fibers were used to form unidirectional carbon-epoxy composite specimens. These specimens were tested in compression using a standard Illinois Institute of Technology Research Institute (IITRI) fixture. Results indicated that, although the tensile strength of the two fiber shapes was nearly identical, the compressive strength of the composites reinforced with the more highly noncircular hexalobal fibers was 16 percent higher than those reinforced with the trilobal fibers. The probable cause for this improvement in compression strength was an increased resistance to fiber buckling caused by the higher moment of inertia of the more noncircular fiber. Thus, the use of noncircular shapes may result in improved compressive properties for both PAN-based and pitch-based carbon fibers. Also, in CC composite applications, the increased surface-to-volume ratio of noncircular shapes may even yield improved interlaminar shear strength (refs. 31 and 33).

Summary The microstructure of the polymeric precursor fibers used to form PAN-based and rayon-based carbon fibers is fibrillar, but mesophase precursor fibers are composed of extended domains of a highly oriented structure. Since, to a large extent, the final microstructure of the carbon fiber replicates that of the precursor fiber, in PAN-based and rayon-based carbon fibers the graphite crystallites are arranged into a fibrillar substructure, whereas pitch-based carbon fibers have an extended graphitic layer structure. Although these larger regions of graphitic structure allow pitch-based carbon fibers to exhibit extremely high moduli, this structure also makes pitch-based carbon fibers more flaw-sensitive, accounting for their lower tensile strengths. However, modifying the flow profile during fiber extrusion has been shown to change the microstructure significantly and even create new microstructures in pitch-based fibers, improving both the tensile and the compressive strength of these fibers. Also, it appears that noncircular fiber cross sections can increase the buckling resistance of both pitch-based and PANbased carbon fibers.

Acknowledgments The authors thank G. P. Daumit of BASF Structural Materials, Incorporated, for generously providing the original photographs used for figure 16. In addition, Dr. D. J. Johnson and IOP Publishing Limited are to be thanked for graciously permitting figures 2, 4, and 5 to be reproduced. Thanks are also due to the Society of Plastics Engineers, Chapman and Hall, Limited, Pergamon Press, Incorporated,

66

Effect of Microstructure and Shape on Carbon Fiber Properties and Kluwer Academic Publishers for permission to reproduce figure 3, figure 7, figure 8, and figures 1 and 11, respectively. Finally, the authors thank the Materials Research Society and T. Hamada for permission to reprint figure 8.

References 1. Riggs, D. M.; Shuford, R. J.; and Lewis, R. W.: Graphite Fibers and Composites. Handbook of Composites, George Lubin, ed., Van Nostrand Reinhold Co., c. 1982, pp. 196-271. 2. Bacon, R.: Carbon Fibers From Rayon Precursors. Chemistry and Physics of Carbon, Volume 9, P. L. Walker and P. A. Thrower, eds., Marcel-Dekker, Inc., 1973, pp. 1-102. 3. Singer, L. S.: High Modulus High Strength Fibers Produced From Mesophase Pitch. U.S. Patent 4,005,183, Jan. 25, 1977. 4. Riggs, D. M.; and Diefendorf, R. J.: Forming Optically Anisotropic Pitches. U.S. Patent 4,208,267, June 17, 1980.

5. Edie, D. D.; and Dunham, M. G.: Melt Spinning Pitch-Based Carbon Fibers. Carbon, vol. 27, no. 5, 1989, pp. 647-655. 6. Edie, D. D.: Pitch and Mesophase Fibers. Carbon Fibers Filaments and Composites, J. Figueiredo, C. A. Bemardo, R. T. K. Baker, and K. J. Huttinger, eds., Kluwer Academic Publ. (Dordrecht, The Netherlands), 1990, pp. 43-72. 7. Fitzer, Erich, ed.: Carbon Fibres and Their Composites. Springer-Verlag, 1985.

8. Blakslee, 0. L.; Proctor, D. G.; Seldin, E. J.; Spence, G. B.; and Weng, T.: Elastic Constants of Compression-Annealed Pyrolytic Graphite. J . Appl. Phys., vol. 41, no. 8, July 1970, pp. 3373-3382. 9. Williams, Wendell S.; Steffens, D. A.; and Bacon, Roger: Bending Behavior and Tensile Strength of Carbon Fibers. J . Appl. Phys., vol. 41, no. 12, Nov. 1970, pp. 4893-4901. 10. Seldin, E. J.: Mechanical Properties of Graphite-Review. Proceedings of the Ninth Biennial Conference on Carbon, 1969, p. 59. 11. Jones, Janice Breedon; Barr, John B.; and Smith, Robert E.: Analysis of Flaws in High-Strength Carbon Fibres From Mesophase Pitch. J . Mater. Sci., vol. 15, no. 10, Oct. 1980, pp. 2455-2465.

67

CC Materials and Composites 12. Reynolds, W. N.; and Moreton, R.: Some Factors Affecting the Strengths of Carbon Fibres. Philos. Trans. Royal SOC.London, ser. A, vol. 294, no. 1411, Jan. 21, 1980, pp. 451461. 13. Johnson, D. J.: Structure-Property Relationships in Carbon Fibres. J . Phys. D: Appl. Phys., vol. 20, no. 3, Mar. 14, 1987, pp. 286-291. 14. Diefendorf, R. J.; and Tokarsky, E.: Polym. Eng. Sei., vol. 15, no. 3, 1975, pp. 150-159. 15. Endo, Morinobu: Structure of Mesophase Pitch-Based Carbon Fibres. J . Mater. Sei., vol. 23, no. 2, Feb. 1988, pp. 598-605. 16. Hamada, T.; Nishida, T.; Sajiki, Y.; and Matsumoto, M.: Structure and Physical Properties of Carbon Fibers From Coal Tar Mesophase Pitch. J . Mat. Res., vol. 2, no. 6, 1987, pp. 850-857. 17. Hamada, T.; Nishida, T.; Furuyama, M.; and Tomioka, T.: Transverse Structure of Pitch Fiber From Coal Tar Mesophase Pitch. Carbon, vol. 26, no. 6, 1988, pp. 837-841. 18. Whitney, W.; and Kimmel, R. M.: Griffith Equation and Carbon Fibre Strength. Nature, vol. 237, no. 75, June 5, 1972, pp. 93-94. 19. Kowalski, Ian M.: New High Performance Domestically Produced Carbon Fibers. Advanced Materials Technology '87-Volume 32 of International SAMPE Symposium and Exhibition, Ralph Carson, Martin Burg, Kendall J. Kjoller, and Frank J. Riel, eds., SOC.for the Advancement of Material and Process Engineering, 1987, pp. 953-963. 20. Product Data Sheet. E Series Pitch-Based Carbon Fibers, E. I. du Pont de Nemours and Co., 1989. 21. Nazem, F.: Process for Controlling the Cross-Sectional Structure of Mesophase Pitch Derived Fibers. U.S. Patent 4,376,747, Mar. 15, 1983. 22. Edie, D. D.; Fox, N. K.; Barnett, B. C.; and Fain, C. C.: Melt-Spun NonCircular Carbon Fibers. Carbon, vol. 24, no. 4, 1986, pp. 477482. 23. Kumar, Satish: Structure and Properties of High Performance Polymeric and Carbon Fibers-An Overview. SAMPE Q., vol. 20, no. 2, Jan. 1989, pp. 3-8. 24. Kumar, Satish; and Helminiak, T. E.: Compressive Strength of High Performance Fibers. SAMPE J., vol. 26, no. 2, 1990, pp. 51-61.

68

Effect of Microstructure and Shape o n Carbon Fiber Properties 25. Dobb, M. G.; Johnson, D. J.; and Park, C. R.: Compressive Behaviour of Carbon Fibres, J . Mater. Sci., vol. 25, 1990, pp. 829-834. 26. Kaverman, Richard: New Shapes Give Glass Fibers More Strength. Res. & Dev., Mar. 1989, p. 20. 27. Le Maistre, Christopher William: The Origin of Structure in Carbon Fibers. Ph.D. Diss., Rensselaer Polytechnic Inst., 1971. 28. Tokarsky, E. W.: The Relationships of Structure to Properties in Carbon Fibers. Ph.D. Thesis, Rensselaer Polytechnic Inst., 1973. 29. Johnson, John W.: Factors Affecting the Tensile Strength of Carbon-Graphite Fibers. Appl. Polymer Symp., no. 9, 1969, pp. 229-243. 30. Gainey, H. E.; Handlos, A. A.; Edie, D. D.; Kennedy, J. M.; and Fain, C. C.: Flaw Distributions in Noncircular Carbon Fibers. FIBER-TEX 1988, John D. Buckley, ed., NASA CP-3038, 1989, pp. 163-167. 31. Daumit, G. P.; and KO, Y. S.: A Unique Approach to Carbon Fiber Precursor Development. High Tech-The Way Into the Nineties, K. Brunsch, H. D. Golden, and C. M. Herhert, eds., Elsevier Science Publ. B. V. (Amsterdam), 1986, pp. 201-213. 32. Niederstadt, Gunter: Moglichkeiten zur Erhohung der Biegesteifigkeit von CFK-Schichtstoffen mittels Hohlfasereinlage. Z. Flugwiss. und Weltraumforsch., Band 5 , Helf 1, 1981, pp. 30-36. 33. Diwan, P.: The Effect of Fiber Shape on the Compressive Properties of Composites. M. S. Thesis, Clemson Univ., 1985.

69

Chapter 4 Textile Preforms for Carbon-Carbon Composites Frank K. KO Drexel University Philadelphia, Pennsylvania

Abstract 72 Introduction 72 Classification of Preforms 72 Linear Fibrous Assemblies 74 Fabric Preforms 75 Structural Geometry of 2-D Fabrics 78 Woven Fabrics 78 Knitted Fabrics 78 Braided Fabrics 80 Structural Geometry of 3-D Fabrics Woven 3-D Fabrics 82 Orthogonal Nonwoven Fabrics

82

84

Knitted 3-D Fabrics 85 3-D Braided Fabrics 89 Structure and Properties of Textile-Reinforced CCC 89 Modeling of Textile Structural Composites 93 Concluding Remarks 99 Acknowledgments

100

References 100 Bibliography 104

71

CC Materials and Composites

Abstract Textile preforms for carbon-carbon composites (CCC) are reviewed in this paper. From the structural geometry point of view, the various levels of fiber architecture are classified into linear, planar (2-D), and 3-D fibrous assemblies. The role of fiber architecture in the processing and strengthening of CCC is discussed. To provide a basis for the mechanistic analysis of CCC reinforced by textile structures, unit cell-based modeling methods are reviewed.

Introduction Textile preforming is the method of placing reinforcing fibers in a desired arrangement prior to formation of a composite structure. Starting with linear assemblies of fibers in continuous and/or discrete form, these micro-fibers can be organized into two-dimensional (2-D) and three-dimensional (3-D) structures by means of textile processes such as interlacing, intertwining, or interlooping. Properly selecting the geometry and the method of placement or geometric arrangement of the fibers can tailor the resulting structural performance of the composite. These fiber placement methods create textile preforms that possess a wide spectrum of pore geometries and pore distribution; a broad range of structural integrity and fiber volume fraction; and fiber orientation distribution as well as a wide selection of formed-shape and net-shape capabilities. In linear form, the carbon threads can serve as stitch yams for stitched structures. These linear structures also can be used as fasteners. Planar systems are suitable for skin structures, although 3-D structures find varying uses ranging from rocket nozzles to large-scale structural components for hypersonic vehicles. Combining with high-performance fibers, matrices, and properly tailored fiber/matrix interfaces, fiber architecture promises to expand the design options for tough and reliable structural CCC. With an integrated network of structural cells in two- and three-dimensional arrangements, textile structures not only provide a mechanism for structural toughening of composites but also facilitate composite processes into net or near net-shape structural parts.

Classification of Preforms On the basis of structural integrity and fiber linearity and continuity, fiber architecture can be classified into four categories: discrete, continuous, planar interlaced (2-D), and fully integrated (3-D) structures. In table I, the nature of the various levels of fiber architecture is summarized (ref. 1). The first category of fiber architecture is a discrete fiber system, such as a whisker or fiber mat, which has no material continuity. The orientation of the fibers is difficult to control precisely, although some aligned discrete fiber systems

72

Level I I1 I11 IV

Reinforcement system Discrete Linear Laminar Integrated

Textile construction Chopped fiber Filament yam Simple fabric Advanced fabric

Fiber length Discontinuous Continuous Continuous Continuous

Fiber Fiber orientation entanglement Uncontrolled None Linear None Planar Planar 3-D 3-D

CC Materials and Composites

Linear Fibrous Assemblies Yams and rovings are linear fibrous assemblies composed of discrete or continuous fibers. The basic differences between rovings and yams are listed:

1. Rovings are large fiber bundles with little or no twist. 2. Yams are finer fiber bundles with some twist. Yarns that are composed of discrete fibers are called staple yarns; yams having continuous fibers are continuous-filament yams. The majority of high-performance yarns are continuous-filament yams having single or multiple strands. For example, Avco SCS fibers are available in single-strand and are referred to as monofilaments. The AS4 12-K yams composed of 12 000 filaments are referred to as multifilaments. Figure 2 shows the geometric features of filament yarns. The technology for converting short fibers into yam assemblies is well established (ref. 2 ) . For high-modulus fibers that exist in staple form (such as single crystal whiskers and asbestos), special yam formation techniques have been developed. For example, researchers at Los Alamos National Laboratory (fig. 3) are growing long S i c whiskers (75 mm to 100 mm in length by 3 pm to 10 pm in diameter) for spinning into staple yams (ref. 3). Similar methods may be applicable for the vapor-grown carbon fibers developed at the General Motors Research Laboratory Linear

3-D

Planar (2-D)

& .. .

Biaxial Triaxial

+

Biaxial Triaxial

Core

M LI It if ilame nt Monofilament

Flat

74

6-ply 4.~1~ impaled impaled

Textured Twisted

Figure 1. Classification of fiber architecture.

Textile Preforms for Carbon-Carbon Composites

0

j

Monofilament

Untwisted filament yarn

Twisted filament yarn

High-bulk filament yarn

Figure 2. Structural geometry of filament yarns.

(ref. 4). Through proper selection of whisker length and yam twist level, the pore structure and mechanical properties of the staple yam can be engineered for weaving into 2-D or 3-D fabrics for CCC. Motivated by the need for finer carbon yams, Heltra, a division of Courtaulds, has also developed a staple yam formation method using a stretch breaking process. Compensating for the inevitable reduction in tensile strength, the protruded fibers on the transverse direction of the staple yams (as shown schematically in fig. 4(a)), can potentially improve the through-the-thickness strength and shear resistance and also the toughness of the CCC. This improvement in properties is derived from the subtle interlocking of the protruding fiber creating a network of in-plane and out-of-plane crack arrestors. An example of a yam manufacturing process that produces protruding fibers in the transverse direction is the family of hispiduous yams created by the chenille process (fig. 4(b)).

Fabric Preforms A fabric is defined as an integrated fibrous structure produced by fiber entanglement of yam interlacing, interlooping, intertwining, or multiaxial placement. Fiber felts, composed of fabrics formed directly from fibers, are an example of fiber-to-fabric structures. This research concentrates on yam-to-fabric structures. Table I1 compares the four basic yam-to-fabric formation techniques, and figure 5 shows examples of fiber architecture created by these techniques.

75

CC Materials and Composites Staple yarn

Warm water

Whisker ribbon

(filter paper) Figure 3 . Spinning of staple yarns from whiskers.

(a)

(b)

Figure 4. Structural geometry of ( a ) spun and (b) hispiduous yarns.

76

Basic direction of yarn introduction Two (0°/900) (warp and fill)

Basic fabric formation technique Interlacing (by selective insertion of 90' yarns into 0' yarn system)

Braiding

One (machine direction)

Intertwining (position displacement)

Knitting

One (0' or 90') (warp or fill)

Interlooping (by drawing loops of yarns over previous loops)

Nonwoven

Three or more (orthogonal)

Mutual fiber placement

Weaving

C C Materials and Composites of three or more yams in the thickness direction, is a fibrous network wherein yams pass from surface to surface of the fabric in all three directions. The key criteria for selecting the fiber architecture for structural composites are: the capability for in-plane multiaxial reinforcement, the through-the-thickness reinforcement, and the capability for formed-shape and/or net-shape manufacturing. Depending on the processing and end-use requirements, some or all of these features are required. In this section, the representative structural geometry of 2-D and 3-D fabrics is introduced.

Structural Geometry of 2-D Fabrics Woven Fabrics The interlacing of yams fabricates the hundreds of possible woven fabric combinations. From the in-plane fiber orientation, woven fabrics can be divided into biaxial and triaxial woven structures. Biaxial weaves consist of 0' and 90' yams interlaced in various repeating patterns or topological unit cells. The three basic weave geometries from which many other patterns evolve are the plain, twill, and satin weaves. A schematic diagram of various views of these three basic weaves is shown in figure 6. The frequency of yam interlacing and the linearity of the yam segments distinguish these three fabrics. The plain weave has the highest frequency of yam interlacing, whereas the satin weave has the least number of yam interlacing, with the twill weave somewhere in between. Accordingly, the plain weave has a higher level of structural integrity and greater ductility because of the crimp geometry produced by yam interlacings. On the other hand, the satin weave has the highest level of fiber-to-fabric strength and modulus translation efficiency because of the low level of yam interlacing and yam linearity. The low level of yam integration in satin weave also allows freedom of yarn mobility, which contributes to higher fiber packing density and, consequently, higher levels of fiber volume fraction. Although cane weaving for cane chairs has existed for a long time, machinemade triaxially woven fabrics weren't available until Norris Dow 's development of the triaxial weave in the early 1970's (ref. 5). The unique feature of triaxial weave is the 90f60' hexagonal yam orientation in one plane, resulting in a high level of in-plane shear resistance. High levels of isotropy and dimensional stability can be achieved with triaxial weave at low fiber volume fraction. Figures 7 and 8 show a schematic diagram of two triaxial weave geometries.

Knitted Fabrics Knitted fabrics are interlooped structures wherein the knitting loops are either produced by the introduction of the knitting yam in the cross-machine direction (weft knit) or along the machine direction (warp knit). As shown in figures 9 and 10, knitting can produce a large number of stitch geometries. By controlling the

78

Textile Preforms for Carbon-Carbon Composites 4

!I

ii

1

4iwiim

7 .Jaxn

(a)

(c)

.$!

rn

1 8 _ _

(b)

Figure 6. Structural geometry of biaxially woven fabrics: ( a )plain weave, (b) twill weave, and ( c ) satin weave. a

Figure 7. Structural geometry of triuxiully woven fabrics (basic weave).

79

CC Materials and Composites

Figure 8. Structural geometry of triaxially woven fabrics (biplain weave). stitch (loop) density, a wide range of pore geometry can be generated. Because of the nature of the interlooped structure, the maximum fiber packing density of knitted structures is lower than that of the woven fabrics. The severe bending of yarns during the knitting process also discourages converting ceramic yarns to knitted structures. However, if a creative combination of ceramic sewing threads or very fine yarns is used to form the stitches, knitted fabrics may be used effectively as a base structure for the incorporation of Oo and/or 90° yarns (fig. 11). The resulting preform is an integrated structure that combines the high conformability of the knitted base structure with a high level of directional reinforcement from the straight lay-in yarns.

Braided Fabrics By intertwining three or more yarn systems together, braided fabrics can be produced in flat or tubular form (fig. 12). The bias interlacing nature of the braided fabrics makes them highly conformable, shear resistant, and tolerant to impact damage. To enhance reinforcement in the 0' direction, triaxial braiding can be used to introduce 0&6' yarns as shown in figure 13. Although braiding and filament winding bear some similarities, subtle differences exist between these two processes and the resulting structures. Table I11 provides a comparison of braiding and filament winding. In reference 6, the subject of braiding is treated in greater detail.

80

Textile Preforms for Carbon-Carbon Composites

+

4-

Mise or float

Tuck

Cross

Figure 9. Weft-knit constructions.

Plain tricot

Sharkskin

Locknit

Queenscord

Velour

Figure IO. Warp-knit constructions.

81

Filament winding 200 to 900, (fiber orientation)

Braiding 0' to 80°,

Maximum fiber volume fraction V f

.80

.68

Fiber placement

Mandrel rotation Uniaxial/ply Linear

Fiber rotation Multiaxial/ply Nonlinear (crimp)

Structural integrity

Noninterlaced

Interlaced

Impact damage resistance

Poor

Excellent

Data base

Well established

Limited

Production rate

100 to 700 lb/hr

60 to 100 lbhr

Product size

Unlimited

70-in. dia (144 carrier)

Product shape

Problems with concave contour

Very comformable

Fiber orientation

+/-e

+/-e

Textile Preforms for Carbon-Carbon Composites

JrwYTY?Y J U U n U n L

\ I

\

I

\

I

\

/

J \n Id \n ln \a IU \ l

, 6 3 . 3 u w m L

q

0W-L

\ / \ I \ / \ /

J n u

\

/

\

u-363-

/

\

l

\

l

‘i )

< H , \ l \ /

Jn#+U=?!?!

(a)

n u nIn \x u n LI JH U $+ U nl n W a pL

J

!A <:iA
(c)

% 1

*#g$fjf@-L&

+ + f & + $ & U j Z ?

+f& - % p& L7 !+ 5p Jw nn-

m J mLJ m k / m

JrJ&n~&~

@ a @ (d)

Figure 11. Weft- and warp-insertion knits: ( a ) weft knit, (b) warp-insertion weft knit, and ( e ) and (d) weft and warp insertion.

Figure 12. Structural geometry of 2 - 0 braid.

83

CC Materials and Composites

Figure 13. Structural geometry of triaxial braid.

panels (fig. 14(b)), core structures simulating a box beam (fig. 14(c)), or truss-like structures (fig. 14(d)). Furthermore, by properly manipulating the warp yams, as exemplified by the angle interlock structure (fig. 14(e)), the through-the-thickness yams can be organized into a diagonal pattern. To address the inherent lack of in-plane reinforcement in the bias direction, new progress is being made in triaxial weaving technology by Dow et al. (ref. 7) to produce multilayer triaxial fabrics, as shown in figure 15.

Orthogonal Nonwoven Fabrics Aerospace companies such as General Electric (ref. 8) pioneered nonwoven 3-D fabric technology, which was developed further by Fiber Materials Incorporated (ref. 9). Recent progress in automating the nonwoven 3-D fabric manufacturing process was made in France by Aerospatiale (refs. 10 and l l ) , Brochier (ref. 12), and SEP (ref. 13) and in Japan by Fukuta of the Research Institute for Polymers and Textiles (refs. 14 and 15). The structural geometries resulting from the various processing techniques are shown in figure 15 (ref. 16). Figures 16(a) and 16(b) show the single bundle XYZ fabrics in a rectangular and cylindrical shape. Figure 16(c) demonstrates the multiple-yam bundle possibilities in the various directions; figure 16(d) shows the multidirectional reinforcement in the 3-D structure plane. Although most of the orthogonal nonwoven 3-D structures consist of linear yam reinforcements in all of the directions, introduction of the planar yams in a nonlinear manner, as shown in figures 16(e), 16(f), and 16(g), can result in either an open lattice structure or a flexible and conformable structure.

84

Textile Preforms f o r Carbon-Carbon Composites

(a

(b) (e) Figure 14. Structural geometry of 3-0 woven fabrics: (a) solid orthogonal panels, ( b ) variable-thickness solid panels, ( e ) core structures, (d) trusslike structures, and (e) angle interlock structure.

Knitted 3-0 Fabrics The knitted 3-D fabrics are produced either by the weft knitting or warp knitting process. An example of a weft knit is the near-net shape structure knitted under computer control by the Pressure Foot@process (ref. 17) (fig. 17). In a collapsed form, this preform has been used for carbon-carbon aircraft brakes. The unique feature of the weft-knit structures is their conformability (ref. 18). When additional reinforcement is needed in the 0' and 90' directions, linear laid-in yarns can be placed inside the knitting loops as shown in figure 18. The most undesirable structural reinforcement feature of weft-knit structures is their bulkiness, which leads to the lowest packing density or lowest level of maximum fiber-volume fraction compared with other fabric preforms. Although the weft-knitted structures have applications in limited areas, the multiaxial warpknit (MWK) 3-D structures are more promising, and they have undergone more development in recent years (refs. 19 and 20).

85

CC Materzals and Composates

Figure 15. Multilayer, triaxially woven fabric (ref. 7).

x

& & X % /

(4

y

2

(b)

(9) Figure 16. Complex structural shapes by 3-0 braiding: (a) rectangular shape, (b) cylindrical shape, ( c )multiple-yarn bundle possibilities, ( d ) multidirectional reinforcement in plane of three-dimensional structure, and (e),(f), and (g)planes moving in nonlinear structure.

86

Textile Preforms for Carbon-Carbon Composites

Figure 17. Knitted 3-0 fabrics.

Figure 18. Three-bar weft-insertion warp knit.

87

CC Materials and Composites

&@g oo

Chain

Tricot

Unit cell

Figure 19. Multiaxial warp-knit fabric. The structural geometry of the MWK fabric systems consists of warp (OO), weft (90°), and bias (,to) yams held together by a chain or tricot stitch through the thickness of the fabric, as illustrated in figure 19. The major distinctions of these fabrics are the linearity of the bias yams, the number of axes, and the precision of the stitching process. Depending on the number of guidebars available and the yaminsertion mechanism, the warp-knit fabric can consist of predominantly uniaxial, biaxial, triaxial, or quadraxial yams. The latest commercial nonimpaled MWK fabric is produced by the Mayer Textile Corporation using a multiaxial magazine, weft-insertion mechanism. The attractive feature of this system is the precision of yam placement with four layers of linear or nonlinear bias yams plus a short fiber mat arranged in a wide range of orientations. Furthermore, stitches are formed without piercing through the reinforcement yams (hence the term, "nonimpaled") at a production rate of 100 mhr. The latest development in the impaled MWK is the LIBA or Hexcel system, shown in figure 20. Six layers of linear yams can be assembled in various stacking sequences and stitched together by knitting needles piercing through the yam layers. While this piercing action unavoidably damages the reinforcing fiber, the powerful knitting needles incorporate a fiber mat as a surface layer for the composite.

88

Textile Preforms for Carbon-Carbon Composites

Knitting yarns

Warp inlay yarns

Figure 20. LIBA impaled MWK system. 3 - 0 Braided Fabrics

The 3-D braiding technology is an extension of the well-established 2-D braiding technology wherein the intertwining or orthogonal interlacing of two or more yarn systems to form an integral structure constructs the fabric. The 3-D braiding is one of the textile processes in which a wide variety of solid complex structural shapes can be produced integrally resulting in a highly damageresistant structural preform. Figure 21 shows two basic loom setups in circular and rectangular configurations (ref. 21). The 3-D braids are produced by a number of processes including the Track and Column (3-D circular loom) method (ref. 22) (fig. 22), the two-step braiding method (refs. 23 and 24) (fig. 23), and a variety of displacement braiding techniques. The basic braiding motion includes the alternate X and Y displacement of yarn carriers followed by a compacting motion. The proper positioning of the carriers and the joining of various rectangular groups through selected carrier movements accomplish shape formation. Examples of the structural shapes successfully demonstrated in the Fibrous Materials Research Laboratory at Drexel University are shown in figure 24.

Structure and Properties of Textile-Reinforced CCC The properties of carbon-carbon are not well publicized. For structural applications, the goal is to produce CCC having a density level of 1.8-2.0 g/cc. Increased density calls for a careful selection of yarn bundle size and fiber architecture. A fine weave is more desirable in obtaining high density, although an interconnected, three-dimensional fiber network is preferred for a chemical vapor deposition (CVD) system. As the density of the composite increases, strength and modulus are expected to increase as well. Because of the low strength of the carbon

89

CC Materials and Composites

Figure 21. Basic loom setups for circular and rectangular configurations.

Figure 22. Atlantic Research Corporation’s Track and Column 3 - 0 circular loom.

90

Textile Preforms f o r Carbon-Carbon Composites

Figure 23. Two-step braiding method. matrix (less than 40 MPa) and the weak interface between the fiber and matrix, one of the key issues in the area of CCC is the improvement of composite throughthe-thickness strength. For structural applications, most reinforcement preforms are 3-D fabrics. The article by McAllister and Lachman (ref. 25) and the book by Tampol’ski et al. (ref. 26) provide an excellent summary of the structure and properties of 3-D fabric-reinforced CCC. In a review article by KO (ref. 16), the properties of 3-D fabric-reinforced carbon matrix composites were reviewed and summarized as shown in table IV.

In Table IV, p is density, V f is fiber volume fraction, at is tensile strength, Et is tensile modulus, uc is compressive strength, E, is compressive modulus of the composite, a f is strength of the fiber, E f is elastic modulus of the fiber, T is shear strength, and a is coefficient of thermal expansion. The dependence of mechanical properties on fiber architecture can be illustrated in table V.

91

Property p, g/cc

Vf

gt,MPa

Et,GPa g,-,MPa

E,,GPa g f ,MPa

Ef,GPa r , MPa a, in/in°C References

Nonwoven 1.62.0 0.45 40-200 11-60 40-220 1245 60-150 10-30 20-40 0.06% 27, 12, 2 8 4 2

3-D woven

3-D braid

-

-

-

0.60 193-372 103-138

18.7-62 7.5-34.5 17.5-82.8 4.841.4 26.7-89.7 4.627.6

-

-

-

-

27

27,+

-

Textile Preforms for Carbon-Carbon Composites

It is interesting to note that the reduction in compacting frequency from weave pattern A to B resulted in a substantial increase in strength and modulus without a significant reduction in shear strength.

Modeling of Textile Structural Composites The mechanical properties of textile-reinforced composites can be predicted with a knowledge of the fiber properties, matrix properties, and fiber architecture through a modified laminate theory approach. Geometric unit cells defining the fabric structure (or fiber architecture) can be identified and quantified to form a basis for the analysis. For 2-D woven fabric-reinforced composites, Dow et al. (ref. 7) and Yang et al. (ref. 43) have developed models for the thermomechanical properties of plain-, twill-, and satin-reinforced composites. Examples of these include the mosaic, crimp, and bridging models developed by Yang et al. (ref. 43). In the mosaic model, fiber continuity is ignored and the composite is treated as an assembly of cross-ply elements. With the crimp model, the nonlinear crimp geometry and the yam continuity are considered. Based on the geometric repeating unit cell, each yarn segment is treated as a laminae. Although the crimp model was found to be suitable for plain weave composites, the bridging model was found to be best for satin weave composites because it takes the relative stiffness contribution of the linear and nonlinear yarn segments into consideration. The modeling of 3-D fiber-reinforced composites also begins with the establishment of geometric unit cells. A summary of the four major classes of 3-D fiber architecture is given in figure 25. Three basic yam components make up the unit cell for a 3-D orthogonal nonwoven fabric, defined according to the yam orientation: 0' (warp), 90' (weft), and through-the-thickness. Five basic yam components make up the unit cell for a general MWK fabric, defined according to the yarn orientation: 0' (warp), 90' (weft), 45' (bias), -45' (bias), and the stitching yam (through thickness). The fractional volume of fiber in each of the directions can be calculated geometrically on the basis of yarn size, yam Spacing, and stitch construction that should reflect a combination of the orthogonal unit cell and crimp geometry. Several yams running parallel to the body diagonal of the cell represent the unit cell for the 3-D braid. However, in some instances, yams are placed in longitudinal (0') and transverse (90') directions of the fabric and are referred to respectively as longitudinal and transverse reinforcements (or lay-ins). Among the 3-D composites, the 3-D braided composites have received much attention because of their improved stiffness and strength in the thickness direction, their delamination-free characteristics, and their near-net shape manufacturing capabilities. Several analytic models have been developed to characterize the elastic moduli and structural behavior of 3-D braided composites.

93

Weave A(l x 1 x 1) B(l x 1 x 3)

Tensile strength 28 x lo3 psi 54 x io3 psi

Elastic modulus 15 x lo6 psi 20 x lo6 psi

Shear strength 3.1 x lo3 psi 2.9 x lo3 psi

Textile Preforms for Carbon-Carbon Composites Yang et al. (see ref. 43) also developed a Fiber Inclination Model according to the idealized, zig-zagging yarn arrangement in the braided preform (ref. 43). They assumed an inclined lamina as a representation of one set of diagonal yams in a unit cell. In this way, four inclined, unidirectional laminae formed a unit cell. Using classical laminate theory, the elastic moduli then can be expressed in terms of the laminae properties. From the preform processing science aspect, KO et al. (ref. 45) developed a fabric geometric model (FGM) based on the unit cell geometry shown in figure 26. The stiffness of a 3-D braided composite was considered to be the sum of stiffnesses of all its laminae. The unique feature of the FGM is its ability to handle 3-D braid and other multidirectional reinforcements including 5-D, 6-D, and 7-D fabrics with straight or curvilinear yams. The product of the FGM is a stiffness matrix that provides a link between applied strains and the corresponding stress responses. With a properly selected failure criterion, the stress-strain relationship of the fiberreinforced composite can be predicted. The development of the FGM has been detailed by KO (refs. 6, 16, and 45); only a brief introduction is presented here. The first step in the modeling process is defining fiber orientation 6 and fiber volume fraction Vf in terms of the preform processing parameters U , V, W, IVY, and D,. The preform processing parameters U , V, and W are track displacement, column displacement, and combing frequency, respectively; IVY is the number of yams in the cross section and D, is the yarn linear density (denier)* from which the total fiber cross-sectional density can be computed. For a given fiber with density p and a give composite cross-sectional area A,, the fiber orientation 6 and fiber volume fraction V f for a (U x V x W ) 3-D braid composite structure are shown (ref. 44):

o=tan-l(

Vf=

V FwT V

)

(1)

NyDysin(tan-'[K* tan{O}/k]) A , . 9 x 105ptan(6)k*

(2)

where

d K=(

m

w

.)

and

K* =

J1+K2

* 1 denier = 1 gm/9000 meters. Specific cross-sectional area of the fiber A,

= No.

Of denier cm2. 9 x 105 x p

95

CC Materials and Composites

Figure 26. Idealized yarn segment in unit cell. With the knowledge of fiber orientation 8, the stiffness matrix [Ci]for each composite yarn segment in the unit cell can be transformed from the stiffness matrix [C]of equivalent unidirectional composites, as follows:

[GI = [~21[C1[Tilt

(3)

The systems stiffness matrix [Cs] for the composite subsequently can be determined by the summation of the according to their fractional volume fraction contribution Ki:

[ci]

[GI = C(Ki[C,l)

(4)

With the material properties and geometry of the 3-D braided composite defined piecewise and linearly, the stress response of the composite system can be determined for each incremental strain {A€} imposed to produce the stressstrain relationship: { A d s , i = [Csl{Ad { d s , i = {dS,Z-l

+ {A&

(5)

where { Am}s is the incremental stress response of the composite system and {a}+ is the total stress on the composite system.

96

Find strain on composite, build system stiffness matrix, C Calculate stress on the composite O = CE

-

Plot ( E d

output

100 -

2

80

w vi

60 -

0 ic

-3 3

-

nonwoven

-0

E -

40 -

-p

20-

F

6

3

._ 0

-I

OO

11 -

m

b

10-

OX vi

9-

B

8-

5

7-

I

I

I

I

I

20

40

60

80

100

XYZ

2

-3 3 E

c v)

2ril -

6-

s

5-

4

4

I

I

I

I

Textile Preforms f o r Carbon-Carbon Composites relationship between G,, and E, for the linear and curvilinear reinforcement systems. For the curvilinear system, G,, increases as E, increases. This increase reflects the respective sensitivity of the linear and curvilinear-reinforced systems to fiber volume fraction distribution and fiber orientation.

Concluding Remarks Textile preforms have much to offer in the toughening and manufacture of nextgeneration high-performance structural composites. With a large family of highperformance fibers, linear fiber assemblies, and 2-D and 3-D fiber architectures, a wide range of composite structural performances may be tailored to meet specific requirements.

An examination of the literature indicated that only a limited number of systematic studies have been carried out on fabric-reinforced CCC’s (carbon-carbon composites). A well-established data base is needed to stimulate the usage of fabric-reinforced CCC’s for structural applications. The literature suggests a trend toward using 3-D fiber architecture for CCC structural toughening which poses important technical challenges. The first challenge is the question of converting high-modulus yarns to textile structures. The processing difficulty with brittle fibrous structures calls for an innovative combination of materials systems such as the concept of material and geometric hybridization. The infiltration or placement of matrix material in a dense, 3-D fiber network also creates new challenges and demands an understanding of the dynamics of the process-structure interaction. Questions that must be answered relate to the optimum pore geometry for matrix infiltration, the pore distribution, and the bundle size. As the level of fiber integration increases, the opportunity of fiber-to-fiber contact intensifies at the crossover points. Guidance is required to select the fiber architecture and matrix placement method best suited to reduce the incidence of localized fiber-rich areas.

To take advantage of the attractive features that textile structural composites offer, a sound data base and design methodologies need to be developed. The fabric geometry models developed so far establish a necessary, but not entirely complete, first step in the modeling of CCC’s. Future work in the modeling of fabric-reinforced CCC’s requires a better understanding of the dynamic interaction among fiber, matrix processing condition, and fiber architecture.

99

CC Materials and Composites

Acknowledgments Much of the work on fiber architecture reported herein, and specifically 2-D and 3-D fiber architecture, has been supported by the Office of Naval Research and the Air Force. The assistance provided by Mitchell Marmel of the Fibrous Materials Research Center in the preparation of this manuscript is greatly appreciated.

References 1. Scardino, F. L.: Introduction to Textile Structures. Textile Structural Composites, T. W. Chou and F. K. KO, eds., Elsevier, 1989, pp. 1-26. 2. Goswami, B. G.; Martindale, J. G.; and Scardino, F. L.: Textile Yarns, Technology, Structure and Applications. John Wiley & Sons, 1977, pp. 273-337. 3. Gac, F. C.: S i c Whisker Staple Yarn Development. Mater. & Process. Rep., vol. 3, no. 1, Apr. 1988, pp. 3-4. 4. Tibbets, G . G.: Vapor Grown Carbon Fibers: Status and Prospects. Res. Publ. GMR-6454, General Motors, Oct. 1988.

5. Dow, N. F.; and Tranfield, G.: Preliminary Investigations of Feasibility of Weaving Triaxial Fabrics (Dow Weave). Textile Res. J . , vol. 40, no. 11, Nov. 1970, pp. 986-998. 6. KO, Frank K.: Braiding. Composites, Volume 1 of Engineered Materials HandbookTM, ASM InternationalTM, c.1987, pp. 519-528. 7. Dow, Noms F.; Ramnath, V.; and Rosen, B. Walter: Analysis of Woven Fabrics for Reinforced Composite Materials-Technical Final Report. NASA CR-178275, 1987. 8. Stover, E. R.; Marx, W. C.; Markowitz, L.; and Mueller, W.: Preparation of an

Omniweave-Reinforced CarbonlCarbon Cylinder as a Candidatefor Evaluation in the Advanced Heat Shield Screening Program. AFML-TR-70-283, U.S. Air Force, Mar. 1971. (Available from DTIC as AD 884 111.) 9. O’Shea, J.: AutoweaveTM: A Unique Automated 3-D Weaving Technology. Proceedings of the Third Textile Structural Composites Symposium, Drexel Univ., 1988, pp. 21&224. 10. Hemck, John W.: Multidimensional Advanced Composites for Improved Impact Resistance. Materials Synergisms, Volume I O of National SAMPE Technical Conference, SOC. for the Advancement of Material and Process Engineering, 1978, pp. 38-50.

100

Textile Preforms f o r Carbon- Carbon Composites 11. Pastenbaugh, J.: Aerospatiale Technology. Proceedings of the Third Textile Structural Composites Symposium, Drexel Univ., 1988, pp. 189-209. 12. Bruno, P. S.; Keith, D. 0.; and Vicario, A. A., Jr.: Automatically Woven Three-Directional Composite Structures. SAMPE Q., vol. 17, no. 4, July 1986, pp. 10-17. 13. Geoghegan, P. J.: DuPont Ceramics for Structural Applications-The SEP Noveltex Technology. Proceedings of the Third Textile Structural Composites Symposium, Drexel Univ., 1988, pp. 225-239. 14. Fukuta, K.; and Aoki, E.: 3-D Fabrics for Structural Composites. Paper presented at the 15th Textile Research Symposium (Tokyo, Japan), 1986. 15. Fukuta, K.; Onooka, R.; Aoki, E.; and Magatsuka, Y.: Application of Latticed Structural Composite Materials With Three Dimensional Fabrics to Artificial Bones. Bull. Res. Inst. Polym. & Text., vol. I, no. 131, 1982, p. 151. 16. KO, F. K.: Three-Dimensional Fabrics for Structural Composites. Textile Structural Composites, T. W. Chou and F. K. KO, eds., Elsevier, 1989, pp. 129-169. 17. Williams, D. J.: New Knitting Methods Offer Continuous Structures. Adv. Compos. Eng., Summer 1978, pp. 12-13. 18. Hickman, G. T.; and Williams, D. J.: 3-D Knitted Preforms for Structural Reaction Injection Moulding (S.R.I.M.). How To Apply Advanced Composites Technology, ASM InternationalTM, 1988, pp. 367-370. 19. KO,Frank K.; and Kutz, John: Multiaxial Warp Knit for Advanced Composites. How To Apply Advanced Composites Technology, ASM InternationalTM, 1988, pp. 377-384. 20. KO, F. K.; Pastore, C. M.; Yang, J. M.; and Chou, T. W.: Structure and Properties of Multi-directional Warp Knit Fabric Reinforced Composites. Composites '86: Recent Advances in Japan and the United States, Proceedings-JapanU.S. CCM-III, K. Kawata, S. Umekawa, and A. Kobayashi, eds., Tokyo, 1986, pp. 21-28. 21. KO, Frank K.: Developments of High Damage Tolerant, Net Shape Composites Through Textile. Fifh International Conference on Composite MaterialsICCM-V, William C. Hanigan, Jr., James Strife, and Ashok K. Dhingra, eds., Metallurgical SOC.,Inc., 1985, pp. 1201-1210.

101

CC Materials and Composites 22. Brown, R. T.; Patterson, G. A.; and Carper, D. M.: Performance of 3-D Braided Composite Structures. Proceedings of the Third Textile Structural Composites Symposium, Drexel Univ., 1988, pp. 180-185. 23. Popper, Peter; and McConnell, Ronald: A New 3D Braid for Integrated Parts Manufacture and Improved Delamination Resistance-The 2-Step Process. Advanced Materials Technology '87-Volume 32 of International SAMPE Symposium and Exhibition, Ralph Carson, Martin Burg, Kendall J. Kjoller, and Frank J. Riel, eds., SOC. for the Advancement of Material and Process Engineering, 1987, pp. 92-103. 24. Popper, Peter; and McConnell, Ronald: U.S. Patent No. 4,719,837, Jan. 19, 1988. 25. McAllister, L. E.; and Lachman, W. L.: Handbook of Composites, Volume 4, Fabrication of Composites, A. Kelly and S. T. Mileiko, eds., North Holland (Amsterdam), 1983. 26. Tampolkki, U. M.; Zhigun, I. G.; and Polyakov, V. A.: Spatially Reinforced Composite Materials Handbook. Moscow Masinostroyeniye, 1987. 27. Adsit, N. R.; Carnahan, K. R.; and Green, J. E.: Mechanical Behavior of Three-Dimensional Composite Ablative Materials. Composite Materials: Testing and Design (Second Conference), ASTM Spec. Tech. Publ. 497, c. 1972, pp. 107-120. 28. Lachman, W. L.; Crawford, J.; and McAllister, L. E.: Multidirectional Reinforced CarbonKarbon Composites. Proceedings ZCCM-I1 , B. Norton, R. Signorelli, K. Street, and L. Phillips, eds., Metallurgical SOC.AIME, 1978, pp. 1302-1319. 29. Sepcarb Data Sheet, SOC. Europeenne de Propulsion, Saint Medard en Jalles, France, 1985. 30. McAllister, L. E.; and Tavema, A. R.: Development and Evaluation of Mod 3 Carbon/Carbon Composites. Materials Review for '72, Volume I7 of National SAMPE Symposium and Exhibition, SOC. of Aerospace Material and Process Engineers, 1972, pp. 111-A-THREE-1-111-A-THREE-7. 3 1. Lamico, P.: Recent Improvement in 4-D Carbon-Carbon Materials for Ablation Resistant Nozzle Parts. AIAA Paper No. 77-822, AIAAISAE Propulsion 13th Conference, 1977.

102

Textile Preforms for Carbon-Carbon Composites 32. Laramee, R. C.; and Canfield, Alan: CarbonKarbon Composites-Solid Rocket Nozzle Material Processing, Design, and Testing. Composite Materials: Testing and Design (Second Conference), ASTM Spec. Tech. Publ. 497, c. 1972, pp. 588-609. 33. Mullen, C. K.; and Roy, P. J.: Fabrication and Properties of AVCO 3-D CarbonKarbon Cylinder Material. Proceedings of the 17th National SAMPE Symposium, 1972, p. 111-A-2. 34. Pierson, H. 0.: Development and Properties of Pyrolytic Carbon Felt Composites. Advanced Techniques for Material Investigation and Fabrication, Volume 14 of Science of Advanced Materials and Process Engineering Proceedings, SOC.of Aerospace Material and Process Engineers, 1968, paper 11-4B-2. 35. Stoller, H. M.; and Frye, E. R.: Processing of Carbon-Carbon Composites. Advanced Materials: Composites and Carbon, American Ceramic SOC.,c. 1972, pp. 165-172. 36. Guo, L.: An Investigation of the 4D Texture in Carbon-Carbon Composites and Its Behaviors. International Symposium on Composite Materials StructuresAbstracts of Papers for Work-in-Progress, T. T. Loo and C. T. Sun, eds., (Beijing, China), 1986, pp. 9-12. 37. Herrick, J. W.: Multidimensional Advanced Composites for Improved Impact Resistance. 10th National SAMPE Technical Conference, 1978, p. 38. 38. Herrick, J. W.: Multi-Directional Advanced Composites for Improved Damage Tolerance. SME Tech. Paper EM84-104, Jan. €984. 39. Herrick, J. W.: Advanced Impact Resistant Multidimensional CompositesFinal Report. Rep. No. FR-79-1-3, Jan. 30, 1978.

40. Herrick, J. W.: Advanced Impact Resistant Multidimensional Composites. FR-79-1-3 (Contract N00019-77-C-0430), Fiber Materials, Inc., Jan. 1979. (Available from DTIC as AD A068 517.) 41. Herrick, J. W.; and Globus, Robert: Impact Resistant Multidimensional Composites. Materials 1980, Volume 12 of National SAMPE Technical Conference, SOC. for the Advancement of Material and Process Engineering, 1980, pp. 845-856.

42. Chou, T. W.; and Yang, J. M.: Structure-Performance Maps of Polymeric Metal and Carbon/Carbon Composites. Metall. Trans. A , vol. 17A, Sept. 1986, pp. 1547-1559.

103

CC Materials and Composites 43. Yang, Jenn-Ming; Ma, Chang-Long; and Chou, Tsu-Wei: Fiber Inclination Model of Three-Dimensional Textile Structural Composites. J. Compos. Mater., vol. 20, no. 5, Sept. 1986, pp. 472484. 44. Ma, C. L.; Yang, J. M.; and Chou, T. W.: Elastic Stiffness of ThreeDimensional Braided Textile Structural Composites. Composite Materials: Testing and Design, ASTM STP 983, American SOC.for Testing and Materials, 1986, pp. 404-421. 45. KO, F. K.; Pastore, C. M.; Lei, Charles; and Whyte, D. W.: A Fabric Geometry Model for 3-D Braid Reinforced FP/Al-Li Composites. Competitive Advancements in MetalslMetals Processing-Proceedings of 1987 International SAMPE Metals Conference, August 18-21, 1987.

Bibliography A. Chou, T. W.; and KO, F. K.: Textile Structural Composites. Elsevier, 1988. B. Hearle, J. W. S.; Grosberg, P.; and Backer, S.: Structural Mechanics of Fibers, Yarns nd Fabrics. John Wiley & Sons, 1969.

C. Kaswel, E. R.: Wellington-Sears Handbook of Industrial Textiles. West Point Pepperell, Inc., 1963. D. Klein, A. J.: Which Way to Weave. Adv. Mater. & Process., vol. 2, no. 3, 1986, pp. 40-43.

104

Chapter 5 Carbon-Carbon Matrix Materials N. Murdie, C. P. Ju, J. Don, and M. A. Wright Southern Illinois University at Carbondale Carbondale, Illinois

Introduction Carbon Fibers

106 108

Fabrication Methods of CC Composites Liquid Phase Infiltration Pitch Matrices

111

112

112

Thermoset Resin Matrices

115

Gas Phase Infiltration Process

118

Isothermal Chemical Vapor Deposition

119

Thermal-Gradient Chemical Vapor Deposition

119

Differential Pressure Chemical Vapor Deposition Matrix Inhibition

120

121

Microstructural Characterization Techniques Optical Microscopy X-Ray Diffraction

123

123 124

Scanning Electron Microscopy

125

Transmission Electron Microscopy Microstructure of CC Matrices

127

Pitch Matrix Composites

127

Resin Matrix Composites

135

CVI Matrix Composites

126

142

Influence of Matrix on Composite Properties

149

105

CC Materials and Composites General Background Elastic Modulus

150

Tensile Strength

151

149

Matrix Dominated Properties

154

Two-Dimensional Reinforcements Three-Dimensional Reinforcements Acknowledgments References

156 156

157

158

Introduction Carbonaceous materials have been used as refractories, as electrodes in both steelmaking and aluminum production, as moderators for nuclear reactors, rocket nozzles and exhausts, as aircraft brakes, and in various chemical and electrical applications (ref. 1). Initially, the carbons used for each structure were comprised of a granular material bonded together with a carbonized resin or pitch matrix. Improvements in the mechanical properties of structural carbons were obtained by depositing carbon layers from the gas phase. Mechanical properties of this material were superior to those exhibited by the pitch- or resin-bonded material, especially so when measured parallel to the basal plane of the graphite (ref. 1). Carbon fibers are a relatively recent development in which the basal planes of the graphite-like crystallites are arranged almost parallel to the fiber axis. This arrangement produces a fiber that exhibits a significant anisotropy of properties. The elastic modulus, strength, and electrical conductivity are large parallel to the fiber axis, while the same properties measured transverse to that direction are an order of magnitude smaller. Conversely, the thermal expansion coefficient is small parallel to the fiber axis, but much larger perpendicular to it (ref. 2). Carbon-carbon (CC) composites are fabricated by infiltrating a carbon precursor into a preform of carbon fibers. Chemical vapor infiltration (CVI) directly produces a carbon fiber-reinforced composite at approximately 1000°C. Alternatively, similar materials can be formed by heating pitch or resin matrix composites to a temperature of about 1000°C. This process eliminates volatile elements from the matrix and “carbonizes” it. The appreciable matrix shrinkage that occurs during carbonization generates pores and cracks. Differences between the coefficient of thermal expansion of the fibers and matrix also generate internal stresses and stress cracks on cooling (ref. 3). The carbonized solid therefore exhibits a low density that must be increased to the desired level either by reinfiltration with and carbonization of additional matrix precursor or by CVI with carbon. This process is

106

Carbon- Carbon Matrix Materials repeated until an acceptable composite density is obtained (ref. 4). Multiple cycles of CVI are also necessary in order to densify the solid. Surface machining has to be included in the process since, in most cases, a thick surface layer is formed which tends to prevent diffusion of the hydrocarbon gas into the internal regions of the preform.

If the composite is to be used at very high temperatures, the material may be graphitized by heating above 20OO0C. Both the fibers and the matrix materials then exhibit increasing graphite-like properties. The graphitic form of carbon has a hexagonal crystal structure in which each carbon atom is associated with four valence electrons, three of which form tight covalent bonds with neighboring atoms; the fourth is more loosely bound. Since each carbon atom is surrounded by three neighbors at equal distances in a single plane, a hexagonal ring structure results. The layers formed by these rings are primarily bonded together by van der Waals forces that are much weaker than the covalent bonds. Adjacent atoms in any layer are closer (1.42 A) than the spacing (3.35 A) between layers. This atomic configuration results in extreme anisotropy in the crystal structure. Graphite is unique among engineering materials in that parallel to the basal planes, its specific elastic modulus (modulus/weight ratio) is the largest known, while much lower values of the same property are recorded normal to these planes (ref. 2). Carbon fibers, when arranged unidirectionally in any carbon matrix, produce a solid that exhibits anisotropic mechanical properties. In this case, since the fibers take most of the load, both the stiffness and strength of the composite are large when measured parallel to the fibers but are small when measured perpendicular to them. In addition, the composite becomes tough because a pseudo-plasticity is exhibited if debonding of the fibers and matrix occurs during failure. The frictional resistance of pulling the fibers out of the matrix behind the crack front then contributes to the work of fracture. Fibers can be aligned unidirectionally, multidirectionally, or they can be present as various fiber weaves; these include braided yarns, stacked two-dimensional (2-D) fabrics, orthogonal fabrics, or multidirectional weaves. An important variable that influences the properties of CC composites is the volume fraction of fibers relative to the carbon matrix. Generally, the higher the volume fraction of strong/stiff fibers in a matrix, the greater the strength and stiffness of the composite. The type of weave can vary, with those more commonly used consisting of either a plain weave or a harness satin weave (ref. 5). The harness satin weaves allow a higher volume fraction of fibers and usually produce higher strengths due to the floating yarns (refs. 4 and 5). Carbon-carbon composites are generally considered among the most competitive materials in structures designed for use at high temperature. Unfortunately, although their stability is extremely good in nonaggressive mediums, their performance quickly degrades if oxidation occurs (refs. 6 and 7).

107

CC Materials and Composites A number of techniques designed to provide some level of oxidation protection have been examined. For instance, the composites (or matrix) can be purified to remove catalytic oxidants (refs. 8 and 9), or they can be treated to passivate reactive sites (refs. 10 through 13). Such inhibited composites appear useful, especially for low-temperature applications. For high-temperature applications, elements or compounds can be added to serve as oxygen getters, providing constituents for protective glassy films, or providing materials that might diffuse and block oxygen transport through cracks in surface films (refs. 14 through 17). Nevertheless, despite all the research activity directed to inhibit oxidation of the matrix, development of an external coating system capable of protecting the total composite (ref. 18) is considered essential. This coating system is particularly important if CC materials are to be used in long-term high-temperature oxidative environments.

Carbon Fibers Efficient, high-performance materials such as those used in aeronautics, astronautics, and other segments of the transportation industry must be strong and stiff and must also exhibit a minimum weight. It is well known that the potential ultimate tensile strength of a material is proportional to the Young’s modulus; thus, the specific modulus (Young’s modulus/density) is the best single measure of the potential mechanical performance of a material. This property is considered useful in the assessment of fibers because it represents a first consideration in selecting a carbon fiber with respect to its electrical conductivity, its coefficient of thermal expansion, and its strength and stiffness. A summary of the major carbon fiber varieties available now is given in tables 1 and 2* (ref. 19). Table 1. Comparison of Properties of Pitch and Pan Fibers

*Based on Arnoco performance products technical data sheets.

108

Density Fiber types Graphite, PAN, low-cost high-strength (LHS): Celion A-S T-300 Type 111 3T Intermediate-modulus (IM): T-400 Type I1 HTS 4T High-modulus (HM): T-50 Type 1 5T 6T Ultrahigh-modulus (UHM), GY-70 Pitch, type P Silicon carbide, 5.6-mil, Carbon core Tungsten core

Ultimate tensile strength

Tensile modulus

lb/in3

kg/m3

ksi

MPa

Msi

GPa

0.0632 0.0654 0.0636 0.0643 0.0650

1749 1810 1760 1780 1799

419 420 400 350 300

2889 2896 2758 2413 2068

34 34 33 33 30

234 234 228 228 207

0.0643 0.0618 0.0658 0.0650 0.0683 0.0603 0.0672 0.0668 0.0686

1780 1711 1821 1799 1891 1669 1860 1849 1899

425 360 410 350 350 300 350 400 420

2930 2482 2827 2413 2413 2068 2413 2758 2896

33 40 36 38 53 57 56 48 58

228 276 248 262 365 393 386 331 400

0.0672 0.0730

1860 2021

250 250

1724 1724

75 50

517 345

0.1109 0.1192

3070 3299

550 550

3792 3792

50 60

414 414

GC Materials and Composites Continuous carbon fibers are fabricated using a process that essentially consists of three steps: spinning, stabilization (oxidation), and carbonization (heating to 1000°C to 1600°C in an inert atmosphere, usually nitrogen). Carbonization involves the pyrolysis of the stabilized fibers to increase the carbon content. Thus, this step involves the elimination of heteroatoms and other molecules such as H2, HCN, NH3, 0, CO, N2, H, CH3, and S. A fourth step, graphitization (heating to about 2500°C in an inert atmosphere) is sometimes performed if very high modulus, high thermal conductivity, or low thermal expansion fibers are required. In the early 1960’s, rayon was a popular carbon fiber precursor, and it is presently used in the nose cap and exhaust nozzle of various aerospace vehicles, e.g., the Space Shuttle (ref. 18). However, large quantities of carbon fibers are no longer routinely produced using this material because it must be stretched at a high temperature in order to produce fibers with optimum properties. This requirement, when combined with the low carbon yield of about 35 percent, results in a process that is not commercially viable (ref. 20). Rayon can be replaced by a precursor based on PAN. High-quality carbon fibers are cheaper to produce from PAN because stretching only needs to be carried out at the low temperatures typically used during stabilization, i.e., between 200°C to 280°C. In addition a carbon yield of about 65 percent is exhibited. Fibers produced from pitch should be the cheapest carbon fibers since pitch fibers require no stretching and can give a carbon yield of about 85 percent. Pitch fibers can be processed from an isotropic pitch or from a material that has been extensively treated to produce mesophase. Although isotropic pitch-based fibers are inexpensive, they are restricted for thermal insulation because of their mechanical properties. More competitive properties are produced from the more expensive mesophase feedstocks. Even in this case, however, the superior modulus values of the resulting carbon fibers are not matched by correspondingly high strengths (refs. 23 through 26). Very high modulus fibers (>75 Msi (517 GPa)) can be more easily produced from pitch than from PAN-based fibers. Thus, they tend to be used where a high modulus or the associated properties of high thermal and electrical conductivity or negative thermal expansion coefficient is desired. However, except for this specialized use, pitch fibers tend to be uncompetitive; PAN fibers are, therefore, the carbon fiber most used as a structural reinforcement (ref. 20). High modulus (HM) fibers are fabricated by heating to very high temperatures. These fibers cost more than low modulus fibers, exhibit smaller longitudinal coefficients of thermal expansion, have high thermal conductivities, and normally have lower surface reactivities. The fiber used in most engineering structures exhibits a low modulus (LM) of about 30 Msi (206 GPa).

110

Carbon- Carbon Matrix Materials If low modulus fibers are heated to a graphitization temperature of approximately 25OO0C, primarily to increase the modulus, the degree of axial alignment and the size of the crystallites also increase. Conversely, the interplanar spacing decreases and the fiber shrinks. In order to avoid the dimensional changes, fibers used to produce CC are often, but not always, graphitized, by heating to about 25OO0C, before being infiltrated with a matrix precursor. Fiber types are obtained in a tow form and can contain a small number of fibers (1000) or a relatively large number of fibers (12000). Normally, the larger tows are cheaper to produce.

Fabrication Methods of CC Composites The classical method for fabricating carbon materials involves combining solid particles of pure carbon (filler coke) with a precursor that can be carbonized to serve as a binder. In CC composites, carbon fibers are used as the primary carbon instead of the filler coke. Unfortunately, the mass loss and shrinkage of the matrix during carbonization result in a final material that exhibits considerable bulk porosity. The objective of repeated infiltration-pyrolysis is to densify the initial porous skeleton. Densification is achieved by impregnation with liquid or gaseous carbon precursor compounds and subsequent pyrolysis. A key factor in the selection of a matrix carbon precursor involves the ability to fully densify the preform and to achieve a high char yield. Three basic methods of fabricating carbon fiber reinforced carbon materials exist. The first two methods are based on thermally degrading a thermosetting resin or a thermoplastic pitch. The third method involves depositing carbon into a fibrous preform using CVI. As discussed in a following section, the choice of fabrication method depends to a major extent on the geometry of the part being processed. Thin sections are prime candidates for CVI processing; however, since this method tends to preferentially deposit in and on the surface layers, it is not suitable for the fabrication of thick sections. Thick sections therefore tend to be produced using resin or pitch infiltration. Thermosetting resins remain solid during carbonization; however, pitches soften and tend to flow from a preform at high temperatures; therefore, they require containment during the carbonization step. Complex shapes are difficult to fabricate using either CVI or pitch matrix materials owing to the difficulty of maintaining the dry fiber preform shape during the initial infiltration. Hybrid densifications are sometimes practiced when a rigid structure is first made using the resin prepreg, autoclave molding process. In this process after carbonization, subsequent infiltrations are made with CVI or pitch.

111

CC Materials and Composites A recent development allows the impregnation of thin-section carbonaceous preforms with 100 percent mesophase (ref. 27). Carbonization can then be achieved after an oxidation step has rendered the mesophase infusible.

Liquid Phase Infiltration Two types of liquid impregnants are used to densify CC performs. The first includes pitch, which may be coal tar or petroleum based, and the other is derived from resins or polymers. Both are used because they have suitable viscosities and the carbon yields are high enough to provide high-density CC composites up to -1.9 g/cm 3 .

Pitch Matrices The basic fabrication method for producing pitch matrix CC composites is to use pressure to force pitch into an evacuated cavity that contains the dry fiber preform. The chemical composition of the pitch is believed to control the microstructure and reactivity of the resulting CC solid (ref. 28). It is known, for example, that while most pitches produce graphitic matrices, other pitches cannot be graphitized because some of the molecules are not capable of forming a hexagonal-type three-dimensional graphitic network. In other pitches, side chains exist, for example, alkyl side chains, which distort the crystal structure and render the resulting carbonized material somewhat reactive (ref. 28). Conversely, many pitches contain trace metal impurities that catalyze graphitization locally (ref. 29), while others are known to catalyze high-temperature oxidation reactions (refs. 30 through 32). The infiltration-carbonization cycle can be performed at atmospheric pressure or at pressures up to 30000 pounds per square inch. At atmospheric pressure, the carbonization of the pitch matrix is carried out by heating to 1000°C under a very slight partial pressure of nitrogen. This carbonization process can be characterized by the viscosity changes experienced by the pitch. Depending upon the elemental composition and the thermal history, the matrix softens until oxygenated functional groups and their sulfur and nitrogen equivalents are released, subsequently resulting in minimum viscosity (ref. 33). Typical viscositytemperature relationships for pitches are shown in figure 1 (ref. 34). Essentially, heating isotropic pitch from room temperature results in the melting of solid pitch and in low viscosity. Continued increase of temperature causes relatively insignificant changes in viscosity although devolatilization of the pitch continues until, at a critical temperature (dependent on the elemental composition and thermal history), the viscosity of the pitch increases rapidly (ref. 35). The kinetics of mesophase formation depend on the composition, reaction temperature, degree of agitation, and the removal rate of the lower molecular

112

Carbon- Carbon Matrix Materials

Figure 1. Variation of apparent viscosity with temperature of various pitch fractions. Reprinted with permission, Materials Technology Distinguished lecture 1986, “The Future of Carbon-Carbon Composites” (Professor Erich Fitzer). weight fractions (sparging) as shown in figure 2. In general, however, as the temperature of the pitch is increased, the rate of mesophase formation and viscosity increase. Mesophase spheres similar to those shown in figure 3 begin to appear at temperatures greater than 4OO0C, growing until they coalesce to become a continuous phase. At this stage, a preferred orientation of large numbers of these liquid crystals imparts a directionality or anisotropy to the properties of the resultant coke (ref. 28). Maintaining the temperature for longer times results in the viscosity increasing rapidly until the pitch becomes a brittle, predominantly crystalline solid (coke). Pitch-based matrices, when graphitized, are dense (- 1.9 g/cm3). Carbonization of isotropic coal tar or petroleum pitch produces a 50- to 60percent coke yield under atmospheric conditions. However, if the carbonization is performed very slowly or under high pressure (29000 psi (200 MPa)), the coke yield can be increased to 70 to 80 percent (ref. 5). If the 100-percent mesophase pitch, first produced from an initially isotropic pitch, is stabilized (oxidized) before carbonization, an even greater yield, up to 92 percent, can be achieved (ref. 27).

113

100

x

Q

e

.-c c g

-

::$/I -

m Ea) 40 2

f /

-

a)

a 20 -

0

/@

/@

7

I@

- - n + i ~ l l l , ~ r ; t I ~ l l ~ l I I l ~ I ( I I ~ l I I I I ~ ~ ~ ~ l l ~

Carbon- Carbon Matrix Materials

Figure 3. Photomicrograph of mesophase spheres in isotropic pitch. a CC composite was obtained with a density of -1.9 g/cm3. These data compare with the data for a density of 1.6 g/cm3 following eight impregnations at atmospheric pressure. Under atmospheric conditions, the maximum achievable density is -1.8 g/cm3.

Thermoset Resin Matrices Phenolics and epoxy resins are two types of commonly used thermosetting resins. Both these resins are cured prior to carbonization; since they are thermosets, they will not flow from the fibrous preforms during first carbonization. Resin matrix composites are fabricated from preimpregnated carbon fiber layers (prepreg) of woven fiber cloth. This type of material is preferred when fabricating complex shapes. The materials technology involves the impregnation of one layer of carbon fibers (or carbon fabric) with a resin. This prepreg is partially cured (B-staged) to a fixed degree of tackiness and can be used immediately or refrigerated for 6 to 12 months. Prepregs are cut and combined according to the need and hot-pressed or cured in an autoclave to produce a rigid solid. They are then pyrolyzed and subsequently densified by repeated infiltration/carbonization cycles. This process is described in detail by Curry et al. (ref. 37). Although

115

Matrix

Virgin matrix experimental char Carbon content, percent

Yield, percent

Resins Polybenzimidazole Polyphenylene Biohenol formaldehyde Furfural alcohol Phenol formaldehyde (phenolic) Epoxy novalac Poly imide

96 92 80 75 78 74 77

73 71 65 63 60 55 49

Pitches Coal tar Electrode binder Vegetable Petroleum

75 92 69 88

60 40 30 21

Synthetic pitches Truxene Isotruxene

95 95

87 70

Resin: Pitch blends Phenolic - 60%, Coal tar pitch

78

75

-

67

-

60

Furfural alcohol - 60%, Coal tar pitch Epoxy novalac - 60%, Coal tar pitch Heated 12'F/min to 1800'F in an inert gas

Autoclave pressure- 204 external to part - 177

- 149 9 -

-121

g

3 -93

-66

21

L

g E

$

CC Materials and Composites examples of frequently used infiltrants. In order to fill the smallest pores, impregnation can be carried out using a low pressure (50 psi; 3445 Pa). Since the density of resin char is relatively low, resin densification requires more cycles than those necessary using other precursors. Up to five cycles of resin impregnation are required to obtain a composite with a density of about 1.5 g/cm3 to 1.8 g/cm3 (ref. 4).

Gas Phase Infiltration Process The CVI process for carbon deposition uses volatile hydrocarbon compounds such as methane, propane, or benzene as precursor gases. Thermal decomposition of any of these gases is achieved on the hot surfaces of the substrate, resulting in a deposit of pyrolytic carbon and the emission of volatile by-products, which consist mainly of hydrogen. One of the problems associated with CVI is that under isothermal conditions, pyrolytic carbon deposits preferentially on the surface of the substrate; such deposits change open porosity to closed porosity and make the filling of internal pores difficult. Surface deposits always tend to build up rapidly if diffusion of the reacting gases to the deposition surface is allowed to control the overall reaction rate. This effect can be reduced, however, if chemical reaction on the external surface and inner surfaces is the rate controlling step. Kinetic studies (ref. 39) indicate that low temperatures tend to limit reaction while having relatively little effect on diffusional transport. Reaction rate control of the deposition process and hence more complete densification is therefore favored by low temperatures. Factors that influence the structure, uniformity, and rate of deposition of a CVI matrix include the nature of the substrate, the carrier gas temperature, the composition, the pressure, and the geometry (particularly the thickness) of the final structure (refs. 39 through 42). Typical effects, summarized in figure 5 (ref. 43) illustrate how the microstructure of the CVI carbon deposit can be made to vary from columnar through laminar to isotropic by altering the gaseous composition and deposition temperature. Essentially, high-temperature, low-propane concentrations favor isotropic deposits, and low-temperature, high-propane concentrations favor a columnar structure. In separate work, Jachlewski and Diefendorf (ref. 41) confirmed the findings of these authors, but also showed that isotropic deposits can be formed at temperatures less than those indicated (
118

Carbon-Carbon Matrix Materia

a,

c m

50

-

20

-

1100

1150

1200

1250

1300

1350

1400

Temperature, "C

Figure 5 . Effect of deposition condition on microstructure of pyrolytic carbon matrix deposition from propane (reJ43). Reprinted with permission. three methods of forming CVI carbon: isothermal, thermal gradient, and pressure gradient.

Isothermal Chemical Vapor Deposition In this process, a woven structure is placed within a furnace susceptor and is heated uniformly. The pressure and the temperature of the furnace are kept constant at typical values of 1 psi (6 KPa) and llOO°C, respectively. The flow rate of hydrocarbon gas is predetermined depending on the surface area of the substrate. Intermittent machining of the surface is required because the chemical vapor deposition (CVD) technique leaves a crust on the outer surface of the substrate. The machining cycle is repeated until the desired density is achieved (ref. 4). Thermal-Gradient Chemical Vapor Deposition In this technique, a carbon preform is supported on a mandrel. Inductive coils heat the surface of the mandrel to a temperature of about llOO°C. The hottest portion of the substrate is in contact with the mandrel, while the other side,

119

CC Muterials and Composites

-

1

.o c cn

w 3 -

s

/

2-

.o c 1-

/

/ '

1300 1400 Deposition temperature, "C

Figure 6. Variation of elastic modulus with deposition conditions from propane (ref 43). Reprinted with permission. in contact with the reacting gas, is cooler; thus, a thermal gradient through the substrate thickness is created. As the hydrocarbon gas passes through the furnace at atmospheric pressure, carbon is deposited on the hottest region of the woven structure. This hot section migrates through the thickness of the structure as the densified region grows toward the colder surface. This technique prevents the formation of a crust on the outer surface of the preform; thus, the machining step is eliminated. Unfortunately, the process tends to be restricted to large individual parts.

Differential Pressure Chemical Vapor Deposition Differential pressure CVD is a variation of the isothermal technique in which the inner portion of the fiber preform is sealed off from the furnace chamber at the base. Hydrocarbon gases are fed into the inner cavity at a positive pressure with respect to the furnace chamber. A pressure difference that forces the hydrocarbon to flow through the pores depositing carbon and exiting as hydrogen (refs. 4 and 39) is created across the wall of the structure. This technique also prevents the formation of an outer crust on the surface of the preform and facilitates densification uniformity.

120

Carbon- Carbon Matrix Materials

1600

1800

2000

Temperature, "C Figure 7. Relationship of density and structure with temperature and pressure (ref. 39). Reprinted by permission of the Society for the Advancement of Material and Process Engineering. Kotlensky, W. V.: A Review of CVD Carbon Infiltration of Porous Substrates. SAMPE J . 16, 1 9 7 1 , ~257. .

Matrix Inhibition The oxidation of carbon becomes quite rapid at temperatures above approximately 5OO0C (ref. 44). The exact rate of oxidation has been shown to depend on the active surface area, the porosity, the degree of crystallinity, the purity, the internal stress, and the absolute temperature (refs. 45 and 46). It has been shown (refs. 6, 7, and 47) that levels of 2 to 5 percent weight loss during oxidation leads to a substantial degradation (40 to 50 percent) of mechanical properties. Although short-term high-temperature applications @e., rocket nozzles and exhausts) may only require limited oxidation protection, long-term exposure normally requires that a surface barrier coating be present (ref. 48). Surface coatings are generally composed of refractory layers designed to separate the oxidizing gases from the composite. Unfortunately, most refractory coatings exhibit coefficients of thermal expansion that differ significantly from the CC substrate. Temperature changes produce stresses in the coating that are sufficient to cause cracking; thus, a glassy layer designed to flow and seal such

121

C C Materials and Composites cracks (refs. 18,44, and 48) is also incorporated. Oxidation protection using barrier layer techniques is the subject of a different chapter in this publication and will not be considered here. However, it is recognized that oxidation of the basic carbon structure should be prevented. Appropriate treatments are therefore applied to the matrix (and fiber) to achieve a minimum oxidation rate. Unfortunately, the details of most of these treatments are proprietary. The density of a composite is particularly important because gaseous penetration of any coating can produce gasification at both internal and external surfaces. A highly porous structure will gasify at a high rate, and it is for this reason that the rate of oxidation of any carbon tends to increase initially because of the expansion of internal surface area. Eventually, however, a decrease in reaction rate tends to occur as the less ordered structure (less graphitic) is gasified. Donnet, in a series of articles (refs. 49 and 50), has emphasized the importance of crystalline order; ". . .the oxidation rate (of carbon) is strongly dependent on the structure of the carbon layers, less organized parts appearing more reactive than better organized parts.. .." In these experiments, it was stated that less graphitic structures always oxidize first, leaving the more graphitized material. Gasification is extremely sensitive to purity; many metallic impurities, even in very small amounts, are known to be very aggressive oxidation catalysts (refs. 31 and 51). Elements such as iron, calcium, lead, copper, vanadium, chromium, manganese, nickel, and cobalt have been shown to increase the rate of gasification of carbons. Baker (ref. 52) has described a series of electron microscope experiments which emphasize the catalytic effect of various metals (and some oxides) on the gasification of pure graphite. Attempts have been made to improve the oxidation resistance by removing impurities. Acid washing, for instance, has been shown to be beneficial (ref. 8), as has purification of the original feedstock and purification of the carbon by high-temperature (30OO0C) halogen treatments (ref. 53). These high-temperature treatments have the additional effect of graphitizing suitable matrices, as well as reducing porosity (ref. 54). As stated previously, a maximum oxidation resistance is obtained by coating the outer surfaces of a structure with a layer of material specifically designed to stop oxygen from contacting the carbon substrate. A viscous glass, frequently silicon based, is usually incorporated into this coating. This material is designed to flow at high temperatures and fill any cracks that might develop. An alternative approach is particularly useful at lower temperatures or for conditions in which oxygen has diffused through the outer coating or along cracks. Low-temperature protection can be achieved by adding a boron containing compound to the matrix precursor which on oxidation forms a viscous borate glass that covers the internal surfaces (ref. 18). This type of protection is similar but involves a material less viscous than the silicate glass formed at higher temperatures (refs. 18, 44, and 48). Other inhibition treatments involve the addition of materials to the matrix that can also produce glass sealants. A JTA quality nuclear graphite contains

122

Carbon- Carbon Matrix Materials zirconium diboride and elemental silicon specifically to produce a glassy phase at about 12OO0C (ref. 55). Carbon-carbon composites cannot usually contain the high amounts of glass-forming materials present in JTA graphite (approximately 40 percent); however, smaller quantities of boron and silicon compounds have been added to the composite in order to provide a source for the continuous generation of new glassy materials (refs. 18 and 56). One of the problems associated with oxidation-resistant additions to organic precursors is that the addition will either be removed or modified during the process of carbonization/graphitization. Boron added as fine particulates will change to boron carbide during high-temperature treatment, whereas during polymerization of the original resin, it will oxidize and could hydrate (ref. 57). Most of these impurity or inhibition reactions may adversely affect the mechanical properties of the composite. An alternative method to improve oxidation resistance at lower temperatures involves protecting active sites of the graphite crystal structure with reactive elements such as halogens and phosphorous compounds (refs. 10 and 11). Active sites most prone to oxygen attack are edge sites associated with the planar graphitic lattice, dislocation sites, and vacancies. Reaction of these sites with elements like the halogen gases suggests the production of stable complexes that prevent oxidation.

Microstructural Characterization Techniques Optical Microscopy Optical microscopy is used to relate the microstructure of carbon materials to their processing conditions (refs. 58 through 63). Hot-stage microscopy enables the observation of mesophase formation (ref. 64) and changes that occur in the crack morphology of CC as a result of increasing temperature (ref. 3). Image analysis systems have enabled quantification of the microstructure (refs. 64 and 65). Microstructural characterization techniques that have become increasingly popular include the use of polarized light (ref. 66) and differential interference contrast (ref. 67). The popularity of these techniques has largely resulted from the anisotropic properties of certain carbons, a property which leads to the generation of structurally related colors. For example, interference colors are generated when using polarized light and a half-wave retarder plate. These data can be used to assess the orientation of the constituent lamellar planes of anisotropic carbon as they terminate at a polished surface (refs. 66 through 68). The size and shape of these isochromatic regions determine their optical texture (ref. 69). Optical textures vary from isotropic (nongraphitizable) to domain anisotropic (highly graphitizable). The definitions of such textures are shown in table 4 (ref. 28). The use of optical microscopy in characterizing carbonaceous materials is limited to a resolution of about 0.5 pm. Any carbon layer less than 0.5 pm thick will appear as part of

123

CC Materials and Composites the surrounding material. For example, the highly graphitic sheath (0.1 pm wide) observed at the surface of some PAN fibers (Type I) using the transmission electron microscrope cannot be resolved using optical microscopy (ref. 20). Table 4. Nomenclature to Describe Optical Texture in Polished Surfaces of Coke (ref. 28*) Isotropic (I) Very fine-grained mosaics (VMF) Fine-grained mosaics Medium-grained mosaics (Mm) Coarse-grained mosaics (Cm) Supra mosaics (Sm)

No optical activity < O S pm diameter

<1.5 pm; > O S pm diameter <5.0 pm; >1.5 pm diameter <10.0 pm; >5.0 pm diameter Mosaics of anisotropic carbon oriented in the same direction to give a mosaic area of isochromatic color.

Medium flow anisotropy Elongated (MFA) Coarse flow anisotropy Elongated (CF) Acicular flow domain anisotropy (AFD) Flow domain anisotropy Elongated (FD) Small domains Isochromatic (SD) Domain-isometric (D)

<30 pm length; <5 pm width <60; >30 pm length; 5 pm width >60 pm length; <5 pm width >60 pm length; >10 pm width <60; >10 pm diameter

>60 pm diameter

Db is from basic anisotro y of low volatile coking vitrains anianthracite. Dm is by growth of mesophase from final phase.

*Reprinted with permission. Marsh, H.; and Menendez, R.: Mechanisms of Formation of Isotropic and Anisotropic Carbons. Introduction to Carbon Science, H. Marsh, ed., Butterworths (London), 1989, p. 37.

X-Ray Diffraction X-ray diffraction (XRD) provides useful information on the crystallinity of carbons and other materials. Since the beam size of the X-ray source is large, 1 mm, the results are an average obtained from a large volume of a specimen that may include many different constituents.

124

Carbon- Carbon Matrix Materials The average interplanar spacing d of the analyzed volume can be calculated from the Bragg equation

X = 2d sin 0 where 0 is the Bragg angle and d is the interplanar spacing. The X-ray diffraction patterns reflect the structural order of the crystal. For carbonaceous materials, the widths of the (002) lines are generally used to calculate the crystallite height (L,) using the Scherrer equation (ref. 69),

L c = [KX]/[Boo2(20)cos 01 where K is the Scherrer constant ( K = 0.9), X is the X-ray wavelength, 0 is the Bragg angle, and Bo02 is the half-height width of the 002 diffraction line. The crystallite diameter (L,) is usually determined from the half-height width, B ~ Q , of the 10 or 11 patterns for turbostratic materials using the Warren and Bodenstein formula (ref. 70) L u = [KXI/[B,p (20)cos 01 where 0 is the Bragg angle, K = 1.77. Crystallites in turbostratic PAN fibers typically exhibit a spacing between 3.42A and 3.45 A as shown by the width of the (002) diffraction line (ref. 71). Generally, heating graphitizable carbons to temperatures approaching 30OO0C increases the degree of graphitization, which can be detected by a decrease in (002) spacing to that of crystal graphite, 3.35 A. Polycrystalline graphite produces a typical Debye ring pattern in which the width of the appropriate lines can be used to estimate the average crystallite size both in the hexagonal c direction (L,) and in the transverse a direction (&).

Scanning Electron Microscopy Scanning electron microscopy (SEM) is used to examine the microstructure and surface topography of carbon specimens at magnifications from -20 to 200 000 times and resolutions more than 100 A (refs. 72 through 75). The main advantage of this technique is the large depth of focus, which is useful for the examination of fracture surfaces. Scanning electron microscopy is used to reveal microstructural information of fractured surfaces of carbon composites or of surfaces that have been etched. Etching is often used to distinguish the different carbon components in a composite. However, over-etching should be avoided because undesired artifacts such as pitting and microcracks may be produced.

125

CC Materials and Composites Transmission Electron Microscopy Transmission electron microscopy (TEM) is a useful microstructural tool, because it enables direct observation of the phase details, interfaces, and degree of crystallinity within carbonaceous materials (refs. 76 through 82). To use this technique, it is necessary to produce samples that are thin enough (<2000 A) for the electron beam to pass through. Grinding and microtome sectioning are the two thinning techniques that have been used almost exclusively in the TEM studies of fibers, mesophase, and other carbonaceous materials. More recently, thin foil sections of CC composites have also been successfully produced using the mechanical dimpling and atom milling techniques (ref. 77). Three TEM imaging techniques are used in microstructural analysis: bright field, dark field (DF), and lattice image. The bright field is a mode in which the transmitted beam is used to produce an image of the specimen. The DF mode is obtained by tilting the illumination system so that the desired diffracted beam coincides with the optical axis of the microscope. An objective aperture is then inserted to exclude all other beams. Dark-field imaging is used in carbon work because all crystallites of similar orientation can be identified. The diffracting crystallites are generally seen as white regions on a dark background. An illustration of this effect is shown in figure 8, taken from a pitch-based carbon fiber using the (002) reflection (ref. 80). In this figure, individual fibrils are shown highly oriented parallel to the fiber axis. Each fibril exhibits a length of a few micrometers and a thickness of 300-400 A. Each bright region (fibril) consists of a number of crystallites with similar basal plane orientation. Lattice imaging is another mode of operation that is well suited for direct observation of the lattice planes in carbonaceous materials. The image obtained in the image plane must be slightly underfocused in order to obtain the lattice image (ref. 83). The lattice image in figure 9 shows the typical (002) basal plane orientation of fiber and matrix regions found in transverse sections. This lattice image was formed using a portion of the (002) ring under a tilted-beam condition. The turbostratic structure of the PAN fiber was observed across the whole fiber thickness, even immediately adjacent to the interface, as seen in figure 9. Compared with the surrounding pitch matrix, the coherent domains in the PAN fiber are much smaller in size and more uniform in structure. The selected area diffraction (SAD) technique also allows for the determination of crystallite dimensions Lc, La, and do02 in a similar way to that described for X-ray diffraction. The major advantage of SAD over XRD in such measurements is that SAD can be used to obtain crystallite information from very small regions (<1 pm). The degree of perfection of basal plane alignment 2 with respect to the

126

Carbon-Carbon Matrix Materials

Figure 8. Dark-field micrograph of fibrils within pitch-based carbon fiber. fiber axis can be obtained from the half-width measurement of the (002) diffraction arc (ref. 2 ) . The SAD results of a PAN fiber in a 3-D CC composite indicate that the structure of the PAN fibers is turbostratic. Both SAD -andDF show that the degree of perfection for basal plane alignment to the fiber axis increases gradually from fiber core (2 = 38") to the surface (2 = 22') (ref. 73).

Microstructure of CC Matrices Pitch Matrix Composites The fabrication of CC composites using pitch matrix materials involves first infiltrating the fiber preform with pitch before heating to carbonization temperature. As discussed previously, during carbonization various gases are evolved, and providing the rate of heating is slow, the isotropic pitch is transformed gradually to mesophase. The molecular structure of mesophase has been described by Brooks and Taylor (ref. 84) to consist of discotic liquid crystals aligned parallel to one another. Although these layers exhibit high in-plane strength and stiffness, they can easily

127

CC Mate-als and Composites

Figure 9. Photomicrograph of lattice fringe image of interface between PAN fiber in pitch matrix. shear relative to each other. Because of this, such layers will tend to preferentially orient themselves parallel to any surface across which they move. White (ref. 27) has also pointed out that wetting of a fiber surface by the mesophase will result in a similar effect. Regardless of the absolute mechanism of orientation, the result is that in a carbon matrix derived from pitch, each fiber will be surrounded by a matrix of highly aligned, strong-stiff graphitic basal planes. Figure 10 is an optical micrograph of a 3-D CC composite composed of PAN (T-300) fibers and a coal tar pitch matrix. The composite was processed using a multiple high pressure impregnation carbonization (HPIC) technique at 30 ksi (206 MPa) pressure to densify the composite. Figure 10 position A shows a transversely oriented fiber bundle surrounded by the interbundle matrix (position C). The composite contains a significant amount of porosity (positions D and E) of which there are two types. The larger voids (position D) were formed as a result of volatilization during carbonization/graphitizationcycles, whereas the cracks/fissures (position E) were formed due to shrinkage of the matrix carbon during cool down from the processing temperature (ref. 3). A large proportion of the voids are infilled with a resinous carbon added to complete the densification process.

128

Carbon-Carbon Matrax Materials

Figure 10. Photomicrograph of CC composite (PAN-basedfibers plus coal tar pitch matrix). The development of the mesophase within the interbundle matrix of the composite is not always uniform. The structure of some of the matrix regions is coarse-grained mosaic whereas the structure of other regions is fine-grained mosaic. This variation in the optical textures of the binder-coke causes differences in mechanical properties and chemical reactivity of the composite (refs. 85 and 86). Variations in processing conditions have reportedly affected the matrix orientation. For instance, Cranmer et al. (ref. 87) concluded that the alignment of the graphitic planes under low-pressure carbonization is primarily controlled by the flow motion of mesophase. High pressures have been reported to favor a transversely aligned matrix (ref. 88). Other factors also seem to be important because Murdie et al. (ref. 89) have recently reported that both parallel and transversely oriented graphite planes were produced in a specimen carbonized under atmospheric pressure. Although local variations in crystallographic orientation are observed, pitch matrix composites generally exhibit microstructures in which the fibers are surrounded by coaxial, graphitic sheaths.

129

CC Materials and Composites

0jJ-$&?-yYM~ @ Q(k

~6y6/ ~ -(-Jy@ %~7 id&-& O)Y,t(?,&?F /Q/” \i&&=sq(-p-&yo *-.?.z? 1,-

1

@ 3 ? + \

$\-z>F ( \?gq-J!.-

%ggh$J)[7=5/5*<=

--

)2&=&55&

\+&%&@&@ f s ==4

1

Figure 11. Matrix alignment in fiber-reinforced pitch composites (ref. 90). Reprinted with permission.

A schematic representation of the alignment of the mesophase pitch matrix within a unidirectionally reinforced material is shown in figure 11 (ref. 90). In this figure, the prismatic edges of the graphitic planes are schematically shown as flow lines distributed between the individual fibers. A similar flow pattern was observed in a PAN-pitch composite by etching the composite surface with chromic acid or atomic oxygen (ref. 59). The example shown in figure 12 is an SEM micrograph of a PAN-pitch composite etched in this way. This micrograph clearly shows that the intrabundle matrix platelets (the basal planes of the platelets are parallel to the platelet broadface) are parallel to the fiber surface. As can be seen from this transverse section, the fiber bundles appear fully densified, although the interface between the fiber and matrix appears to be quite reactive since it has been etched preferentially. The presence of a central core in the PAN fiber shown in this figure has been discussed in the literature and is attributed to graphitization of partially stabilized PAN fibers (ref. 91). The interface between an individual fiber and the interbundle mesophase matrix is shown in figure 13. In this case, two rather large fissures are shown that occur

130

Carbon- Carbon Matrix Materials

Figure 12. SEM micrograph of transverse section of fiber bundle showing jiberlintrabundle matrix interface after etching with atomic oxygen at 100°C for 3 hr. at some distance from the fiber surface. It appears, therefore, that in some areas the bonding between the graphite-like layers formed within the pitch matrix is weaker than that between the pitch and fiber. This feature is quite different from that which occurs in pitch fiber-reinforced phenolic resin matrix materials where shrinkage during first carbonization causes the matrix to pull away from the surface of the fiber. Figure 14 is a scanning electron micrograph of the same composite showing the interface between the interbundle matrix and a longitudinally oriented fiber. This micrograph shows that near the fiber-matrix interface, the matrix structure is parallel to the fiber surface in the longitudinal orientation for a considerable distance. This alignment only exists for a distance of 2 pm to 6 pm from the fiber surface and becomes more random in the bulk of the interbundle matrix. Zimmer and Weitz (ref. 92) has indicated that even when a strong magnetic field is applied to attempt reorientation of the mesophase, the mesophase is still aligned parallel to fiber surfaces for radial distances up to 6 pm. The results of this SEM study indicate that each fiber is surrounded by a highly oriented thin sheath of matrix. This sheath

131

CC Materials and Composites

Figure 13. SEM micrograph of fiber after etching with atomic oxygen at 100°C for 3 hr.

Figure 14. SEM micrograph offractured surface of interbundlelfiber interface after etching with atomic oxygen at 100°C for 3 hr.

132

Carbon- Carbon Matrix Materials thickness (0.2 pm to 4.0 pm) can vary within a fiber bundle owing to complex flow patterns generated during mesophase formation and different interfiber distances.

A TEM bright-field image showing a typical transverse morphology of an intrabundle matrix region in the same composite is shown in figure 15. As can be seen, the crystallites of the intrabundle matrix exhibit a flow-type morphology with basal planes oriented roughly parallel to the fiber surface. This geometry is similar to that described in Zimmer and White’s model (ref. 93). The crystallites of the matrix near the fiber surface (position A) or between two closely spaced fibers exhibit better alignment with the fiber surface than those farther away from fibers. Microfissures, indicating weak bonding, are clearly revealed within the matrix and along the fiber-matrix interface. A typical morphology of an interfacial region in the same composite is shown in figure 16. The bonding between the matrix and fiber appears discontinuous because some matrix graphite crystallites are well bonded to the fiber while other crystallites are poorly bonded. This interfacial morphology is similar to the fissuretype interface classified by Ragan and Marsh (ref. 86) in their study of bulk binder, coke-filler coke composites. This higher magnification micrograph shows more clearly the shape, size, and distribution of the numerous microcracks along and near the fiber-matrix interface (position B). The matrix crystallites adjacent to the interface are rather small and irregularly shaped, while those farther away from the fiber surface appear larger. The microcracks in the matrix, formed between and parallel to the graphitic platelets, generally have a sharp, lenticular shape. Near the fiber surface, these microcracks become smaller and denser. The crystallites in the pitch matrix appear much larger and more graphitic than those in the turbostratic PAN fiber; this is clearly revealed in the lattice fringe images shown in figure 17. A rough estimate from these and other lattice images indicates that the crystallite length in the mesophase pitch matrix is at least 10 times larger than in the PAN fiber. In summary, the microstructural investigation of pitch matrix composites indicates

1. The carbon produced from pitch precursor materials is highly graphitic and generally exhibits a flow-type morphology. 2. Although the bonding between the PAN fibers and the mesophase pitch matrix appears continuous at low magnifications, high-magnification TEM studies reveal numerous microcracks along the fiber-matrix interface, each separated by what appear to be well-bonded regions.

3. The microcracks within the matrix are smaller and denser when they are near the interface.

133

CC Materials and Composites

Figure 15. Photomicrograph of interfaces between PAN fibers in pitch matrix composite.

Figure 16. Photomicrograph of interface between PAN fiber and pitch matrix.

134

Carbon-Carbon Matrix Materials

(a)Longitudinal PAN fiber.

-

100A

(b) Longitudinal pitch fiber. Figure 17. Photomicrograph of comparison of lattice fringe images taken from (a) PAN fiber and (b)pitch fiber. 4. The crystallite size in mesophasic matrices is large compared to that in PAN fibers.

5. The crystallites in the matrix are aligned parallel to the fiber surface in a range approximately 0.2 pm to 6 pm from the fiber surface; thereafter, they tend to become more random.

Resin Matrix Composites Carbon-carbon composites formed from reinforced resins are processed by carbonizing (and graphitizing) carbon fiber-reinforced phenolics (or other suitable resin matrices). A typical microstructure of a phenolic resin reinforced with pitch fibers is shown in figure 18. The composite shows good bonding characteristics between the fibers and the resin matrix. Polarized light microscopy, combined with a retarder plate, rendered the phenolic matrix purple. The color of the phenolic matrix structure did not change with the rotation of the microscopic stage, thus indicating that the resin is optically isotropic. A carbon fiber reinforced phenolic

135

C C Materials and Composites

Figure 18. Photomicrograph of as-moldedpitchfibers in phenolic resin (hot pressed).

resin of this type typically contains less than 3 percent macroporosity, and exhibits a density of about 1.5 g/cm3. Carbonization of resin matrix composites involves heating the material in an inert environment to a temperature of 1000°C. During this heat treatment process, volatiles are emitted in the form of vapors and gases such as H20, methane, and hydrogen. This process causes the resin to shrink in volume so that the porosity of the composite increases to 30 to 40 percent, while the bulk density typically decreases to between 1.2 g/cm3 to 1.4 g/cm3. depending on fiber weave and volume fraction. The resin contracts extensively on carbonization (illustrated in fig. 19). Shrinkage has been so excessive that the resin has shrunk away from almost all of the fibers present in this area. Interference colors indicated that most of the resin structure remained isotropic; however, in some regions of the intrabundle matrix, it was obvious that an anisotropic structure had developed. The development of this anisotropic structure was restricted to those areas where the distance between two fibers was very small, i.e., <2 pm.

136

Carbon- Carbon Matrix Materials

Figure 19. Photomicrograph of carbonized pitch fiber-reinforced phenolic resin.

Figure 20. Photomicrograph of pitch fiber-reinforced resinlCVI composite etched with atomic oxygen.

137

CC Materials and Composites An SEM micrograph of a graphitized pitch fiber/phenolic resin plus CVI matrix CC composite etched with atomic oxygen is shown in figure 20. Based on the orientation of the microcracks within the resin char (position B), the matrix appears highly oriented. The interface between the CVI and resin char is crenulated and can be expected to facilitate mechanical bonding through an interlocking mechanism. No evidence for a three-dimensional graphite structure existed within the resin char in this specimen, processed at relatively low temperatures. However, Hishiyama and his co-workers (ref. 94) reported that graphitization can indeed occur when phenolic resin precursor CC’s are heated to very high temperatures. Furfural alcohol matrices align relatively easily. A typical example of this alignment is shown in figure 21, an optical micrograph of a rayon fiber cloth impregnated with multiple cycles of furfural alcohol resin. The fiber bundles appear fully densified. The intrabundle matrix region exhibits a high degree of basal plane alignment parallel to the fiber surface. The interbundle matrix also exhibits a high degree of basal plane alignment parallel to the fiber bundle/matrix boundary (position F). In an attempt to increase the yield of resin matrix composites and to reduce shrinkage, the addition of fillers has been investigated (ref. 95). Figure 22 is an optical micrograph showing the structure of the as-molded composite, composed of rayon fibers and phenolic/carbon-black matrix. The matrix exhibits a granular appearance owing to the presence of the carbon black (filler). In addition, filler rich (position C) and filler depleted (position D) regions exist, thus supporting the known observations of the difficulties associated with obtaining a uniform distribution of particles in filler containing resin matrices. The microcrack present at position E has propagated along individual fiber-matrix interfaces, indicating that the fibermatrix bond is weaker than the matrix strength. This fact is confirmed by the TEM study. These shrinkage cracks can extend across the total fiber bundle before stopping at the junction with neighboring fiber bundles of different orientation. At this magnification, there appears to be little difference in bonding characteristics and level of macroporosity between doped and undoped resin matrix composites.

A TEM bright-field image of a phenolic resin matrix doped with carbon black is shown in figure 23. Resin platelets aligned parallel to the faceted carbon-black surfaces are shown at position D. This alignment suggests that the carbon-black particles “grow” (position G). The carbon-black particles not only restrict shrinkage of the resin but also act as nuclei for the development of preferred orientation during processing. In some areas, preferred orientation does not develop as illustrated at position H. The microcracks present in the specimen shown in figure 23 occur between the basal planes of the oriented resin and are concentric to the center of the carbon-black particles. Such cracks are not observed in neat (pure) resin composite matrices. It should be noted that although the presence of carbon black inhibits the shrinkage

138

Carbon- Carbon Matrix Materials

Figure 21. Photomicrograph of rayonlfurfural alcohol CC composite.

~.

I 7

'

.

.+

P t 1

? '

,

'E .'.

Figure 22. Rayon fiber-reinforced resin with filler (as molded).

139

CC Materials and Composites

Figure 23. Bright-field image of graphitized phenolic resin containing carbon-black particles. of the matrix, it does not eliminate it. The matrix still shrinks away from the fiber surface during processing, although for much smaller distances. This shrinking effect is shown in the TEM micrograph of a resin-CVI hybrid matrix composite, (fig. 24(a), ref. 96). In this image, it can be seen that carbonization causes the fillerdoped resin to shrink away from the fiber surface for a distance of about 1 pm. This space was subsequently infilled with a CVI layer. The interface between the fiber and CVI layer appears continuous while that between the CVI and the filler-doped resin contains short cracks. The dark-field image of figure 24(b) indicates the continuity of the fiber-matrix interface and the position of the cracks along the CVI-resin interface. Also indicated in this micrograph are the highly aligned basal planes of the filler modified matrix regions. The diameter of these regions varies between 0.1 pm and 0.5 pm. The selected area diffraction patterns indicate the highly oriented nature of the fiber (fig. 24(c)), the isotropic nature of the CVI deposit (fig. 24(d)), and the isotropic (the diffraction spots are due to the presence of the carbon black) nature of the resin (fig. 24(e)). It is interesting to note that an isotropic CVI layer bonds well to either pitch or PAN-based fibers, whereas CVI deposits with a laminar microstructure

140

Carbon- Carbon Matrix Materials

Figure 24. TEM bright field (a), dark field (b),and SAD’S (c), (d), and ( e ) of pitch fiber-reinforced CVI@lled resin composite (ref. 96). Reprinted with permission.

do not. The higher magnification TEM micrograph (fig. 25) indicates that the CVI-resin interfacial cracks can propagate along the interface. In summary, the microstructural investigation of resin matrix composites indicates 1. The carbon produced from resin precursor materials is usually isotropic but can become highly oriented. However, the degree of preferred orientation generated depends on resin type and process conditions. Phenolic resins are more difficult to orient than furfural alcohols. Local regions of phenolic and furfural resins can become graphitic when heated to graphitization temperatures. 2. Photomicrographs suggest that the bonding between resin char matrices and pitch-based fibers is relatively weak. (Fiber-surface treatments may improve this bonding.)

141

CC Materials and Composites

Figure 25. Photomicrograph of interfacial cracks propagating through graphitized resin interface. 3. Addition of carbon black to phenolic resin causes the preferred orientation to develop, concentric with the original carbon-black particles during processing. 4. Microcracks are generated within the oriented resin char around the carbonblack particles that presumably contribute to the reduction of apparent shrinkage of the bulk resin.

CVI Matrix Composites Infiltrating a carbon fiber preform with a reactive gas leads to the deposition of a carbon matrix and directly produces a CC composite. No subsequent thermal degradation step is necessary. However, pyrolytic carbon tends to deposit preferentially on the surface of the composite being fabricated, thereby blocking surface porosity. An example of this surface layer is shown in the optical micrograph (fig. 26). A very thick surface layer with associated growth cones is observed at position K. Infiltration can only be continued after terminating the densification process and removing the surface crust. Most isothermal depositions are carried out at a temperature low enough to limit the rate of surface deposition, yet high enough to assure an economically viable commercial process. The temperature of reaction has been shown to sensitively affect the density and the 142

Carbon-Carbon Matrix Materials

Figure 26. Optical photomicrograph of surface region of PAN-CVI CC composite (graphitized).

microstructure of the deposit, which in turn will affect the resulting properties of the CC product. A typical photomicrograph of CVI carbon infiltrated into a PAN fiber preform is shown in figure 27. This micrograph shows the presence of large voids in some parts of the specimen. The CVI deposit consists of two structures: an isotropic phase immediately adjacent to the fiber surface (1 pm thick) and a second highly oriented lamellar structure. The extinction contours generated indicate that the basal planes of this second structure are circumferentially oriented approximately parallel to the fiber surface. The thickness of the deposit varies with fiber spacing and is thickest within the large open pores. A higher magnification micrograph of the same composite is shown in figure 28, which shows more clearly the structure of the CVI deposits. In this particular region, the CVI deposit appears to have pinched off the pore so that subsequent infiltrations would not be expected to be effective. The bonding between the isotropic CVI layer and the fiber, and between the two CVI layers looks continuous, with no evidence of fissures or cracking.

143

CC Materials and Composites

Figure 27. Optical photomicrograph of interior region of PAN-CVI CC composite (graphitized).

Figure 28. Optical photomicrograph of closed porosity in PAN-CVI CC composite (graphitized).

144

Carbon- Carbon Matrix Materials The continuous nature of the interface between the fiber and the isotropic matrix is confirmed by the TEM dark-field micrograph (fig. 29). The SAD studies suggest that the isotropic layer is composed of small randomly oriented crystallites, while the laminar structure consists of larger highly oriented crystallites. It is interesting that the interface between PAN fibers and isotropic CVI and between pitch fibers and isotropic CVI is continuous. Thus, the development of such an interface appears to depend more on CVI structure than on individual fiber type. The laminar CVI carbon contains many narrow slit-shaped microfissures that are generally <1 pm in length and <0.1 pm in width. This type of cracking is similar to that observed in pitch matrices and indicates that the strength of the bonds between individual crystallite platelets within the lamellar matrix is weaker, in some places, than the bond between the isotropic CVI and the PAN fiber.

Figure 29. TEM dark-field image of CVI deposit on PANfiber showing gradual transition from grainy isotropic to laminar structure. A photomicrograph of pitch fibers infiltrated with CVI carbon is shown in figure 30. In this case, conditions of deposition were such that no detectable isotropic layer was observed adjacent to the fiber surface. The CVI carbon exhibits optically anisotropic characteristics and has a rough lamellar structure. At this low magnification, it is difficult to observe if the interface between the fiber and matrix is microcracked, or if microcracks are present within the CVI deposit. Figure 31

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CC Materials and Composites

'I Figure 30. Photomicrograph of pitch fiber-reinforced CVI composite.

shows an SEM micrograph of the same composite after etching with atomic oxygen. The interface between the fiber and CVI matrix is cracked. It should be noted that while the fiber matrix interface in this specimen is cracked, the corresponding interface shown in figure 29 is not. Since both composites were exposed to similar heat treatments, it can be concluded that the bonding between isotropic CVI and fiber is stronger than the bonding between lamellar CVI and fiber. A TEM bright-field micrograph of the same specimen is shown in figure 32. Cracks are observed in both CVI layers. The interface between the two CVI layers appears relatively continuous, whereas extensive interfacial cracking occurs between the pitch fibers and the first CVI layer. Although etching preferentially attacks interfaces and can exaggerate the size of existing cracks, the technique is useful in detecting different structures and their boundaries. For instance, after etching the previously described specimen in atomic oxygen, examination with the SEM revealed the existence of two distinct CVI layers that were not detected optically (fig. 31). The effect of such a weak fiber-matrix bond is clearly shown in the fracture surface of this composite (fig. 33) where the fibers have been completely pulled out from their surrounding CVI matrix.

146

GC Materials and Composites

Figure 33. SEM micrograph showing fiber pullout. In summary, the microstructural investigation of CVI matrix composites indicates 1. Pyrolytic carbon tends to deposit preferentially on the surface of the composite being fabricated, thereby blocking surface porosity. 2. The structure of carbon produced from CVI deposits can vary from isotropic to highly oriented anisotropic, depending on processing conditions.

3. Bonding between isotropic CVI matrices and fibers is stronger than the bonding between lamellar CVI matrices and fibers.

4. The lamellar CVI deposits contain many narrow, slit-shaped microfissures that are generally <1 pm in length and <0.1 pm in width. The isotropic CVI deposits contained no cracks.

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Carbon- Carbon Matrix Materials

Influence of Matrix on Composite Properties General Background When selecting a resin, metal, ceramic, or carbon-matrix composite for a given application, three criteria are generally used: the composite must have the desired physical and mechanical properties; it must be capable of being processed or manufactured into the desired shape; and it must be economical to produce. For most composites, the primary consideration for a given application concerns the properties of the reinforcing fibers. The mechanical properties (strength and modulus) of carbon fibers are also related to the physical properties (thermal and electrical conductivity and coefficient of thermal expansion). Because of the interrelationships that exist between the microstructure and the elastic and physical properties (thermal and electrical conductivity and coefficient of thermal expansion), the choice of a fiber based on one property usually determines the value of the other properties. For instance, the microstructure of very highmodulus fibers usually consists of long fibrils that are almost perfectly aligned parallel to the fiber axis. As a result, the transverse modulus will be relatively low, the thermal and electrical conductivity will be high in the longitudinal direction, and the thermal expansion coefficient will be small or negative.

A composite matrix usually serves to protect the reinforcement fibers from damage or reaction with the environment, to provide some measure of support in compression, to provide adequate matrix-dominated properties, and to provide a continuity of material. This last property is important in electrical and thermal applications and is particularly important in mechanical applications since load must be transferred to the fibers through the matrix. In this respect, a load can be transferred to the fibers across a chemically or physically bonded interface or across a mechanically interlocked one formed by the matrix shrinking onto and thereby gripping the fiber surface. The properties of the matrix dominate the properties of composites in any direction in which fibers are not aligned. Such properties include the transverse tensile strength and modulus, interlaminar shear strength, thermal expansion coefficient, and electrical conductivity. Composite materials reinforced with continuous fibers are complex materials that, for a given weight, exhibit specific mechanical properties almost always superior to those exhibited by conventional metals and alloys. Three basic composite types can be conveniently discussed in terms of maximum temperature of use: reinforced polymers, which are restricted to use at relatively low temperatures; reinforced metals, which can be used at intermediate temperatures; and reinforced carbons and ceramics, which can be used up to very high temperatures.

149

CC Materials and Composates In order for the mechanical properties of a continuous fiber-reinforced composite to be superior to an unreinforced material, the modulus and strength of the fiber reinforcement must be greater than the matrix. In addition, there must be chemical, physical, or mechanical bonds formed between the fibers and matrix that are strong enough to transfer load between individual fibers and between fiber layers. In estimating the mechanical properties of specific composites, assumptions are made with respect to the contribution of the matrix. For instance, the elastic modulus of typical polymers is usually so small that the contribution to the mechanical properties of the composite is ignored. Conversely, the modulus of metals is much larger; thus, the stiffness and strength of the composite will reflect a significant contribution to the tensile modulus. The modulus of polycrystalline carbon is small; however, the basal planes of carbon within the matrix of a composite can become highly oriented with respect to the fiber axis, showing a large increase in modulus. The modulus of carbon is very large in the direction of its basal plane, but very small in the out-of-plane direction; thus, any contribution to the modulus of the composite will depend on the alignment of the basal planes. In the previous section of this chapter, we have shown that a carbonaceous matrix highly oriented with respect to the axis of the fibers is most often produced from a pitch precursor. A range of structures (isotropic and anisotropic) can be produced from CVI and resin. However, basal plane alignment can be more easily achieved in CVI.

Elastic Modulus When measured parallel to the fiber direction, the elastic modulus of a unidirectionally reinforced composite can be calculated to a first approximation from the rule of mixtures

+

Ec = Em(1 - V f ) E f V f

(1)

where Em and E f are the moduli of the matrix and fiber measured parallel to the fiber axis and V f is the fiber volume fraction. When the modulus of the matrix is much smaller than the fibers, equation (1) reduces to

E, = E f V f

(2)

The above expressions give reasonable values for resin and metal matrix composites. Experiments carried out by Perry and Adams (ref. 97), using a variety of fiber types with SC- 1008, FF-26 phenolic resins, or polyphenylquinoxaline and various furfural or furfural blends as impregnants, measured modulus values for CC which were much larger. Further evidence for the large contribution of matrix to the stiffness of CC composites has been reported by other workers (refs. 98 through 101). For instance,

150

Carbon- Carbon Matrix Materials Fitzer and Huttner (ref. 98) reported modulus values for pitch-char matrix materials that were twice that computed using equation (2), implying that the stiffnesses of the matrix and fiber were equal. The modulus values for resin char matrices were 40 percent higher. Further confirmation of this large stiffness contribution of the matrix was provided by the present authors in testing a unidirectionally reinforced sample of PAN fiber T-300 reinforced pitch material cut from a graphitized multidirectionally reinforced composite. The composite sample containing about 50 volume percent fiber reinforcement exhibited a modulus value of 45 Msi (310 GPa). Since the work reported by Becker (ref. 102) indicates that the modulus of graphitized T-300 fibers is about 58 Msi (400 GPa), it can be inferred from equation (1) that the modulus of the carbon matrix in the direction of the fiber axis was 32 Msi (220 GPa). In the previous section dealing with microstructures, we established that an appreciable degree of preferred orientation develops on heat treating pitch, resin (or CVI) matrix precursors. Therefore, we concluded that this directionality significantly influences the modulus values of the resulting composite, and for a given volume fraction of fibers with similar bonding, the strong orientation tendency of pitch and highly oriented CVI-based matrix materials will produce the stiffest composite when measured parallel to the fibers.

Tensile Strength

In a mechanically loaded composite, it is usually assumed that the fibers and the matrix are bonded together and that no differential movement exists between them. The strain in each part of the composite can then be considered identical and equal to the total strain on the composite. For an elastic response, the stress on a unidirectionally reinforced composite loaded parallel to the fiber axis a, can be computed by multiplying the elastic modulus in that direction E, by the measured strain E,, i.e., ac = E,€c

(3)

or, combining (1) with (3) f7,= E,€,

= E,€,(l

-

V,)

+ Ef€fVf

(4)

The failure strain of a polymer or a metal E , is substantially larger than that of the reinforcing fibers E,, and failure of this type composite occurs when the fibers break. Equation (4) is therefore used to predict failure of the composite by assuming at fracture E , = 6,. In contrast, the failure strain of the matrix E , in most CC’s is smaller than the failure strain of the fibers Ef and, in a well-bonded composite, failure occurs when the the matrix fails. Failure occurs rapidly and results from the catastrophic propagation of a crack initiated at some weak point in the composite.

151

CC Materials and Composites The failure of CC composites can again be discussed with reference to equation (4). If it is assumed that the maximum failure strain of carbon matrices is approximately 0.3 percent, then the composite will fail at that strain. Also, if it is assumed, for simplicity, that the matrix carries no load, the strength of a carbon matrix reinforced with 50 volume percent of 75 Msi (517 GPa) high modulus fibers would be about 112 ksi (772 MPa). Conversely, the strength of the same matrix reinforced with 30 Msi (206 GPa) low modulus fibers would be about 45 ksi (310 MPa). Since higher modulus fibers are usually weaker, we note the curious result discussed by Fitzer and Huttner (ref. 98) that for composites that fail at the matrix failure strain, the strengths can be larger when stiffer; usually weaker fibers are used as the reinforcement. The failure of a composite whose failure is matrix-dominated can occur at lower strains when differences in thermal expansion between the fiber and matrix cause the matrix to be prestressed in tension on cooling from carbonization temperatures. Since the matrix and fiber must be bonded together in order to prestress the matrix, very low strengths might be expected from well-bonded materials. Thomas and Walker (ref. 103) studied phenolic resin precursors reinforced with three types of commercially available carbon fibers. Some of the fibers were surface treated to promote bonding; some were not. In every case, the matrix-dominated properties (transverse flexural strength, modulus, and interlaminar shear) of well-bonded phenolic char-matrix composites were superior to those exhibited by less well-bonded composites. Unfortunately, the longitudinal strengths of the well-bonded material were extremely poor and the work of fracture was low. In contrast, the matrix-dominated properties were poor, and the longitudinal strength of fully processed CC composites reinforced with nonsurface treated fibers was superior to those containing surface-treated fibers. An additional series of experiments were performed by Fitzer et al. (ref. 104) who studied the effect of the degree of fiber oxidation on the mechanical properties of carbonized resin materials reinforced with SIGRAFIL HF and SIGRAFIL HM carbon fibers. After initial carbonization, but before subsequent densification, they found, relative to material reinforced with nonoxidized fibers, that increasing oxidation improved the strength of the material reinforced with the HM fibers, but decreased the strength of the material containing the high-strength fiber. Manocha et al. (ref. 105) performed a similar experiment using surface-treated Toray M40 fibers and a matrix of furfural alcohol condensate. It was found that after the initial carbonization, the strength of the composite containing the surfacetreated fibers (oxidized) was poorer than that containing untreated fibers. The effect was similar to that observed by Fitzer et al. on SIGRAFIL HF materials. Manocha et al. graphitized their first carbonized material without densifying it. The strength of the surface-treated fiber composites increased by a factor of 3 or more while the strength of the nonsurface-treated materials decreased by nearly the same amount. In effect, Manocha showed that the composite reinforced with the surface-treated fibers exhibited the highest strength when graphitized.

152

Carbon- Carbon Matrix Materials In Fitzer's work, subsequent pitch densification cycles increased the strength of the untreated SIGRAFIL HM material to almost 100 percent (calculated assuming no contribution of the matrix), whereas only a 40-percent strength realization was obtained from the nonsurface-treated HF material. A reasonable conclusion appears to be that even without surface treatment, the SIGRAFIL HF fibers reacted to form a strong bond with the phenolic novolac matrix, whereas the HM fiber did not. A similar effect had been reported by Thomas and Walker (ref. 103) working with phenolic resin subsequently densified with CVI. It is reasonable to conclude from the discussed data that matrix precursor-fiber combinations that promote strong bonds (with or without surface treatment) will not produce a high-strength composite when measured parallel to the direction of the fibers. The addition of graphite powder is often used to decrease the shrinkage tendency of phenolics during carbonization and to improve the carbon yield (ref. 106). Since the decrease in the thermal contraction should also produce a corresponding decrease in the internal stress generated in the matrix on cooling from carbonization temperatures, the strength of composites whose failure is matrix dominated should be improved. The results of Fitzer et al. (ref. 107) confirmed this when they demonstrated that the strength translation of HF fiberreinforced material increased from 40 to 60 percent when SO-percent graphite powders were added to the resin matrix before carbonization. We have shown in a previous section (Microstructure of CC Matrices) that (in the absence of fiber surface treatments) weak interfaces will occur naturally in some matrixfiber combinations, while strong interfaces will occur in others (refs. 67, 77, 78, and 95). Specifically, the nature and strength of the fiber-matrix bond appear dependent on the reactivity of the fiber (i.e., fiber type), the reactivity of the matrix, and the type and degree of fiber-surface treatment. Fitzer has shown that an improvement in properties is achieved with densification (ref. 107). He found that the strength of composites containing nonsurface-treated HM fibers could be increased from 10 percent after the initial carbonization to 90 percent of the theoretical strength after four densification steps, while the strength of the composite containing nonsurface-treated HF fibers increased to only about 40 percent. These data suggest that a reaction occurred naturally between the fiber and matrix in the case of HF fibers and a limited reaction occurred throughout the densification process for the HM materials. Perry and others (refs. 97, 108, and Herrick") working with phenolics and different fibers found a maximum of only 40- to 50-percent translation of strength indicating some degree of reaction; however, since their preparation technique involved a number of graphitizations, other factors may have contributed to the low strength values.

*J. W. Herrick, Fiber Materials, Inc. Private communication.

153

CC Materials and Composites Fitzer also found (ref. 98) that resins which form strong fiber-matrix bonds and also exhibit the highest shrinkage should be avoided if high translation of fiber strengths is required. He suggests the use of high carbon-yielding precursors that do not form strong bonds and that exhibit minimum shrinkage. Data presented indicate that pitches (rather than resins) will produce better translation of fiber strengths for fewer densification cycles.

Matrix Dominated Properties The out-of-plane, off axis and transverse properties of unidirectionally reinforced composites depend mostly on the properties of the matrix or matrix-fiber bond. Specific properties of interest include transverse modulus, strength, and interlaminar shear strength. Using a simple model, an approximate value for the transverse modulus of a unidirectionally reinforced composite ECT can be estimated ECT = & n E f / ( v m E f-k VfEm) (5) where E , and E f are modulus values for the transverse direction of the composite; V, and Vf are the respective volume fractions of the matrices. Although this expression neglects porosity and some other important effects (ref. 109), it can be used to reach the reasonable conclusion that in well-bonded composites, for the usual fiber volume fractions ( S O to 60 percent), the transverse modulus will never be greater than about twice the modulus of the matrix. In the absence of constraint, transverse strength can never be greater than the strength of the matrix (and will degrade at a rate that depends on volume fraction of poorly bonded fibers). It is believed that in well-bonded composites the transverse strength will be independent of volume fraction of fibers and approximately equal to the strength of the matrix. Conversely, in a poorly bonded composite, a dependence on volume fraction will be observed and strengths appreciably lower than the strength of the matrix will be expected. The results of Perry and Adams (ref. 97) who studied three types of resin matrices and different fiber types support the conclusion that larger values of transverse strength, modulus, and interlaminar shear are obtained from composites reinforced with well-bonded fibers. Although several fiber types were investigated, it appears that bonding is controlled by the reactivity of the matrix and the fiber. All properties improved with densification up to the fourth cycle and, in general, heating to graphitization, although degrading the transverse tensile strength caused no degradation in interlaminar shear values. Carbon-fiber-reinforced carbonized Borden SC-1008 phenolic exhibited a maximum transverse strength of about 870 psi (6 MPa) which degraded to about 700 psi (4.8 MPa) or less when heated to graphitization temperatures. A similar degradation from a maximum of 680 psi (4.7 MPa) (nongraphitized) to less than 500 psi (3.4 MPa) (graphitized) occurred for carbon-fiber-reinforced PPQ resin char and from 760 psi (5.2 MPa) to 450 psi (3.1 MPa) for the reinforced FF-26 resin char. In most cases, values of interlaminar

154

Carbon- Carbon Matrix Materials shear varied between 3000 psi to 3600 psi (20.7 MPa to 24.8 MPa) for the SC1008, about 2100 psi (14.5 MPa) for the PPO, and about 2600 psi (19.9 MPa) for the FF-26. Interlaminar shear strengths for carbonized pitch densified materials appear to be superior to resin-char matrices. Values of 3915 psi (27 MPa) have been quoted for resin pitch-densified material (ref. 104). In the late 1970’s, Thomas and Walker (ref. 103) studied CVD-densified carbonized phenolic resin chars reinforced with a series of commercially available fibers, some of which were surface treated. Apart from naturally reactive fibers (high strength, low modulus), surface treatment improved the transverse flexural strength slightly while the interlaminar shear was increased significantly. Subsequent carbonization and redensification always produced composites with matrixdominated properties lower than the reinforced precursor; however, the superiority of the surface-treated material always remained after heat treatment. Unfortunately, as discussed before, a good fiber-matrix bond reduced the value of the fiberdominated longitudinal strength, and the fracture mode became very brittle. Depending on fiber type and the presence or absence of surface treatment, values of transverse flexural strength exhibited by the CC material vaned from about 2030 psi (14 MPa) to 4350 psi (30 MPa); values of interlaminar shear varied from 2320 psi (16 MPA) to 3915 psi (27 MPa).

The apparent superiority in matrix-dominated properties of the CVD-densified material when compared to resin-densified material is supported by the work of Walker and Lee (ref. 1IO) who indicate that both transverse tension and interlaminar shear of 2-D CC are superior if the matrix is prepared by chemical vapor deposition. These authors indicated interlaminar tensile strengths of about 790 psi (5.4 MPa) and interlaminar shear strengths of 2340 psi (16.1 MPa) for carbonized resin precursor material compared with 1460 psi (10.1 MPa) and 2740 psi (18.9 MPa) for CVD material. In this work, it was pointed out that the superiority of CVD was only true for thin materials (0.1 in.) because reductions of both interlaminar tensile and interlaminar shear strengths occurred when the thickness of the CVD-densified material approached 0.5 in. This property degradation was shown to be related to the degree of difficulty in fully densifying thick sections using CVD. The results reported in the literature are difficult to interpret because varying one parameter while keeping other factors constant is difficult. The matrix-dominated properties depend on fiber type, surface reactivity, matrix type (resin precursor, CVI microstructure, and pitch precursor), matrix reactivity, fiber volume fraction, processing condition, degree of densification, fiber orientation, and laminate thickness. By optimizing all these parameters, the matrix-dominated properties can be maximized; however, the greatest improvement in these properties is obtained by placing fibers in the appropriate directions.

155

CC Materials and Composites Two-Dimensional Reinforcements Two-dimensional composites can be fabricated from unidirectionally reinforced prepreg material in which individual plies or fabrics are stacked in precomputed orientations. Fabric is the preferred material used to process CC material. Fabrics contain undulating fibers, however, and some loss in strength is realized when compared to simple cross-plied unidirectionally reinforced types. Typically, strengths and modulus of 2-D CC composites measured parallel to the fibers average 30 000 psi (206 MPa) and 16 Msi (1 10 GPa), respectively, while interlaminar tension and shear strengths are 1000 psi (6.9 MPa) and 1200 psi (8.3 MPa). These results represent an increase of more than 30 times in transverse strength and modulus depending on fiber volume fraction. The data showed a drop to half, or even a quarter, of the longitudinal tension strength values formerly exhibited by unidirectionally reinforced composites. Although the properties in the transverse direction have improved significantly, the properties through the thickness are still relatively poor (ref. 110).

Three-Dimensional Reinforcements Three-dimensional composites have been designed in order to improve the poor, through-the-thickness, interlaminar tension and shear strength values exhibited by the 1-D and 2-D CC composites. Various methods of doing this include pattern nesting, angle interlock, and ply interlock as described by Walker and Lee (ref. 111) and 3-D weaving by Ransone et al. (ref. 110). Methods for producing thick sections have also been discussed by McAllister and Lachman (ref. 4). For an orthogonal construction, Schmidt (ref. 112) quoted by McAllister and Lachman (ref. 4) obtained tensile strengths of 15 ksi (103 MPa) in the 2 (through-thethickness) direction and 5 ksi (34.5 MPa) in both the X and Y directions on pierced fabric-reinforced CC. Girard (ref. 113) obtained 10 ksi (70 MPa) in all directions for CVD-densified material with slightly better properties obtained from resindensified CVD precursor material. Levine* reported values of 45 ksi (310 MPa) in the Z direction and 15 ksi (103 MPa) in both the X and Y directions of the 3-D materials that he examined. Correspondingly high values for modulus were also reported. The absolute values of strengths for the 3-D materials cannot be directly compared with unidirectionally 2-D or other 3-D materials because the fiber volume fraction in each direction varies; however, the conclusion can definitely be made that placing fibers in the three orthogonal directions will improve otherwise poor, matrix-dominated properties. *Levine, A.: High Pressure Densified Carbon-Carbon Composites-Part 11: Testing. Extended Abstracts-12th Biennial Conference on Carbon, American Carbon SOC., 1973. Unpublished paper.

155

Carbon- Carbon Matrix Materials The following points can be made in summary:

1. The elastic modulus of CC composites is dependent on the modulus of the carbon matrix and the modulus of the fibers. 2. Good bonding between the fiber and the carbon-matrix precursor will result in a better bonded CC composite.

3. Bonding depends on surface treatments, type of fiber, and type of matrix precursor.

4. Fibers subjected to higher processing temperatures (high modulus) generally do not bond as well as those subjected to lower processing temperatures. 5. Good bonding between the fibers and matrix results in a composite that fails in longitudinal tension at the failure strain of the matrix.

6. Good bonding tends to prestress the matrix and can cause failure to occur at strains lower than the failure strain of the unrestrained matrix. 7. High longitudinal strength is obtained by providing weak interfaces (or cracks) perpendicular to the path of the catastrophically propagating crack.

8. Better matrix-dominated properties are obtained from well-bonded materials. 9. The matrix-dominated mechanical properties of CVD-densified material are superior to that densified with phenolic resin. 10. Compared to resins and thick CVD-densified composites, pitch matrix materials produce higher translation of fiber strengths and higher interlaminar shear strengths.

Acknowledgments We wish to acknowledge J. He, C. Kocher, and E. Forrester for their contributions in preparing the specimens for examination in the optical scanning and electron microscopes, J. Bazola for his help as Director of the Electron Microscope Center at Southern Illinois University at Carbondale (SIUC), and the staff of the Electron Microscope Laboratory of the Materials Research Laboratory, University of Illinois, Champaign-Urbana. We also wish to thank K. Rankin, who diligently typed this manuscript.

157

CC Materials and Composites

References 1. Fitzer, Erich: The Future of Carbon-Carbon Composites. Carbon, vol. 25, no. 2, 1987, pp. 163-190. 2. Donnet, Jean-Baptiste; and Bansal, Roop Chand: Carbon Fibers. Dekker, Inc., c.1984.

Marcel

3. Jortner, Julius: Macroporosity and Interface Cracking in Multi-Directional Carbon-Carbons. Carbon, vol. 24, no. 5 , 1986, pp. 603-613. 4. McAllister, Lawrence E.; and Lachman, Walter L.: Multidirectional CarbonCarbon Composites. Fabrication of Composites, A. Kelly and s. T. Mileiko, eds., Elsevier Science Publ. Co., Inc., 1983, pp. 109-175.

5. Klein, A, J.: Which Weave To Weave. Adv. Mater. & Process., vol. 3, 1986, pp. 40-43. 6. Zhao, J. X.; Bradt, R. C.; and Walker, P. L., Jr.: Effect of Air Oxidation at 873 K on the Mechanical Properties of a Carbon-Carbon Composite. Carbon, vol. 23, no. 1, 1985, pp. 9-13. 7. Pickup, I. M.; McEnaney, B.; and Cooke, R. G.: Fracture Processes in Graphite and the Effects of Oxidation. Carbon, vol. 24, no. 5, 1986, pp. 535-543. 8. Hippo, E. J.; Murdie, N.; and Kowbel, W.: The Effect of Acid Treatments on Subsequent Reactivity of Carbon-Carbon Composites. Carbon, vol. 27, no. 3, 1989, pp. 331-336. 9. Hippo, E. J.; and Walker, P. L., Jr.: Reactivity of Heat-Treated Coals in Carbon Dioxide at 90OoC. Fuel, vol. 54, 1973, pp. 245-248. 10. Hippo, Edwin, J.; Murdie, Neil; and Hyjazie, Aamer: The Role of Active Sites in the Inhibition of Gas-Carbon Reactions. Carbon, vol. 27, no. 5 , 1989, pp. 689-695. 11. McKee, D. W.; Spiro, C. L.; and Lamby, E. J.: The Inhibition of Graphite Oxidation by Phosphorous Additives. Carbon, vol. 22, no. 3, 1984, pp. 285-290. 12. Reynolds, W. N.: Structure and Physical Properties of Carbon Fibers. Chemistry and Physics of Carbon, vol. 11, Marcel Dekker, Inc., c.1973, pp. 1-67. 13. McKee, D. W.; and Spiro, C. L.: The Effects of Chlorine Pretreatment on the Reactivity of Graphite in Air. Carbon, vol. 23, no. 4, 1985, pp. 437444.

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Carbon- Carbon Matrix Materials 14. Hirai, Toshio; Niihara, Koichi; and Goto, Takashi: Oxidation of CVD Si3N4 at 1500' to 1650'. J . American Ceram. SOC.,vol. 63, no. 7-8, July-Aug. 1980, pp. 419424. 15. Sheehan, J. E.: Ceramic Coatings for Carbon Materials. Proceedings of the Fourth Annual Materials Conference, M. Genisio, ed., Southern Illinois University at Carbondale, May 1987, pp. 56-58. 16. Ehrburger, P.; Baranne, P.; and Lahaye, J.: Inhibition of the Oxidation of Carbon-Carbon Composite by Boron Oxide. Carbon, vol. 24, no. 4, 1986, pp. 495499. 17. St. Pierre, G.: Basic Thermodynamic Considerations Related to Oxygen Protection of Structural Carbons. Proceedings of the Seventh Annual Materials Technology Center Conference (Composite-Technology), M. Genisio, ed., Southern Illinois University at Carbondale, Apr. 1991, pp. 181-220. 18. Strife, J.; and Sheehan, J.: Ceramic Coatings for Carbon-Carbon Composites. Ceram. Bull., vol. 67, no. 2, 1988, pp. 369-374. 19. Advanced Composites Design Guide, Volume II-Analysis. Third ed. Air Force Materials Lab., U.S. Air Force, Jan. 1973. (Available from DTIC as AD 916 680L.) 20. Johnson, D. J.: Carbon Fibers: Manufacture, Properties, Structure and Applications. Introduction to Carbon Science, H. Marsh, ed., Butterworths (London), 1989, p. 197. 21. Advanced composites Design Guide-Volumes I-V, Third ed. U.S. Air Force, Jan. 1973. (Available from DTIC as AD 916 679L to 916 683L.) 22. Schwartz, M. M.: Composite Materials Handbook. McGraw-Hill Book Co., c.1984, pp. 2.1-2.105. 23. Rand, B.: Carbon Fibres From Mesophase Pitch. Strong Fibres, W. Watt and B. V. Perov, eds., Elsevier Science Publ. Co., Inc., 1985, pp. 495-575. 24. Johnson, David J.: Structural Studies of PAN-Based Carbon Fibers. Chemistry and Physics of C a r b o e A Series of Advances, vol. 20, Peter A. Thrower, ed., Marcel Dekker, Inc., c.1987, pp. 1-58. 25. White, J. L.; and Buechler, M.: Microstructure Formation in Mesophase Carbon Fibers and Other Graphitic Materials. Petroleum Derived Carbons, J. D. Bacha, J. W. Newman, and J. L. White, eds., American Chemical SOC., 1986, pp. 62-83.

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CC Materials and Composites 26. Dresselhaus, Mildred S.; Dresselhaus, Gene; Sugihara, KO; Spain, Ian L.; and Goldberg, Harris A.: Graphite Fibers and Filaments. Springer-Verlag, 1988. 27. White, J. L.; and Sheaffer, P. M.: Pitch-Based Processing of Carbon-Carbon Composites. Carbon, vol. 27, no. 5 , 1989, pp. 697-707. 28. Marsh, H.; and Menendez, R.: Mechanisms of Formation of Isotropic and Anisotropic Carbons. Introduction to Carbon Science, H. Marsh, ed., Butterworths (London), 1989, p. 37. 29. Oya, A.; and Marsh, H.: Review-Phenomena of Catalytic Graphitization. J . Mater. Sci., vol. 17, no. 2, Feb. 1982, pp. 309-322. 30. Marsh, H.; and Kuo, K.: Kinetics and Catalyzing of Formation of Carbon Gasification. Introduction to Carbon Science, H. Marsh, ed, Butterworths (London), 1989, p. 108. 31. McKee, Douglas W.: Effect of Metallic Impurities on the Gasification of Graphite in Water Vapor and Hydrogen. Carbon, vol. 12, no. 4, Aug. 1974, pp. 453464. 32. Morgan, W. C.; and Thomas, M. T.: The Inverse Oxidation Phenomenon. Carbon, vol. 20, no. 1, 1982, pp. 71-78. 33. Rand, B.; Hosty, A. J.; and West, S.: Physical Properties of Pitch Relevant to the Fabrication of Carbon Materials. Introduction to Carbon Science, H. Marsh, ed., Butterworths (London), 1989, p. 75. 34. Bhatia, G.; Fitzer, E.; and Kompaliu, D.: Mesophase Formation in Tolvene Soluble and Insoluble Coaltar Pitch Fractions and Their Mixtures. Extended Abstracts-lnternational Carbon Conference (Bordeaux, France), 1984, pp. 330-331. 35. Marsh, Harry; and Walker, Philip L., Jr.: The Formation of Graphitizable Carbons Via Mesophase: Chemical and Kinetic Considerations. Chemistry and Physics of Curbor+-A Series of Advances, vol. 15, Philip L. Walker, Jr., and Peter A. Thrower, eds., Marcel Dekker, Inc., c.1979, pp. 229-286. 36. Forrest, M. A.; and Marsh, H.: The Effects of Pressure on the Carbonization of Pitch and Pitch/Carbon Fibre Composites. J. Mater. Sei., vol. 18, no. 4, Apr. 1983, pp. 978-990.

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Carbon- Carbon Matrix Materials 37. Curry, Donald M.; Scott, H. C.; and Webster, C. N.: Material Characteristics of Space Shuttle Reinforced Carbon-Carbon. The Enigma of the Eighties: Environment, Economics, Energy, Book 2, SOC.for the Advancement of Material and Process Engineering, 1979, pp. 1524-1539. 38. Schmidt, D. L.: Carbon-Carbon Composites. SAMPE J., May-June 1972, pp. 9-19. 39. Kotlensky, W. V.: A Review of CVD Carbon Infiltration of Porous Substrates. SAMPE J . 16, 1971, p. 257. 40. Kaae, J. L.: The Mechanism of the Deposition of Pyrolytic Carbons. Carbon, vol. 23, no. 6, 1985, pp. 665-673. 41. Jachlewski, T.; and Diefendorf, R. J.: Chemical Vapor Deposition of Carbon Matrices. Extended Abstracts-15th Biennial Conference on Carbon (Philadelphia, Pennsylvania), 1981, pp. 284-285. 42. Bokros, J. C.: Deposition, Structure and Properties of Pyrolytic Carbon. Chemistry and Physics of Carbon, Volume 5, P. L. Walker, Jr., ed., Marcel Dekker, Inc., 1969, pp. 1-118. 43. Oh, Seh-Min; and Lee, Jai-Young: Structures of Pyrolytic Carbon Matrices in Carbon/Carbon Composites. Carbon, vol. 26, no. 6, 1988, pp. 763-768. 44. Maahs, Howard G.; Ohlhorst, Craig W.; Barrett, David M.; Ransone, Philip 0.;and Sawyer, J. Wayne: Response of Carbon-Carbon Composites to Challenging Environments. Materials Stability and Environmental DegradatioVolume 125 of Materials Research Society Symposium Proceedings, A. Barkatt, E. D. Verink, Jr., and L. R. Smith, eds., Materials Research SOC., 1988, pp. 15-30. 45. Walker, P. L., Jr.; Rusinko, Frank, Jr.; and Austin, L. G.: Gas Reactions of Carbon. Advances in Catalysis and Related Subjects, Volume X l , D. D. Eley, P. W. Selwood, and Paul B. Weisz, eds., Academic Press Inc., 1959, pp. 133-221. 46. Jones, L. E.; Thrower, P. A.; and Walker, P. L., Jr.: Reactivity and Related Microstructures of 3D Carbon/Carbon Composites. Carbon, vol. 24, no. 1, 1986, pp. 51-59. 47. Knibbs, R. H.; and Morris, J. B.: Some Effects of Oxidation on the Properties of Graphite. Proceedings of 3rd International Conference on Carbon and Graphite, SOC.of Chemical Industry (London), 1971, pp. 297-308.

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C C Materials and Composites 48. Dickinson, Richard C.: Carbon/Carbon Composites: Fabrication and Properties and Selected Experiences. Materials Stability and Environmental Degradatio-Volume I 2 5 of Materials Research Society Symposium Proceedings, A. Barkatt, E. D. Verink, Jr., and L. R. Smith, eds., Materials Research SOC.,1988, pp. 3-13. 49. Donnet, J. B.: Structure and Reactivity of Carbons: From Carbon Black to Carbon Composites. Carbon, vol. 20, no. 4, 1982, pp. 267-282. 50. Ehrburger, P.; and Donnet, J. B.: Carbon Fibers in Polymer Reinforcement. Carbon, vol. 15, 1977, pp. 143-152. 51. Rakszawski, J. F.; and Parker, W. E.: The Effect of Group IIIA-VIA Elements and Their Oxides on Graphite Oxidation. Carbon, vol. 2, 1964/65, pp. 53-63. 52. Baker, R. T. K.; and Harris, P. S.: Controlled Atmosphere Electron Microscopy. J . Phys. E.: Sei. Instrum., vol. 5 , no. 8, Aug. 1972, pp. 793-797. 53. Hutcheson, J. M.: Polycrystalline Carbon and Graphite. Modern Aspects of Graphite Technology, L. C. F. Blackman, ed., Academic Press, Inc., 1970, pp. 1-47. 54. Lewis, J. B.: Thermal Gas Reactions of Graphite. Modern Aspects of Graphite Technology, L. C. F. Blackman, ed., Academic Press, Inc., 1970, pp. 129-190.

55. Goldstein, E. M.; Carter, E. W.; and Kluz, S.: The Improvement of the Oxidation Resistance of Graphite by Composite Technique. Carbon, vol. 4, no. 2, July 1966, pp. 273-279. 56. Sheehan, James E.: Oxidation Protection for Carbon Fiber Composites. Carbon, vol. 27, no. 5 , 1989, pp. 709-715. 57. Strife, J.; and Sheehan, J.: Ceramic Coatings for C-C Composites. Ceram. Bull., vol. 67, no. 2, 1988, pp. 364-374.

58. Grint, A.; Swietlik, U.; and Marsh, H.: Carbonization and Liquid Crystal (Mesophase) Development. Fuel, vol. 58, 1979, p. 642. 59. Forrest, M. A.; and Marsh, H.: Structure in Carbon/Carbon Fibre Composites as Studied by Microscopy and Etching With Chromic Acid. J . Mater. Sci., vol. 18, no. 4, Apr. 1983, pp. 973-977. 60. Knibbs, R. H.: The Use of Polarized Light Microscopy in Examining the Structure of Carbon Fibers. J . Microsc., vol. 94, 1971, pp. 273-281.

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61. Ehrburger, P.; Lahaye, J.; and Bourgeois, C.: Characterization of CarbonCarbon Composites-I. Carbon, vol. 19, no. 1, 1981, pp. 1-5. 62. Marsh, H.: Carbonization and Liquid-Crystal (Mesophase) Development-Part I: The Significance of the Mesophase During the Carbonization of Coking Coals. Fuel, vol. 52, 1973, pp. 205-212. 63. Atkinson, C.; and Marsh, H.: The Formation of Mesophase from a Range of Pitch Precursors. Extended Abstracts-Sixth International Carbon and Graphite Conference (London), 1982, p. 198. 64. Murdie, Neil; Edwards, Ian A. S.; and Marsh, Harry: Changes in Porosity of Graphite Caused by Radiolytic Gasification by Carbon Dioxide. Carbon, vol. 24, no. 3, 1986, pp. 267-275. 65. Mays, T. J.; and McEnaney, B.: A Quantitative Study of Pores and Cracks in a 3D Carbon Fiber Reinforced Carbon Composite. Extended Abstracts-1 8th Biennial Conference on Carbons, American Carbon SOC.,1983, p. 438. 66. Marsh, H.; and Smith, J.: Analytical Methods for Coal and Coal Products, Volume ZZ, Clarence Karr, Jr., ed., Academic Press, Inc., 1978. 67. Murdie, N.; Don, J.; Kocher, C.; Liew, R.; and Ju, C. P.: Microstructural Analysis of C-C Composites Used in Aircraft Braking Applications. The 13th Conference on Metal Matrix, Carbon, and Ceramic Matrix Composites, John D. Buckley, ed., NASA CP-3054, Part 1, 1990, pp. 349-365. 68. Forest, R. A.; and Marsh, H.: Reflection Interference Colours in Optical Microscopy of Carbon. Carbon, vol. 15, 1977, pp. 348-349. 69. Edwards, I. A. S.: Structure in Carbons and Carbon Forms. Introduction to Carbon Science, H. Marsh, ed., Butterworths (London), 1989, p. 1. 70. Warren, B. E.; and Bodenstein, P.: The Shape of Two-Dimensional Carbon Black Reflections. Acta Crystallogr., vol. 5, pt. 20, May 10, 1966, pp. 602605. 71. Johnson, D. J.; and Tyson, C. N.: The Fine Structure of Graphitized Fibres. British J . Appl. Phys. D , ser. 2, vol. 2, no. 6, June 1969, pp. 787-795. 72. Marsh, H.; and Forrest, M.: Structure of Matrix Carbons in C-C Fiber Composites. Extended Abstracts-15th Biennial Conference on Carbon, 1981, p. 270.

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CC Materials and Composites 73. Kowbel, W.; Hippo, E. J.; and Murdie, N.: Electron Microscopy Studies of Structural Heterogeneity in Carbon-Carbon Composites. Metal Matrix, Carbon, and Ceramic Matrix Composites-1986, John D. Buckley, ed., NASA CP-2482, 1987, pp. 267-281. 74. Wicks, B. J.; and Coyle, R. A.: Microstructural Inhomogeneity in Carbon Fibres. J . Mater. Sci., vol. 11, no. 2, Feb. 1976, pp. 376-383. 75. Bamet, F. Robert; and Norr, Marriner K.: Carbon Fiber Etching in an Oxygen Plasma. Carbon, vol. 11, no. 4, Aug. 1973, pp. 281-288. 76. Oberlin, A.: Carbonization and Graphitization. Carbon, vol. 22, no. 6, 1984, pp. 521-541. 77. Don, J.; and Ju, C. P.: Basal Plane Orientation of Polyacrylonitrile Fibers in a Commercial Polyacrylonitrile-Pitch Carbon-Carbon Composite. Mater. Sci. & Eng. A , vol. A124, no. 2, Apr. 1990, pp. 259-266. 78. Ju, C. P.; and Don, J.: Carbon-Carbon Thin Foil Preparation Techniques and Problems. Mater. Charact., vol. 24, 1990, pp. 77-82. 79. Marsh, H.; and Crawford, D.: Structure in Graphitizable Carbon From Coal-Tar Pitch HTT 750-1 148 K Studied Using High Resolution Electron Microscopy. Carbon, vol. 22, no. 415, 1984, pp. 413422. 80. Bennett, S. C.; and Johnson, D. J.: Electron-Microscope Studies of Structural Heterogeneity in PAN-Based Carbon Fibres. Carbon, vol. 17, no. 1, 1979, pp. 25-39. 81. Auguie, D.; Oberlin, M.; Oberlin, A.; and Hyvemat, P.: Microtexture of Mesophase Spheres as Studied by High Resolution Conventional Transmission Electron Microscopy (CTEM). Carbon, vol. 18, no. 5 , 1980, pp. 337-346. 82. Shiraishi, M.; Terriere, G.; and Oberlin, A.: Electron Microscopic Study on Graphitization of Bulk Mesophases. J . Mater. Sci., vol. 13, no. 4, Apr. 1978, pp. 702-710. 83. Crawford, D.; and Johnson, D. J.: High Resolution Electron Microscopy of High Modulus Carbon Fibers. J . Microsc., vol. 94, 1971, pp. 51-62. 84. Brooks, J. D.; and Taylor, G. H.: The Formation of Some Graphitizing Carbons. Chemistry and Physics of Carbo-A Series of Advances, Volume 4, Philip L. Walker, Jr., ed., Marcel Dekker, Inc., 1968, pp. 243-286.

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85. Markovic, V.; Lander, J. R.; Marsh, H.; and Taylor, D.: Kinetics of Gasification of Carbon of Different Optical Textures Using Oxygen and Carbon Dioxide. Extended Abstracts--15th Biennial Conference on Carbon (Philadelphia, Pennsylvania), 1981, pp. 415416. 86. Ragan, S.; and Marsh, H.: Fracture Mechanisms in Microstrength Testing of Carbon Artifacts. Carbon, vol. 18, no. 12, Dec. 1983, pp. 3712-3720. 87. Cranmer, J. H.; Plotzker, I. G.; Peebles, L. H., Jr.; and Uhlmann, D. R.: Carbon Mesophase-Substrate Interaction. Carbon, vol. 21, no. 3, 1983, pp. 201-207. 88. Meyer, R. A.; and Gyetvay: Carbon-Carbon Composites: Matrix Microstructure and Its Possible Influence on Physical Properties. Petroleum Derived Carbons, J. D. Bacha, J. W. Newman, and J. L. White, eds., American Chemical SOC.,1986, pp. 380-394. 89. Murdie, N.; Don, J.; Kowbel, W.; Shpik, P.; and Wright, M. A.: Effect of Fabrication Parameters on the Mechanical Properties of Various C/C Composites. Materials Processes-The Intercept Point, Volume 20 of International SAMPE Technical Conference Series, SOC. for the Advancement of Material and Process Engineering, 1988, pp. 73-82. 90. Zimmer, J. E.; and Weitz, R. L.: Disclinations and Fracture. Proceedings of the Fourth Annual Materials Conference, M. Genisio, ed., Southern Illinois University at Carbondale, May 1987, pp. 142-169. 91. Smith, R. Nelson; Young, David A.; and Smith, Roger A.: Infra-red Study of Carbon-Oxygen Surface Complexes. Trans. Faraday SOC.,vol. 62, 1966, pp. 2280-2286. 92. Zimmer, J. E.; and Weitz, R. L.: Magnetic Field Alignment of the Matrix in a C/C Composite. Carbon, vol. 26, 1988, pp. 579-588. 93. Zimmer, J. E.; and White, J. L.: Mesophase Alignment Within Carbon-Fiber Bundles. Carbon, vol. 21, no. 3, 1983, pp. 323-324. 94. Hishiyama, Y.; Inagaki, M.; Kimura, S.; and Yamada, S.: Graphitization of Carbon Fibre/Glassy Carbon Composites. Carbon, vol. 12, no. 3, June 1974, pp. 249-258. 95. Yasuda, Eiichi; Tanabe, Yasuhiro; Manocha, Lalit M.; and Kimura, Shiushichi: Matrix Modification by Graphite Powder Additives in Carbon Fiber/Carbon Composite With Thermosetting Resin Precursor as a Matrix. Carbon, vol. 26, no. 2, 1988, pp. 225-227.

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CC Materials and Composites 96. Ju, C. P.; Murdie, N.; Wright, M. A.; Don, J.; He, J.; and Fortunato, F. A.: Microstructure of Worn Pitch/Resin/CVI Carbon-Carbon BrakesTransmission Electron Microscopy. The 14th Conference on Metal Matrix, Carbon, and Ceramic Matrix Composites, John D. Buckley, ed., NASA CP3097, Part 2, 1990, pp. 497-506. 97. Perry, John L.; and Adams, Donald F.: An Experimental Study of CarbonCarbon Composite Materials. J . Mater. Sci., vol. 9, no. 11, Nov. 1974, pp. 1764-1774. 98. Fitzer, E.; and Huttner, W.: Structure and Strength of Carbon/Carbon Composites. J. Phys. D: Appl. Phys., vol. 14, no. 3, Mar. 14, 1981, pp. 347-371. 99. Evangelides, J. S.: Influence of Pyrolysis Pressure on Microstructure of CarbonCarbon Composites. 13th Biennial Conference on Carbon-Extended Abstracts and Program, American Carbon SOC.,1977, pp. 76-77. 100. White, Jack L.: The Formation of Microstructure in Graphitizable Carbons. SAMSO-TR-74-93, U.S. Air Force, Apr. 15, 1974. (Available from DTIC as AD 777 814.) 101. Zimmer, J. E.; and White, J. L.: The Core Structure of Disclinations. 14th Biennial Conference on Carbo-Extended Abstracts and Program, American Carbon SOC.,1979, pp. 429430. 102. Becker, Eric J.: NSWC Development Programs in Carbon-Carbon Composites. Metal Matrix, Carbon, and Ceramic Matrix Composites-1988, John D. Buckley, ed., NASA CP-3018, 1988, pp. 171-180. 103. Thomas, C. R.; and Walker, E. J.: Effects of PAN Carbon Fibre Surface in Carbon-Carbon Composites. Proceedings of 5th London International Carbon Graphite Conference, SOC.of Chemical Industry (London), 1978, pp. 520-531. 104. Fitzer, E.; Geigl, K.-H.; and Huttner, W.: The Influence of Carbon Fibre Surface Treatment on the Mechanical Properties of Carbon/Carbon Composites. Carbon, vol. 18, no. 4, 1980, pp. 265-270. 105. Manocha, Lalit M.; Yasuda, Eiichi; Tanabe, Yasuhiro; and Kimura, Shiushichi: Effect of Carbon Fiber Surface-Treatment on Mechanical Properties of C/C Composites. Carbon, vol. 26, no. 3, 1988, pp. 333-337. 106. Newling, D. L.; and Walker, E. J.: High-Performance ‘Graphitized’ CarbonCarbon Composites. International Conference on Carbon Fibres, Their Composites and Applications, Plastics Inst., 1971, paper 37.

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Carbon- Carbon Matrix Materials 107. Fitzer, E.; Geigl, K.-H; and Huttner, W.: Studies on Matrix Precursor Materials for Carbon Carbon Composites. Proceedings of 5th International Conference on Carbon and Graphite (London),Volume 1, 1978, pp. 493-506. 108. McAllister, L. E.; and Taverna, A. R.: A Study of Composition-Construction Variations in 3-D Carbon-Carbon Composites. Proceedings of International Conference on Composite Materials, Volume 1, Material SOC.of AIME, 1976, pp. 307-326. 109. Tsai, Stephen W.; and Hahn, H. Thomas: Introduction to Composite Materials. Technomic Publ. Co., Inc., 1980. 110. Walker, M. S.; and Lee, S. C.: Material and Process Effects on the Interlaminar Strength of Structural Carbon/Carbon Materials. Metal Matrix, Carbon, and Ceramic Matrix Composites-1987, John D. Buckley, ed., NASA CP-2482, 1987, pp. 327-342. 111. Ransone, P. 0.;Maahs, H. G.; Ohlhorst, C. W.; and Sawyer, J. W.: Interlaminar Tensile Testing of Thin Carbon-Carbon Composites. Proceedings of the Fifth Annual Materials Conference, Southern Illinois University at Carbondale, Apr. 1988. 112. Schmidt, Donald L.: Carbon/Carbon Composites. SAMPE J . , vol. 8, no. 3, May/June 1972, pp. 9-24. 113. Girard, H.: The Preparation of High Density Carbon-Carbon Composites. Proceedings of 5th International Conference on Carbon and Graphite, SOC. of Chemical Industry (London), 1973, pp. 483492.

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Chapter 6 Mechanics of Multidirectional Carbon-Carbon Composite Materials John J. Kibler Materials Sciences Corporation Blue Bell, Pennsylvania

Abstract

170

Symbols

170

Introduction

170

Background

172

Characteristics of CC Materials Description of Model Method of Analysis

175

176 179

Multidirectional Composite Model Degraded Properties Model Thin CC Composites

179

181

182

Property Predictions and Data Correlation Effects of Degraded Properties Concluding Remarks References

183

188

191

193

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CC Materials and Composites

Abstract This chapter describes the mechanics of analyzing carbon-carbon (CC) composite materials from the structural analyst’s point of view. The peculiarities of CC material as well as their influences upon the material behavior are discussed. Many assumptions must be made regarding some of the material parameters; an intelligent approach to making those assumptions is presented. Material models for evaluating thermoelastic properties of fiber composite materials having a spatial distribution of fiber orientations are discussed. The approach is to construct a material model with a finite number of fiber orientations in a repeating unit cell. This basic material model is a mini-mechanical model based upon the concept that a material system can be regarded as an assemblage of unit cells.

Symbols EA

modulus in axial fiber direction, Msi

ET

modulus in transverse fiber direction, Msi

E,

composite modulus in a fiber ( 2 ) direction, Msi

GA

fiber axial shear modulus, Msi

&/lA

thermal expansion in axial fiber direction

Al/lT

thermal expansion in transverse direction

VAT

Poisson’s ratio in axial direction

UTT

Poisson’s ratio in transverse direction

Introduction The successful use of CC materials in numerous applications has led the engineering community to view CC as a viable material for other applications. Because this material is expensive and not a simple off-the-shelf material, the type of composite which is best suited for any given application must be considered; that is, one cannot simply place an order for a given number of pounds of CC composite. As with most composites, the reinforcement type and complexity must be defined to determine which is best suited for the application. Then the user must define the environment and the complexity of the stress states which must be carried by the material. Given a definition of the desired material type, the user must work with the material manufacturers to develop the specific material and processing approach for the given application. The complexity of the material needed for a given application must be defined. For instance, if the loading is primarily in one plane and attachment loads will

170

Mechanics of Multidirectional Carbon- Carbon Composite Materials not be significant, then a two dimensionally reinforced material may be the right choice. However, three-dimensional reinforcement may be a better choice if the loading is complex or if the part must be relatively thick. Three-dimensional (3-D) materials, braids, and multiaxial weaves all provide reinforcement in more than two directions and therefore may be capable of carrying more complex loading conditions. The penalty for using a multidimensional weave is the increase in material costs and the reduction of in-plane properties in order to add fibers in the extra directions. Consequently, material selection is a major part of the development of new applications for these complex materials. These comments are true not only for CC materials but also for multidirectional composites in general. Therefore, the ability to synthesize material properties and to predict the performance (and, hence weaknesses) of a structure made of these materials is both necessary and important to the development of new materials. It should be obvious that the efficient use of materials and material development dollars requires that the material be designed for (at a minimum) each generic component. Analysis of these materials requires that the analyst understand the structure and the performance of the materials under given conditions. In this chapter, CC from an engineering mechanics point of view is addressed; descriptions of analytical models that are capable of modeling the material at the subcell level also are discussed. Even if the material to be used for a given application is defined, the structural analyst is faced with the question of how to analyze the structure using this material and how to determine its potential failure modes. The peculiarities of CC material which make it unique from an analyst’s point are discussed first. The influence of these characteristics on mechanical and thermal properties is noted, followed by approaches to modeling the material. Thermal protection materials for such high-performance applications as rocket nozzles and reentry vehicle nose tips can be fabricated with a wide range of constituent mechanical and thermal properties and with various internal geometries utilizing three-dimensional arrays of fibers. The performance of the resulting material will be a function of such variables as choice of fiber (or fibers, in combination), weave parameters, type of matrix material, and the temperature, pressure history, and number of cycles experienced during the fabrication process. Because of the large number of variables and the high unit cost for test specimens (particularly when the material under consideration is of the CC class), direct evaluation of the relative performance capability of candidate materials is expected to be a time-consuming and costly approach. The alternative approach is the utilization of reliable material guidance models.

171

CC Materials and Composites The mini-mechanics approach to modeling composite materials is a balance between the micro- and macro-mechanics approaches. A macro-mechanics material model treats a composite as a homogeneous but anisotropic material. Using this effective material, the heterogeneity of the composite material is ignored and, consequently, realistic (microlevel) failure mechanisms are difficult to be treated as part of the design failure criterion. A macro-mechanics approach cannot be utilized as a predictive tool for the thermomechanical properties of materials not yet fabricated or tested, and this approach is of limited assistance in the development of new materials. The macro-mechanics approach, however, does permit structural design analyses to be performed using state-of-the-art analytical techniques. At the other extreme, a micro-mechanics model treats each distinct subregion of the material in an attempt to find a detailed solution for the local stresses and strains. With this approach, failure prediction becomes dependent upon details of the internal geometry such as fiber placement within a fiber bundle. This results in a complex methodology which is sensitive to local variables that are not readily defined by the material construction parameters. However, micro-mechanics solutions for specific regions within the unit cell can provide an important function in the construction of a material model. The mini-mechanics model developments described herein were incorporated into computer codes for treatment of properties of CC materials reinforced by fibers in three or more spatial directions. Comparisons between the predicted and experimental results demonstrate excellent agreement for elastic constants and thermal expansion coefficients. A combination of micro- and mini-mechanics approaches is used at Materials Sciences Corporation (MSC) for modeling these materials. In particular, a model for fiber bundle and matrix properties was incorporated into a laminate analysis model for predicting properties. The success of this approach is discussed in a following section.

Background A large number of alternative methods to define the effective properties of composite structural materials exist; however, the work in the area of threedimensional composites is quite limited. Certain methods for structural composites can play a role in material modeling for the more complex composites. These composites may be defined as quasiregular arrays of yarns (characteristically carbon or graphite) oriented in three or more nonplanar directions and embedded within a matrix of phenolic, carbon, or graphite. There is a regularity to the internal geometry; however, numerous geometric imperfections do exist. Nonuniformity is found in filament spacing within the impregnated yarn bundle and in voids and microcracks within the matrix.

172

Mechanics of Multidirectional Curbon-Curbon Composite Muteriuls The direct approach to define effective properties of this material is to describe quantitatively the geometry, to obtain an exact solid mechanics boundary value solution, and to integrate the state variables to obtain the constraints that relate the average values of these state variables. However, to define every geometrical detail of filament cross section and spacing geometries in a deterministic fashion is clearly a task of undesirable complexity. Even if such a feat could be performed for the specific configuration, the actual geometry will vary from yarn bundle to yarn bundle and also along any single yam bundle. Thus, the problem is to find a realistic yet practical alternative to the exact solution. Several alternatives to obtaining exact solid mechanics boundary value solutions to the exact geometry are available. These alternatives are presented schematically in figure 1.

Exact solution for actual fiber bundlelunit cell geometry

Exact solutions for very special geometry cases

Exact solutions

Cylindrical fibers in uniform packing configurations

,

Approximate solutions to approximate geometries

Exact solutions to approximate geometries

/

/ '

Closed-form solutions for special cases

Finiteelement I solutions 1

1

Composite cylinder assemblage solutions

I

Represent geometry with statistical models

Approximations to actual geometry

,I I

\

Mechanics of materials

f!kB!

1" I

Self-consistent

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173

CC Materials and Composites relations between the thermal expansion coefficients of a two-phase material and other known properties, including effective elastic moduli of the composite material and constituent properties (see also refs. 3 and 4). Approximate geometries appear to be the most fruitful approach to modeling CC composites. Several models have been used to predict the behavior of approximate composite geometries. These analytical models include exact solutions to boundary value problems, rigorous bounding approaches (whereby effective properties can be bounded from above and below), and approximate solutions such as first-order strength-of-materials analyses or the self-consistent scheme. The most important exact solution to approximate geometries is represented by several finite-element solutions (e.g., ref. 5). These results are for regular geometrical arrays (rectangular or hexagonal) of unidirectional composites (figs. 2(a) and (b)) and the composite cylinder assemblage (ref. 6 and fig. 2(c)). Both geometrical idealizations have been used extensively and successfully by many researchers (e.g., refs. 6-9). Solutions to both of these approximate geometries have been demonstrated to give results that are close to each other and to actual experimental data. Also, both approaches can be successfully applied to the modeling of CC materials. A more detailed discussion of the different methods in use for the evaluation of elastic constants of unidirectional fiber composites can be found in reference 10.

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For multidirectional composites, the literature is more limited. A comparison of theoretical models of CC materials has been performed by Jerina (ref. 11). Various approaches were compared with data on a laboratory material system, which was a nylon-fiber-reinforced rubber. Because this field is relatively new, the comparison does not treat many of the state-of-the-art models that have been recently developed. This review did treat the commonly used constant strain and constant stress models for composite properties, which are the Voigt (ref. 12) and Reuss (ref. 13) models. Paul (ref. 14) showed that these models are actually upper and lower bounds on the actual stiffnesses. Jerina showed that an approximate model developed by Pagano (ref. 15), based on a laminate approach, yielded good agreement for many composite elastic properties.

174

Mechanics of Multidirectional Carbon- Carbon Composite Materials Among the other approaches to unit cell modeling, which were not treated in the review of reference 11, are the mechanics of materials approaches of Greszczuk (e.g., ref. 16) and of Zimmer et al. (ref. 17), in which stiffnesses in series and parallel are used as models. More complex models, using combinations of closedform approximate solutions and finite-element models, were formulated by Crose (ref. 18). The model described in this paper resulted from a series of developments by Materials Sciences Corporation over a number of years. The general methodology for evaluating unit cell properties stems from an approach method originally presented by Dow and Rosen (ref. 19). The approach utilized geometry models that are extensions of those presented in references 20 and 21. In a series of unpublished company papers, the authors and their colleagues at the Materials Sciences Corporation developed a methodology to treat multidirectional, threedimensional composite materials reinforced in any number of spatial directions. In recent developments, the effects of localized regions of high microcracking were incorporated via the concept of a degraded property region. These regions of high microcracking are of particular importance to CC materials, which develop severe microcracking between constituents during the fabrication process. These developments have been reported in reference 22.

Characteristics of CC Materials One of the principal difficulties in predicting properties of CC materials is in defining the constituents to be used as input to the models. This difficulty is a problem because the constituents that one starts with are modified by the high temperatures during processing. Consequently, the process history becomes a major influence upon the final in situ constituent properties. Previous chapters described processing approaches and indicated how processing can affect the constituents. The analyst must use previous experience or similar processed material properties to define a starting point for constituents. The following points should be considered when beginning to define input properties:

1. High graphitization temperatures plus tensile stresses during processing will generally increase the in situ fiber modulus and lower the fiber expansion characteristics. These properties are highly dependent upon the fiber type used in the material.

175

CC Materials and Composites 2. First cycle pitch impregnation will promote high alignment of the matrix material surrounding the fibers, leading to an increasingly effective composite modulus.

3. Pyrolytic graphite will result locally in very anisotropic matrix properties, which can influence the matrix-dominated composite properties.

4. Cooldown from graphitization cycles will result in cracks developing between subcell regions, thus reducing the composite shear modulus and expansion coefficients. Whether a final graphitization cycle is or is not used will have a large effect on the properties during the first subsequent heat-up of the material. If multiple temperature cycles must be sustained by the composite, then final graphitization should be taken into account when deciding on the final processing conditions.

Description of Model The basic mini-mechanical material model proposed for multidirectional reinforced composite materials is based upon the concept that a composite material system can be regarded as an assemblage of unit cells. This fact is illustrated in figure 3. The basic unit cell is defined by fiber and matrix properties and the phase geometry. The effective property analysis is performed on the representative volume element (RVE) which contains a large number of unit cells. The material variability can be incorporated by treating an assemblage of RVE’s having some statistical variability of, for example, fiber volume fraction. The unit cell models are constructed from impregnated yarn bundles oriented in space and from interstitial matrix regions as shown schematically in figure 3 . Several basic versions of this unit cell model have been analyzed. To be effective, a model of the material must include the following material characteristics:

1. Anisotropic constituents, including the effects of oriented matrix material such as in a fiber “sheath” 2. Definition of fiber content within each fiber bundle, as well as definition of total fiber content within a repeating element

3. Definition of orientation of each fiber bundle within the repeating element 4. Definition of the interface properties between subcell regions to model the subcell cracking that exists in most CC materials

5. Dispersed voids, of both random and elongated shapes 6 . Temperature-dependent constituent properties

176

Mechanics of Multidirectional Carbon- Carbon Composite Materials

Figure 3. Section of 3-0 orthogonally reinforced material showing assembly of unit cells. The MSC NDPROP (N-Directional Composite Analysis Program) model has been developed for the analysis of very general composite unit cells. The unit cell is composed of fiber bundles oriented in space by defining direction numbers that determine the orientation of each fiber bundle. The cell can contain any number of bundles, provided that the total bundle and interstitial matrix volume fractions do not exceed unity. The material does not need to be orthotropic, but can be fully anisotropic, if necessary. Each of the fiber materials can be different, thus allowing treatment of hybrid composites. The constituent fiber and matrix materials may have transversely isotropic elastic constants and thermal expansion coefficients. Material porosity can be treated and matrix properties in each fiber bundle and in the interstitial matrix regions can differ. The deficiency in the NDPROP model is the lack of ability to model explicitly the interface regions between bundles. Rather, an effective matrix is defined which includes the effects of cracks between bundles. The DCAP (Directional Composite Analysis Program) model, which was developed under U.S. Navy funding and is available to qualified government contractors through the Naval Surface Weapons Center, focuses on the effects of

177

CC Materials and Composites the significant localized microcracking that occurs in CC composites. This model considers a unit cell that is reinforced with three locally orthogonal fiber bundles and which contains regions of material with variable stiffness between the unit cell subregions. Actual crack geometry and crack-tip stress intensity effects are not considered. Rather, the transmission of loads across the interfaces in the unit cell is considered. This approach to the effect of microcracking leads to a definition of the efficiency of the unit cell interfaces in transmitting loads. The unit cell efficiency can be considered equivalently as either the transmission of a load, or as a measure of the uncracked contact area along the interfaces within the unit cell (see fig. 4).

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Figure 4. 3-0 CC unit cell model: (a) unit cell geometry and (b) weak regions or surfaces of weakness. The drawback with the DCAP model is that it is developed especially for three directionally reinforced materials and cannot handle multidirectional materials or thin two-dimensional material easily. Multidirectionally reinforced materials require an analysis approach similar to NDPROP, which assumes an effective matrix material. Thin-section two-dimensional laminated-type materials can be analyzed with a laminate analysis code, as described in the following sections.

178

Mechanics of Multidirectional Carbon- Carbon Composite Materials

Method of Analysis The basic procedure used in both material models is the application of the stress state associated with the particular elastic constant of interest to the unit cell. This stress state results in an average strain state defined by the stresses and the unknown effective elastic constant. In a homogeneous material, this strain state would define the actual displacements throughout the material. In the composite, those displacements are taken as the displacements existing on each of the internal surfaces separating one subcell region from another. In this manner, the displacements on the entire boundary surface of each subcell region are defined. Such displacements define average strain components with respect to the principal elastic axes of each subcell region. If the elastic constants of each subcell are known (the method by which they are found is discussed subsequently), the average stresses in each subcell can be found. By proper coordinate transformations, the average stresses over the entire unit cell can be found from these subcell stresses. The unknown elastic constant is now found by equating the average stress state computed in this fashion to the applied stress state. This procedure is equivalent to using an admissible displacement field to obtain a bound on the strain energy, from which a subsequent bound on the elastic constant can be obtained. Details of the analysis are presented in references 22 and 23.

Multidirectional Composite Model The unit cell for these materials is regarded as an assemblage of subcells (see fig. 2), each of which contains either a unidirectional fiber composite or a matrix material (with or without dispersed voids). Utilization of the model requires definition of the effective thermoelastic properties of each subcell region as a function of the constituent properties. When voids are present in the matrix, the effective matrix properties are computed for various combinations of spherical and cylindrical voids as appropriate. A composite sphere or a composite cylinder assemblage of voids (refs. 6 and 10) is assumed in the matrix material. The voids in a fiber bundle matrix are probably highly elongated in the fiber direction. Effective moduli for such a matrix are computed from the composite cylinder assemblage results (ref. 6) by assuming that the fibers have zero stiffness. Even an isotropic matrix will become transversely isotropic when cylindrical voids are added. The interstitial matrix is likely to have voids that are approximately spherical. For randomly dispersed spherical voids, the lower bound given in reference 10 has been utilized. This bound results in the matrix modulus being reduced by a factor that is a function only of the void volume fraction. As shown in reference 22, even small amounts of voids have a significant effect on the matrix properties.

179

CC Materials and Composites Given the set of effective matrix properties, the one-dimensional (1-D) impregnated fiber bundle properties (shown as bundles 1, 2, or 3 in fig. 4) are computed using the composite cylinder assemblage results for composite cylinders with transversely isotropic constituents (ref. 10) in the form utilized in reference 22. In the composite cylinder assemblage model, a unidirectionally reinforced material is modeled by an assemblage of composite cylinders of variable sizes, which fill out space. Each composite cylinder consists of a fiber and concentric matrix shell. In an actual fiber-reinforced material, the fibers are more or less identical in crosssectional area and are randomly placed. Finite-element models utilizing circular fibers with equal diameters have been used to obtain exact solutions for regular fiber arrays. The composite cylinder assemblage model treats the geometric randomness that exists in real composites, but requires the fibers to have variable diameters. The primary advantages of this approach over the regular array model stem from the fact that it can be analyzed analytically, with resulting simple closed-form expressions for detailed stresses, strains, and composite properties. This analysis is important for both researcher and designer because it allows ready observation of trends due to changes in constituent properties and volume fractions. Furthermore, the simplicity of the model makes it possible to analyze not just elastic properties, but also thermal, thermoelastic, and viscoelastic properties. The model is also readily adaptable to the important case of fiber and matrix cylindrical anisotropy. A drawback of the model is that one elastic parameter (the transverse shear modulus) cannot be calculated exactly and can be bounded only from above and below. However, comparison with experiments shows that for the usual stiff fibers, the upper bound agrees well with the experiments and can be used as an approximate result. A comparison of finite-element analyses of regular arrays and composite cylinder assemblage equations shows that property and stress results are very close. In fact, hexagonal array results and composite cylinder assemblage results are practically indistinguishable. Both geometrical idealizations have been used extensively and successfully by many researchers (e.g., refs. 2 1-24). Solutions to both of these approximate geometries have been demonstrated to give results that are close to each other and to actual experimental data. Also, both approaches can be successfully applied to modeling of CC materials. The greater flexibility of the composite cylinder assemblage model and its ability to achieve significantly lower cost results are important factors in choosing it to model unidirectional composites.

Once the matrix and impregnated fiber bundle properties are available, the unit cell can be constructed as in figure 4. Both the NDPROP and DCAP computer programs allow a large amount of freedom in constructing the unit cell analysis. This feature allows the user to supply one set of constituent properties to the program and to construct a very wide range of unit cell configurations, as desired.

180

Mechanics of Multidirectional Carbon-Carbon Composite Materials The unit cell is constructed by specifying the fiber bundles in each direction and in the interstitial matrix material, along with volume fractions of each material. In the case of NDPROP, the bundle direction numbers define the orientation of each bundle relative to the global axes. Any number of bundles can be defined, each with its own fiber, matrix, and orientation, provided that the total volume fraction of all fiber bundles is less than unity.

Degraded Properties Model Cracking between subregions within the unit cell occurs as a result of cooldown from the graphitization cycle during fabrication of CC composites. The fibers in current CC materials have a much lower axial thermal expansion coefficient than the matrix material. Because the material is graphitized at high temperature, the material becomes stress free at that temperature. Upon cooldown, the matrix is loaded in tension since it tries to contract at a much higher rate than do the fibers. This loading results in cracks occurring between the fiber bundles and the matrix material. The degraded properties model was developed specifically to include the effects of cooldown cracking upon resulting composite properties. Actual crack geometry and crack-tip stress intensity effects have not been considered. Rather, the model considers the ability of the material to transmit loads across the interface between subregions within the unit cell. The weak regions are approximated as thin plate-like regions in which the thickness of such regions is assumed to be small in comparison to the dimensions of the subregions and the unit cell. In order to develop a simple analytical model that yields a suitable representative volume element, all parallel planes of weakness are assumed to have similar stress transfer characteristics, although each of the unit cell interconnecting surfaces may have its own unique stress transfer characteristics defined, if desired. The inclusion of regions of degraded load transfer capability in DCAP requires some definition of the effectiveness (efficiency) of the degraded material layers in transmitting axial and shear loads between subregions. The efficiency of the degraded material may be considered equivalently as the stress transfer capability (i.e., an approximate measure of the uncracked contact area between subregions) or as a measure of the stiffness of the degraded material. The analysis of the effect of these added regions upon unit cell properties is described in reference 22. Both the degraded property model (DCAP) and the basic multidirectional composite model (NDPROP) treat thermomechanical properties including the effects of temperature-dependent constituent properties. Thermal expansion data are defined for the constituents in the form of free thermal expansion versus temperature. The model then can compute secant thermal expansion coefficients for each subcell of the unit cell from the analysis temperature to the reference stress-free temperature. The secant thermal expansion coefficients for the subcells

181

CC Materials and Composites are then combined to form the composite thermal expansion coefficients. The resulting values for the composite, from the stress-free temperature to the given temperature, are used to obtain a set of predicted composite thermal expansion coefficients versus temperature. The output of the models provides all of the thermoelastic properties of the unit cells for each desired temperature. In addition, if composite stresses are applied, the unit cell strains and average strains and stresses in each subregion can be computed for each temperature. An example of property prediction and correlations with experimental data follows.

Thin CC Composites Historically, the development of CC composites has focused on multidirectionally reinforced materials. Focusing was a result of material being used for basically nonstructural applications in which heat resistance was the primary material function. These components were thick; consequently, the poor shear and throughthe-thickness direction properties required additional reinforcement directions. Future CC applications will continue to require multidirectionally reinforced materials, but there is also a growing interest in thin, laminate construction, CC composites for structural applications requiring light weight and good hightemperature behavior. For these thin-section CC materials, one asks how the properties might be predicted. The author and Materials Sciences Corporation have investigated this area and found that standard laminated plate analysis codes are adequate for predicting the behavior of thin-section CC composites. The code utilized for thin-section CC materials can be simply a standard laminate analysis code. The author has been utilizing CLASS, which is a personal computer based laminate code. The advantage of this code is that the composite cylinder assemblage model is built into the code so that the user can predict layer properties for a variety of fibers and matrix materials and then can construct the laminates from these layers. In addition, the code has the ability to combine particles such as needles, spheres, or platelets in a matrix by using upper and lower bound solutions as well as the differential scheme approach. This code allows creating matrices with various inhibitors or short fiber reinforcements and then adding long fibers to create a layer. Also, this code permits transversely isotropic constituents, all of which are temperature dependent. All of the constituent and layer properties are stored within a small data base for later access by the code. Thus, properties can be entered once and used over again. The basic analyses of thin CC material are identical to those of standard epoxy laminates. The difference is in the difficulty of defining the constituent properties

182

Mechanics of Multidirectional Carbon-Carbon Composite Materials and degree of sublayer cracking. As described next, if data can be obtained on one material and that material used to define the constituent properties through data correlations, then the properties of other laminated constructions can be predicted with excellent accuracy. The only caveats are that the new material must contain the same fibers and that it must be processed in a very similar fashion to that of the basic material. These limitations, however, should not be considered major drawbacks because as a material is developed, standard lay-ups are generally employed to characterize that process. Alternate lay-ups then can be employed with the optimized processing conditions. In this case, the analytical prediction of composite properties can be the most useful method to use. Material development can take place on a basic material construction, and tests from material taken from optimized processing conditions can be used to tune the input constituent properties. The code then can be utilized to predict the properties of other constructions and to define the material lay-up best suited for the given application. Finally, the optimized construction is fabricated and tested to verify the analytical modeling assumptions.

Property Predictions and Data Correlation Both models just described have been utilized to predict composite properties of a broad class of materials. The basic model has been used to define unique characteristics of materials having four or more fiber orientation directions. The degraded property model has been used to predict the behavior of CC materials. Material modeling of CC composites is a complex task. The choice of the appropriate analytical code for the prediction of material properties is only an initial step in the material modeling process. Definition of the appropriate modeling parameters must be determined so that the best material representation is created. The required modeling parameters are analytical code dependent and may include such things as fiber and matrix properties, degree of matrix cracking, and fiber bundle sheath content. Definition of these parameters is most appropriately obtained through extensive data correlations. Prior to initiating a data correlation study, a thorough understanding of the history of the material is necessary. Not only ‘is a definition of the weave architecture and types of constituents that comprise the material required, but also a concise understanding of the process history must be known as well. Fibers utilized in fabrication undergo property changes due to the high-temperature heating cycles occumng in the CC process (fig. 5). These property changes are dependent on the type of fiber and they vary based upon the maximum process temperature. For example, the process temperature may result in an increased stiffness for the fiber while degrading its strength (ref. 25). A typical set of fiber properties which was developed through data correlation studies is shown in table 1 for T-300 fibers.

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Mechanics of Multidirectional Carbon-Carbon Composite Materials

The effective properties of the matrix material are very much affected by the process procedure. These effects can be observed in figure 6. Figure 6 shows the measured composite thermal expansion response for two 2-D CC materials that have been subjected to different densification procedures. The densification process causes the impregnant to undergo various phase changes. The process history is also important in characterizing other factors. The processing of CC materials may result in a final composite that can contain a significant amount of matrix microcracking. This microcracking results from cooling the densified material to room temperature. The degree of cracks in the material is related to the final process temperature of the composite. Matrix microcracking affects the matrix-dominated properties of the composite, which include shear and thermal expansion. The use of thermal expansion and shear data can help quantify the degree of cracking in the material. For example, by utilizing the DCAP code, an estimate of the amount of matrix microcracking can be accomplished by correlating thermal expansion data with values of the DCAP unit cell efficiency parameter S . This influence of S is shown in figure 7. Another important factor is the method of impregnation. For example, an initial pitch impregnation cycle will cause a sheath of highly orientated matrix material to form around the fiber bundles. This sheath of high modulus matrix material acts to increase the stiffness of the bundles. The use of fiber direction composite tensile modulus measurements helps define the amount of sheath present. The effectiveness of the fiber bundle sheath is shown in figure 8. A primary requirement for the accurate modeling of CC materials is the determination of fiber and matrix properties. Associating the correct representation of fiber and matrix properties with particular process methods is vital in the modeling of these complex materials; therefore, it is important that the processing history of the material be factored into material property modeling and that the appropriate material representation be formulated. The verification of the fiber and matrix properties can be accomplished through correlation of experimental data with predicted properties. Through the use of data correlation, an appropriate material model can be defined for use in the efficient design of these materials. Differences between the predicted and measured behavior of materials are used to improve the understanding of both the constituents and the interactions between the components of the composite. The influence of the matrix material on composite properties is observed in figure 9. Here, the effect of matrix modulus is shown to influence the matrixdominated properties (i.e., shear and thermal expansion) of a 2-D CC material. Through the use of experimental data, correlations can be performed to determine

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The efficiencies are an approximate measure of the reduction of the composite Young’s moduli from the value of the modulus when no weak regions are present. If one assumes that the weak regions contain flat disk-shaped (crack-like) stressfree regions, then the efficiency is also an approximate measure of the fraction of the contact area between the subregions that are considered perfect in transferring stress. Therefore, when Sij = 0, all the weak surfaces perpendicular to the coordinate direction are fully cracked. When S.. = 1,no cracks exist in the weak ”. surfaces perpendicular to direction i. The efficiencies are similarly approximate measures of the reduction of shear modulus Gij or the fraction of the contact areas effective in transferring shear stress rij. Details of the analytical formulations and assumed displacements within the unit cell have been given in reference 23. The upper- and lower-bound formulations were also shown to give very close results for most weave configurations; because of this, the upper-bound results are sufficient for most purposes.

Concluding Remarks Both the basic model (NDPROP) and the degraded properties model (DCAP) have been shown to be capable of accurately predicting carbon-carbon (CC) composite material properties. DCAP, in particular, has been demonstrated to

191

CC Materials and Composites have the ability to model various CC materials with enough accuracy to permit reliable engineering studies of material performance with little or no experimental data. When developing analytical models for CC composites, the philosophy should be to choose an appropriate approach for each of the specific subproblems at hand. It must be emphasized that the objective of analytical modeling is to provide a cost-effective procedure for enhancing the efficiency of the materials development process. Analytical guidance of innovative material concepts will require extensive analytical screening of potential materials. Thus, the material model must be an engineering tool that treats many complex material problems in a manner that recognizes the primary objective.

In the course of developing these models, various methods have been used, including composite cylinder assemblages, variational methods, and approximate solutions suited to the problem. The name materials engineering has been coined to describe this engineering approach. The materials engineering approach to the development of a material model for quantitative material synthesis utilizes a modular concept that provides a convenient capability to upgrade the quality of predictions without substantial overhaul of the total material model. Also, during the developmental stage when certain modules of the code are approximate or preliminary in nature, the code may be used for comparative rather than for quantitative predictions and/or for data enhancement. The DCAP model has been used to demonstrate the feasibility of the concept of pre- and post-processor interaction with material performance codes. The DCAP code has been written so that it can interact with such thermostructural finiteelement codes. The interaction between codes is effective; yet, the interface with the DCAP code remains accessible such that the program can be made to interface with other performance codes. Material models of these types should prove valuable to the materials manufacturer in developing current and future material systems for advanced applications. Also, the models will provide the structural analysis community with an effective material property prediction tool to be used as a preprocessor to aerospace structural analysis codes. Carbon-carbon materials are complex and possess substantial variability in properties. This complexity coupled with variability should not be looked upon as an impossible task for analytical modeling. Rather, the analysis models should be viewed as tools that can help explain the trends found in data and guide the material development process so that the next batch of material will possess those characteristics that are desirable for the structure being designed.

192

Mechanics of Multidirectional Carbon- Carbon Composite Materials Once material properties for a given material construction are available, whether through testing, prediction, or a combination of both, the design analysis proceeds in a normal fashion. The only uniqueness of CC in this regard is its severe planes of weakness when matrix dominated properties control the failure modes. The designer needs to be aware of these planes of weakness so that those stress components are examined carefully and so that proof tests are designed that will demonstrate the failure mode.

References 1. Hill, R.: Theory of Mechanical Properties of Fibre-Strengthened Materials-I. Elastic Behaviour. J . Mech. & Phys. Solids, vol. 12, 1964, pp. 199-212. 2. Levin, V. M.: Thermal Expansion Coefficients of Heterogeneous Materials. Mekhanika Tverd. Tela, vol. 2, no. 1, 1967, pp. 88-94. (English translation in Mech. Solids, vol. 2, no. 1, 1967, pp. 5 8 4 1 . ) 3. Rosen, B. W.: Thermomechanical Properties of Fibrous Composites. Proc. Royal SOC.London A , vol. 319, no. 1536, Oct. 6, 1970, pp. 79-94. 4. Rosen, B. Walter; and Hashin, Zvi: Effective Thermal Expansion Coefficients and Specific Heats of Composite Materials. Zrtt. J . Eng. Sci., vol. 8, no. 2, Feb. 1970, pp. 157-173.

5. Picket, G.: Elastic Moduli of Fiber Reinforced Plastic Composites. Proceedings of Fundamental Aspects of Fiber Reinforced Plastic Composites, R. T. Schwartz and H. S. Schwarta, eds., Wright-Patterson Air Force Base, Jan. 1967, pp. 1330. 6. Hashin, Zvi; and Rosen, B. Walter: The Elastic Moduli of Fiber-Reinforced Materials. Trans. ASME, Ser. E: J . Appl. Mech., vol. 31, no. 2, June 1964, pp. 223-232. 7. Tsai, Stephen W.: Structural Behavior of Composite Materials. NASA CR-71, 1964. 8. Adams, Donald F.; Doner, Douglas R.; and Thomas, Rodney L.: Mechanical Behavior of Fiber-Reinforced Composite Materials. AFML-TR-67-96, U.S. Air Force, 1967. (Available from DTIC as AD 654 056.) 9. Chen, C. H.; and Cheng, Shun: Mechanical Properties of Fiber Reinforced Composites. J . Compos. Mater., vol. 1, 1967, pp. 30-41.

193

C C Materials and Composites 10. Hashin, Zvi: Theory of Fiber Reinforced Materials. NASA CR-1974, 1972. 11. Jerina, K. L.: Effective Moduli of Three Dimensionally Reinforced Fibrous Materials. Ph.D. Diss., Washington Univ. Sever Inst. of Technology, May 1974. 12. Voigt, W.: Uber die Beziehung Zwischen den Beiden Elastieitatsconstanten Isotroper Korper. Ann. Phys. (Leipzig), vol. 38, 1889, p. 573. 13. Reuss, A.: Berechnung der Fliessgrenze von Mischkristallen auf Grund der Plastizitats-bedingung fur Einkristalle. Z . Angew. Math. & Mech., vol. 9, 1929, p. 49. 14. Paul, B.: Predictions of Elastic Constants of Multiphase Materials. Trans. Metall. Soc. ofAIME, vol. 218, Feb. 1960, pp. 3 6 4 1 . 15. Pagano, N. J.: Exact Moduli of Anisotropic Laminates. Micromechanics, G. P. Sendeckyj, ed., Academic Press, Inc., 1972.

16. Greszczuk, Lcngin B.: Mechanics of Failure of Composites. MDC G5365 (Contract N00019-73C-0405), McDonnell Douglas Astronautics Co., May 1974. (Available from DTIC as AD 708 233.) 17. Zimmer, J. E.; White, J. L.; Evangelides, J. S.; and Meyer, R. A.: Carbon NoseTip Materials Technology. Volume 1: Microstructure and Fracture of Carbon Systems. TOR-0076 (6726-02)-2 VOL. 1 (Contract No. F04701-75-C-0076), Aerospace Corp., Sept. 1975. (Available from DTIC as AD BO07 379L.) 18. Crose, James G.: Carbon-Carbon Nosetip Program (CCNP). PDA Rep. No. 1021-00-04 (Contract No. N60921-74-C-0158), Prototype Development Assoc., Inc., July 30, 1975. (Available from DTIC as AD BO07 054L.) 19. Dow, Norris F.; and Rosen, B. Walter: Zero Thermal Expansion Composites of High Strength and Stiffness. NASA CR-1324, 1969. 20. Rosen, B. Walter; and Shu, Larry S.: On Some Symmetry Conditions for Three Dimensional Fibrous Composites. J . Compos. Mater., vol. 5, Apr. 1971, pp. 279-282. 21. Brazel, James P., ed.: Advanced Hardened Antenna Window Materials Study. AMMRC CTR 72-1, U.S. Army, Feb. 1972. (Available from DTIC as AD 741 384.)

194

Mechanics of Multidirectional Carbon- Carbon Composite Materials 22. Kibler, J. J.; and Chatterjee, S. N.: Development of a Minimechanics Model for 3 - 0 CarbonlCarbon Materials. TFR/75 10, U.S. Navy, July 1975. (Available from DTIC as AD BOO8 415L.) 23. Chatterjee, S. N.; and McLaughlin, P. V.: Strength and Thermoelastic Properties of 3 - 0 CarbonlCarbon Composites. TFR 5 13, Materials Sciences Corp., Oct. 1975. 24. Smith, Robert E.: Ultrasonic Elastic Constants of Carbon Fibers and Their Composites. J . Appl. Phys., vol. 43, no. 6, June 1972, pp. 2555-2565. 25. Schultz, D. A.: Carbon Fiber Property Changes Resulting From Simulated Carbodcarbon Composite Processing. Proceedings of 18th Biannual Conference on Carbon, Worcester Polytechnic Inst., July 1987, pp. 16-18.

195

Chapter 7 Manufacturing and Design of Carbon-Carbon Composites Robert L. Burns Fiber Materials, Incorporated Biddeford, Maine

Introduction 198 Composite Design 200 Discontinuous Reinforcement 200 Filament Wound CC 201 Laminates 202 Through-the-Thickness Reinforced CC 203 Thick-Walled Constructions 204 Design Summary 209 Carbon-Carbon Composite Densification 2 11 Laminates 2 12 Thin-Walled 3-D Composites 213 Thick-Walled 3-D Composites 214 Composite Properties 2 15 References 22 1

197

C C Materials and Composites

Introduction The manufacturing and design of carbon-carbon (CC) composites are considerable topics that require extensive treatment. This chapter is intended as an overview of the principles involved with carbon-reinforced, carbon-matrix composite materials. The manufacturing and design of CC composites are founded upon principles common to the design and fabrication of all composites with additional considerations specific to carbon and graphite. Principles common to the design of all composites include the following: the choice of composite preform is driven by the intended application; the reinforcing maqerial and the matrix material must be compatible over the range of temperatures the composite will experience both in processing and intended application; and the properties of the matrix material and the reinforcement material (individually) must be such that the two materials together can interact to provide the desired composite properties. The term preform means the shape and distribution of the reinforcement in the composite. For example, a composite consisting of a uniform distribution of spheroidal silicon carbide particles in an aluminum matrix comprises a metal matrix composite having macroscopically isotropic properties. In contrast, a cured epoxy reinforced with continuous glass fibers all aligned in the same direction comprises a uniaxial composite which results in high anisotropy of properties. Compatibility refers to chemical and thermal properties. Chemical reactivity or a large mismatch in coefficients of thermal expansion (CTE) between matrix and reinforcement usually will render the materials incompatible if elevated temperatures are experienced by the composite either in processing or in use. Compatibility is typically not an issue with CC composites because of the chemical and physical property equivalence of the resulting matrix with the reinforcement upon completion of processing and because of the very low reactivities between the carbon reinforcement and the matrix precursor(s) during processing. Considerations specific to CC center on the unique properties of carbon and graphite. Graphite is the most refractory of materials and combines low density with desirable mechanical, thermal, and electrical properties. The low strain capability and susceptibility to thermal shock of monolithic graphite are overcome by CC composites, which retain most of the advantageous characteristics of monolithic graphite. Even though strain capabilities of CC composites are sometimes similar to monolithic graphite, load-carrying capability is significantly higher and catastrophic failure is less likely to occur. Some properties of a CC and a commercial graphite are compared in table 1 (ref. 1). The very low coefficients of thermal expansion exhibited by CC composites can impart unparalleled stability for applications such as space structures that

198

Property Density, g/cm3 (Ibmlft3 )

3-D Mod 3

ATJ-S

Carbon-Carbon

Graphite

1.60-1.70 (99.89-106.13)

1.82 (113.62)

X direction

3700 (25.5)

4000-4800 (27.6-33.1)

2 direction

15 000 (103.4)

Tensile strength, psi (MPa)

Tensile modulus, psi (GPa) Z direction

3.3 X lo6 (22.75)

6 1.2-1.56 X 10 (7.72-10.75)

7000 (48.26)

12 000-12 150 (82.74-83.77)

3.3 X lo6 (22.75)

0.97-1.26 X lo6 (6.69-8.69)

0.3

1.0-1.5

Compressive strength, psi (MPa)

Z direction Compressive modulus, psi (GPa)

Z direction CTE, Z direction in/in/'F

X 1O-6, room temp. to 2000'F

Thermal conductivity, Btu/hr-ft-'F at 500'F

30

66

at 1500'F

20

43

CC Materials and Composites temperature experienced by the composite during graphitization. Pyrolysis and graphitization are processes unique to the manufacture of carbon-matrix composites. Of these, only carbon-reinforced carbon-matrix composites, or CC, have been important engineering materials up to now. The remainder of this chapter will treat CC composites exclusively.

Composite Design To design a CC composite is to predetermine its woven preform configuration, target density, and degree of stabilization. Matrix morphology (pore structure and degree of process-induced cracking) can also be part of the material design. Density, degree of stabilization, and matrix morphology are determined by selection of the matrix precursor in combination with thermal processing; thus, manufacturing methodology is an important factor in the design of the composite. The most fundamental aspect of CC design, the element which typically is considered first, is the woven preform. Substructure or preform, to a greater extent than any other design consideration, delineates the properties of a CC composite. For continuous fiber-reinforced CC composites, substructure or preform design is referred to as fiber architecture. Fiber architectures include filament wound patterns, yarn diameters, and spacing in unidirectional tapes, rods, and cylinders, patterns in woven fabrics, stacking patterns of laminates, braid patterns, and multidirectional patterns woven into preform constructions. Multidirectional woven preforms may have an orthogonal three-directional (3-D) fiber architecture, more complex n-D architectures where n 2 3, or a cylindrical architecture having reinforcing fibers in the axial, radial, and circumferential directions. The majority of CC applications have typically required continuous fiber reinforcements.

Discontinuous Reinforcement Carbon-carbon composites reinforced with discontinuous fibers have had a very limited range of applications. Perhaps the most important application is that of high-temperature insulation for vacuum or inert atmosphere furnaces. Porous carbon is cost competitive with ceramic furnace insulation and offers an economical alternative to refractory metal heat shields. Short carbon fibers impart sufficient rigidity and toughness to high (-90 percent) porosity carbon for machining and shape retention. Since thermal conductivity in CC is considered a fiber dominated property, all the fibers should be oriented randomly in-plane or in concentric circles in the case of cylindrical pieces, and should be perpendicular to the thickness of the material. Freestanding CC insulating materials are available in various shapes for use at temperatures up to 30OO0C.* A typical component is illustrated in figure 1.

*High-Temperature Insulation. Fiber Materials, Inc., Biddeford, Maine, 1983.

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Manufacturing and Design of Carbon- Carbon Composites

Figure I . Porous CC furnace insulation. (Photo courtesy of Fiber Materials, Incorporated.)

Filament Wound CC Among the continuous fiber-reinforced CC composites, filament wound structures are perhaps the least common. Filament winding was investigated as a means of overcoming a weakness of nonintenvoven pattern ply laminates, namely, their tendency to delaminate when subjected to shear loads in thermal environments (refs. 3 and 4). The use of helical winding patterns at intervals along the structure provides a degree of interlocking of the continuous fibers. Filament winding of relatively thin wall bodies of revolution is economical and is commonly done for the manufacture of other types of composites. Alternative approaches, such as tape wrapping, involute construction, or weaving 3-D preforms, have been preferred to filament winding for overcoming delamination problems in CC composites. One current application of filament wound CC is for hot-pressing molds. The filament winding of continuous yam into a right circular cylindrical pattern imparts high hoop tensile strength to the composite. Filament wound hot-pressing molds are manufactured under the trademark of FilcarbTM*. These filament wound molds *Filcarb-High Strength Long Life Hot-Pressing Molds and Pistons. Fiber Materials, Inc., Biddeford, Maine, 1978.

201

CC Materials and Composites have tensile strengths 5 to 10 times those of monolithic graphite, are resistant to thermal shock, and exhibit a long lifetime compared to bulk graphite molds. Filament winding of unidirectional tape also has been used to apply circumferential reinforcement to 3-D cylindrical CC structures (ref. 5 ) and is discussed later in this chapter.

Laminates Carbon-carbon laminates find extensive use in the aerospace industry with applications including brake assemblies and structural panels. Various fabrics woven of graphite yam typically comprise the reinforcement, but uniaxial tape may be preferred in certain cases. The fiber architecture of the laminate is determined by the weave pattern in the fabric and by the relative orientation of the fabric layers in the stacking sequence. Fabrics are available in a variety of styles (ref. 6). Fabric styles vary according to basic weave pattern, the type of yam, yarn spacing, yam packing efficiency, and the percent of yarn in each direction. The two weave patterns most widely used in CC laminates are plain weave and satin weave. A plain weave fabric and a five-harness satin weave are illustrated in figure 2. Uniaxial tape obviously provides the most highly directional properties in a layer, while a balanced plain weave results in the most evenly matched orthotropic distribution of properties within a layer. Satin weaves, which produce smooth fabrics having good drape, are preferred for applications in which the layers must conform to topological features such as depressions and bends in the macroscopic shape of the end product (e.g., a panel with complex curvature). Laminate structures afford the designer the ability to tailor in-plane properties within the limits of the in situ reinforcement properties. Strength in a given direction is determined by the yam strength and the volume fraction of yam oriented in that direction. In general, off-axis properties are difficult to predict. The assumption underlying such an approach is one of efficient load transfer between yams of different orientations. In CC composites, load transfer between fibers is poor, depending on shear modulus. In some instances, depending upon the processing approach used, the designer will assume no load transfer at all. Not surprisingly, CC laminates exhibit low out-of-plane tensile strength and low interlaminar shear strength. Out-of-plane thermal conductivity is lower than in-plane conductivities, since the carbon fibers comprise the high conductivity “heat pipes” in a CC composite. Anisotropy in thermal expansion can be as high as 15:l depending on the selection of reinforcing fibers. Laminate structures generally are termed 2-D CC even though rotation of fabric plies could be termed multidirectional. For example, the adjacent layers of a certain

202

Manufacturing and Design of Carbon- Carbon Composites

Fill (Y)

LJQLRJL d--v%-k (a)

Fill (Y)

(b)

Figure 2. Schematic offabrics: (a)plain weave; (b)five-harness satin weave (reprinted with permission from ref. 7).

laminate may be oriented so that the fibers in'one layer are directed at 45' to those in the next layer. Thus, a 2-D CC structure can have fibers oriented in more than two directions.

A 2-D CC is suitable for most applications in which interlaminar shear strength and out-of-plane tensile strength are not critical; 2-D CC also is typically less costly than multidirectionally reinforced @e., n > 2) CC, discussed subsequently. The cost differential originates in weaving-related expenses, since fabrics typically used to reinforce 2-D CC are mass produced to give economies not normally achievable with the weaving of individual fiber reinforcements used for 3-D preforms. The higher cost of 3-D CC has inspired alternative fabrication schemes to permit the use of 2-D material in situations in which 3-D would otherwise be necessary.

Through-the-Thickness Reinforced CC Through-the-thickness reinforcement (typically termed 3-D) overcomes some of the limitations of 2-D CC. Braiding and weaving are the two methodologies employed to fabricate preform structures having through-the-thickness reinforcement.

203

CC Materials and Composites

Figure 3. Braided construction rocket motor exit cone. (Photo courtesy of Fiber Materials, Incorporated.) Braiding is appropriate for making curved thin-walled structures such as the small rocket motor exit cone, shown in figure 3. Braided components can be fabricated with either through-the-thickness reinforcement or triaxial reinforcement which consists of helical and axial fibers. A triaxial braided fiber pattern is shown in figure 4. The braid pattern imparts good longitudinal strength to a thin-walled tube, as would a uniaxial structure resulting from axial fibers. The off-axis fibers in a braided construction tie the structure together and impart bending and torsional stiffness to the structure. These attributes represent significant improvements that make braided tubes preferable to uniaxially reinforced tubes. Thin-walled parts can be constructed of 3-D tape, as illustrated in figure 5. Tape thickness for a single-layer construction depends on the cross section of the yarn used. Typical 3-D tapes are multilayered and can be varied in thickness. Threedirectional tape constructions are appropriate for flat panels applications in which improved interlaminar shear strength or resistance to thermal shock is required.

Thick-Walled Constructions Thick-wall multidirectional CC structures include nose tips for atmospheric reentry vehicles and components of solid propellant rocket engines. Weave 204

Manufacturing and Design of Carbon- Carbon Composites

Figure 4. Schematic of braid pattern. architectures for the former are, for the most part, orthogonal 3-D constructions with fine yam spacings. A typical 3-D construction is illustrated in figure 6 (ref. 7). In addition to the high-temperature mechanical property retention afforded by CC composites, reentry vehicle nose tips are required to exhibit unparalleled resistance to thermal shock, predictable thermochemical ablation performance, and the capacity to survive erosive conditions presented to high-velocity vehicles by ice, snow, rain, and dust. Atmospheric reentry environments and their impacts on CC design have been studied extensively over the years (refs. 8 to 11). For optimum mechanical properties, high fiber volume fraction (42 to 48 percent) is desired. Small center-to-center fiber spacings (<0.040 in.) are preferred (especially in the axial direction) for smooth erosion/ablation performance. Low matrix porosity (<6 percent) is also desired for resistance to both thermochemical ablation and particle erosion. A woven orthogonal 3-D preform is shown in figure 7. Carbon-carbon reentry applications other than nose tips may have other performance requirements. Control surface components requiring higher in-plane shear moduli may require a 4-D architecture, such as shown in figure 8, or a 5-D architecture, as shown in figure 9. In both cases the additional in-plane fibers impart higher X-Y shear moduli and overall greater isotropy than can be realized from an orthogonal 3-D architecture. The result of these fiber architectures is a lower fiber volume fraction in each in-plane direction of reinforcement with a resultant loss of tensile and compressive properties in those directions.

205

CC Materzals and Composztes

Figure 5 . Schematic of 3 - 0 tape.

Z

Figure 6 . Typical 3-0 block construction (reprinted with permission from ref. 7).

206

Manufacturing and Design of Carbon- Carbon Composites

F

Figure 7. Woven orthogonal 3-0 preform. Materials, Incorporated.)

(Photo courtesy of Fiber

Where off-axis properties are of primary importance and on-axis property requirements can be sacrificed, a 4-D architecture of the type shown in figure 10 may be appropriate. A seven-directional CC has been manufactured, and 11-D CC constructions are feasible. In genera;, an increase in the number of directions of reinforcement serves to enhance off-axis properties at the expense of on-axis properties, rendering the CC composite more isotropic in its properties and lowering the maximum possible fiber volume fraction, hence, the preceding conditions influence the directional properties in any one direction. For thick-walled constructions that are cylindrical or conical in shape, a polar 3-D architecture, shown in figure 11, is appropriate. A fully densified cylindrical ring having this fiber architecture is shown in figure 12. The Z-direction fibers of a polar preform provide axial tensile and axial compressive strength; the circumferential fibers provide hoop tensile strength; the radial fibers impart radial compressive strength and torsional shear strength to the construction. While weaving of dry carbon yams is one method of constructing 3-D cylindrical preforms, it is not the only method that has been used. An alternative approach uses prefabricated (Le., pultmded) rods that form the radial array and automatic winding of the axial and hoop directions using dry carbon yam (ref. 5). The cylindrical preform is then processed to yield a CC structure.

207

CC Materials and Composites

X

Figure 8. Schematic of 4 - 0 construction.

v---%wFigure 9. Schematic of 5-Dconstruction.

208

Manufacturing a n d Design of Carbon- Carbon Composites

Figure 10. Schematic of off-axis4 - 0 construction. Orthogonal 3-D constructions may also be fabricated by a method other than weaving continuous yarns through the length, width, and height of the preform. This method stacks layers of woven fabric according to a desired 2-D orientation sequence and pierces the layers of fabric by locating graphite yarns that provide the Z-direction reinforcement (refs. 12, 13, and 14). This technique permits higher in-plane fiber volume fractions than can be realized with 3-D orthogonal weaving. A drawback of this technique is that in-plane fibers are broken during the placement of the Z-yarn bundles. The 3-D CC composite whose properties are listed in table 1 was constructed by this method.

Design Summary The foregoing discussion concludes a brief survey of carbon preform design. The range of possible substructures is as broad as for any other fiber-reinforced composite system. Carbon-carbon is primarily an aerospace material because of its advantageous properties, including its unparalleled ability to withstand hightemperature environments. Applications imposing the most complex thermal and mechanical loads are best served by CC composites having through-the-thickness reinforcements. These materials typically result in higher cost than laminates as a result of the higher weaving costs.

209

CC Materials and Composites

Radial

Circumferential

Figure 1I . Three-dimensional cylindrical weave construction.

Figure 12. Fully densified 3-0CC cylindrical ring. (Photo courtesy of Fiber Materials, Incorporated.) Because of the high costs of densification processing, multidirectiond (3-D) CC composites are only used where their unique properties will serve specific

210

Manufacturing and Design of Carbon-Carbon Composites requirements such as in reentry missile nose tips, or where a long-term savings realized from using CC makes it cost-competitive with less expensive materials or fabrication approaches.

Carbon-Carbon Composite Densification Densification processing, like the design of a woven carbon preform, offers many approaches and alternatives to reach the final composite product. Selection of the process is dependent on final application, preform architecture, and cost considerations. Densification processing is typically carried out using either a gaseous (i.e., chemical vapor deposition) or liquid (i.e., resin or pitches) precursor material. Each of the processes entails numerous variations including treatment of the impregnants, temperatures, gas ratios, pressures, and process sequences. The commonality is that the objective of the process is to fill the voids and interfaces of the woven preforms with matrix material as rapidly and cost effectively as possible while meeting the requirements of the application. The most widely used impregnants for densification are gases (methane) and resins or pitches. The resulting properties of the carbons these precursors form vary significantly, depending on the process parameters and the basic chemical nature of the impregnant.

In almost all instances, the objective of densification is to fill a void or volume formed by the woven structure. Void volumes typically range from 35 to 60 percent. The actual mass that is added to the composite will vary, depending on the selection of the impregnant. For example, pitch precursor materials such as petroleum pitch or coal tar pitch form high density matrices in excess of 2.1 gm/cc when heat treated to more than 20OO0C. Alternately, thermosetting type resins such as phenolic or furfural-based resins result in lower specific gravities owing to their inability to form intermediate graphitic crystals. Rapid and cost-effective densification is a result of the volumetric yield the impregnant provides when fully processed. An untreated pitch precursor nominally provides 50 percent carbon yield when pyrolyzed using atmospheric pressure and can provide as high as 90 percent yield when pyrolyzed at pressures exceeding 100 atmospheres. However, the volumetric yield is low when considering the change in density from an impregnant precursor (1.2 gm/cc) to the carbon formed (2.1 g/cc) subsequent to pyrolysis and heat treatment. Therefore, repeated impregnation and pyrolysis are required to obtain a low volumetric porosity in the woven structure. The following section discusses typical processes used to densify or fabricate various forms of CC composites. Again, note that variations to the processes discussed are applicable.

211

CC Materials and Composites Laminates Densification of 2-D CC reinforced laminates starts with prepregging of fabric or aligned yams. Prepregging of woven fabric is readily accomplished on a largescale basis and is typically automated. For graphite fabric destined for CC, the resin impregnant is commonly a high char (carbon) yielding phenolic, although other resin systems may be used. The process consists essentially of passing the fabric through a controlled resin bath and depositing the desired amount of resin pickup, followed by curing or staging prepregged fabric through thermal treatment. The process of making laminates from prepreg consists of cutting the prepreg material to the desired size, stacking the layers in a predetermined sequence of orientation, and molding. During molding, pressure is applied to the stack while heat is applied to cure the resin. The time-temperature-pressure requirements for molding depend upon the characteristics of the prepreg. For phenolic resins, molding may be conducted at 160°C to 18OOC under a pressure of 200 psi to 300 psi for 30 min to 60 min. Molding is followed by a lengthy 24 hr to 100 hr postcure to temperatures higher than those of the molding temperature but well below those of the resin decomposition temperature. A postcure temperature of 2OO0C is typical for a phenolic resin. Following postcure, the temperature is raised to the 65OOC to 8OOOC range to pyrolyze the resin, leaving a carbon matrix. Postcuring and pyrolysis are carried out under an inert atmosphere to prevent oxidation. After the resin has been pyrolyzed, the laminate is a CC composite. A stabilization or graphitization heat treatment follows to a temperature that will match or exceed the maximum temperature that the composite will experience in subsequent processing or in use. The purpose of any subsequent processing would be to reduce porosity and improve mechanical properties. Repeated resin impregnation, pyrolysis, and graphitization; several cycles of pitch impregnation, pyrolysis, and graphitization; CVD densification; or a combination of these, may be done. Pitch-based processing will be described in the context of processing 3-D CC. Subsequent resin impregnation of 2-D CC panels typically is done by a resin transfer or submerged impregnation technique whereby the part, by itself or with others, is placed in a chamber that is then evacuated to remove gases from the pores of the composite. Following this, resin is admitted to the chamber, filling it. A pressure of 50 psi to 100 psi nitrogen is then applied to assist the penetration of the resin into the panel. The excess resin is then drained off, or the panel is removed from the resin bath and a low pressure (-50 psi nitrogen) resin cure is done. Because the fiber volume fraction has been fixed by the initial molding process, higher pressures are not required for this cycle. Postcure, pyrolysis, and graphitization follow the cure cycle. The sequence of impregnation and thermal processing may be repeated a number of times until the target density of the laminate is achieved. Figure 13 illustrates the manufacture of 2-D CC from prepregged fabric or unidirectional tape.

212

Manufacturing and Design of Carbon- Carbon Composites

Figure 13. Typical 2 - 0 CC process schematic. An alternative densification process, or a final one following a series of resin transfer and stabilization cycles, is chemical vapor deposition or CVD (ref. 15). In this process a carbon-carrying gas such as methane is diffused into the CC substrate and thermally decomposed to deposit carbon onto the interior and exterior surfaces of the composite. The process can be economically accomplished by processing a large number of parts in a single cycle. For example, CC brakes are CVD densified (ref. 2). The CVD-densified thin-walled CC tubes have exhibited improved strength and stiffness over otherwise identical tubes densified by the resin transfer method (ref. 16). The improvement has been attributed to small pore filling and a degree of interfilament bonding in the CVD-densified tubes.

Thin-Walled 3-D Composites The resin transfer process is the preferred method of initially impregnating thinwalled braided construction and 3-D tapes. The process can be carried out with minimum pressure gradients across delicate unrigidized parts, thereby avoiding warpage or other damage to the parts. Subsequent to the first set of thermal processing cycles, the parts can be further densified by any of the available methods. The CVD process is especially effective for densifying thin-walled CC pieces because of the short diffusion distances presented by the thin-walled structures.

213

C C Materials and Composites Thick-Walled 3-D Composites Thick-walled through-the-thickness reinforced preforms can be densified by means of CVD, liquid resin, or pitch. Uniform densification of large, thick pieces by CVD is difficult. Limited penetration of the carbon-bearing gas prior to its decomposition usually results in matrix buildup near the surface of the part and high porosity in the middle. Liquid-resin impregnation normally requires a larger number of cycles than pitch densification to achieve the same density in the composite. Higher char and volumetric yields are obtainable from pitch by high-pressure carbonization. Pitch impregnation is carried out at elevated temperature (-250°C) to reduce the pitch viscosity sufficiently for it to flow and penetrate the preform. A graphite frame with numerous openings to accommodate the passage of viscous liquid is employed to protect the preform and maintain its shape. The furnace chamber containing the part is evacuated and heated; hot pitch is then admitted to the chamber and penetrates the evacuated preform. Subsequent pyrolysis or carbonization of the pitch is best done under pressure. The principal benefit is a higher carbon/volumetric yield, which translates into fewer impregnation cycles to achieve the final target density. The impregnated billet assembly is placed in a steel can, which is evacuated and sealed. The evacuated assembly is placed in a hot isostatic press and heated to ~ 6 5 0 O Cunder pressure, typically 1000 atmospheres. The pressure-assisted pyrolysis process is commonly referred to as PIC (pressure-impregnation-carbonization). Although the terminology is not precise, it is widely used. The term “lo-PIC’’ refers to pitch impregnation followed by pressure-assisted carbonization under a relatively low pressure of 300 atmospheres, while a “hi-PIC’’ process is one in which a pressure of at least 700 atmospheres and usually 1000 atmospheres is used. Following PIC, the part is removed from the can, cleaned, and graphitized. Pitch-impregnated billets usually are graphitized to temperatures in the range of 240OoC to 25OO0C, and occasionally in excess of 270OOC. After five cycles of pitch impregnation, “hi-PIC”, and graphitization, composite densities are in the range of 1.85 g/cm3 to 1.90 g/cm3. The pitch-based densification processing of CC is illustrated in figure 14. As discussed previously, the selection of impregnant is made for reasons of processing efficiency and final composite properties. Figure 15 illustrates the effects of pyrolysis pressure on the carbon or char yield of commonly used coal tar or petroleum pitch impregnants. These data are then used to determine the effective volumetric yield of the impregnant as shown in figure 16. Note that figure 16 is theoretical because the density of the impregnant under actual processing conditions is difficult to determine owing to its

214

Manufacturing and Design of Carbon- Carbon Composites

Low-pressure cycle Vacuum impregnate

Carbonize

-

--

Vacuum

[ machine Rough 1 _..~ ~

~~

ressure-carbonization cycle Pressure-carbonization

E l

__ Clean or e machine

Repeat to final density

A

Figure 14. Thermal processing of 3-0 CC billet.

own thermal expansion in a liquid state. For example, selected petroleum pitches have been shown to expand to nearly twice their volume between 100°C to 35OoC under atmospheric pressure. Therefore, the density of the impregnant is not known for processing conditions, especially when high external pressures are used during pyrolysis.

Composite Properties The properties of CC composites are dominated by the reinforcements used and their physical properties, which result from the processing selected and woven preform configuration.In almost all circumstances, physical strength characteristics of the fibers used are reduced as a result of their handling during weaving and the severe environment of densification. Therefore, prediction of composite characteristics typically uses selected test coupon values to establish basic characteristics. Figure 17 indicates that CC composite properties consisting of tensile strength, modulus, strain, and property value are retained as a function of test temperature and fiber volume fraction. These typical composite data are developed based on PAN precursor fibers used in a 3-D orthogonal construction in which the fibers exhibit a tensile modulus of nominally 50 Msi and tensile strength of nominally 400 ksi prior to weaving and densification. The fiber values referred to for tensile

215

105 .-

g !?

3 v) v)

?

2 0 .c

l d

.-N C 0

--

--

103--

/

--

-

-

42 m 102-0

I

104~

I

Manufacturing and Design of Carbon-Carbon Composites and modulus properties are in the direction of interest. For example, a 100-percent fiber volume references a unidirectional (Le., 1-D) composite that has a fiber fraction of typically 50 percent. When considering a 3-D orthogonal CC composite, a Z-direction fiber volume of 40 percent is typical, with a resulting fiber fraction of under 20 percent when considering that total fiber fraction of the composite is less than 50 percent. Therefore, the predicted tensile strength would be nominally 50 ksi to 60 ksi. One of the unique characteristics of CC composites, similar to monolithic graphite, is the increase in strength properties with increasing temperature to approximately 40OO0C. The predictions of properties indicated in figure 17 are intended to act as a guideline. Significant variations may be obtained by both the selection of reinforcement and, to some extent, the densification process. In summary, CC composites may be fabricated for a variety of constructions ranging from unidirectionally reinforced composites to n-directions of reinforcement. Selection of the woven configurations is mandated by final application and associated costs. Manufacturing processes such as braiding, filament winding, polar coordinate weaving, chopped fiber casting, involute, fabric weaving, and multidirectional weaving have all been successfully applied. The associated densification processes of CVD, resin/pitch impregnation, or combinations are again selected for the requirements of the applications. The basic methods of preform fabrication and densification have been described to indicate the techniques developed over the past 20 years. Numerous variations of these processes are applicable. Selection of the appropriate approaches (reinforcement, preform fabrication, and densification) is dependent on the desired composite properties and intended environment of the composite applications.

217

CC Materials and Composites

35 30 E 25 2 cn 20 3 -

2 U

r"

15 10 5 0

0

Fiber volume, percent (a) Modulus as function of fiber volume in direction of applied load.

+ 1.2 L

W

w'l.0 U i

a .8

-

.-c

a .6 2 v) 3 .4 3

U

0

2

.2

0

1000

2000 3000 Temperature, O F

4000

(b) Retention ratio of modulus as function of temperature. Figure 17. Typical CC properties.

218

5000

Manufacturing and Design of Carbon-Carbon Composites

175 .Y a

150

5-125 m

5 100 4 I -I

a a .-

75

a E

50

?

25 0

0

Fiber volume, percent ( c ) 7c~ll.silc. S t I ~ C I l , ~ (1.Y IIl

4-I

, ~ ~ 1 1 ~ ~ /)f~,fil?cr 1 ~ 0 1 1 I ~ o l r l r l l ci~l l d ; l ~ c f ~ l i o/l~l f ' ( l / ~ p l ; ilo(r(/. yl

2.0

L

T m

- 1.5

U

a

.-c

2 a 1.0 L

5 m

c .5

F

5 0

1000

2000 3000 Temperature, O F

4000

5000

( d ) Retention ratio of tensile strength and compressive strength in direction of applied load. Figure 17. Continued.

219

CC Materials and Composites

.6 .5

.4 +

c

a,

2

.3

a,

Q

.! .2i

2 65

.1

05

-. 1 0

1000

2000

3000

4000

5000

Temperature, O F (e) Typical thermal strain of CC as function of temperature. Figure 17. Concluded.

220

Manufacturing and Design of Carbon- Carbon Composites

References 1. DiCristina, V.: Hyperthermal Ablation Performance of Carbon-Carbon Composites. AIAA Paper No. 71-416, Apr. 1971. 2. Arnold, M. R.: Carbon Brakes-Their Development and Service History. World Aerospace Profile 1986, Arthur Reed, ed., Sterling Publ. Ltd. (London), 1986, pp. 237-239. 3. Frye, E. R.: Carbon-Carbon Materials for Ablative Environments. Technol., vol. 12, no. 1, Sept. 1971, pp. 93-107.

Nucl.

4. Stoller, H. M.; and Frye, E. R.: Processing of CarbonlCarbon Composites: An Overview. SC-DC-.713653, Sandia Labs., Apr. 1971.

5. Mullen, C. K.; and Roy, P. J.: Fabrication and Properties Description of AVCO 3D Carbon/Carbon Cylinder Materials. Materials Review for ’72, Volume 17 of National SAMPE Symposium and Exhibition, SOC. of Aerospace Material and Process Engineers, 1972, pp. 111-A-TWO-1-111-A-TWO-8. 6. Engineers’ Guide to Composite Materials. American Soc. for Metals, c. 1987. 7. McAllister, Lawrence E.; and Lachman, Walter L.: Multidirectional CarbonCarbon Composites. Fabrication of Composites, A. Kelly and s. T. Mileiko, eds., Volume 4 of Handbook of Composites, Elsevier Science Publ. Co., Inc., 1983, pp. 109-175. 8. McHenry, M. R.: AFWAL Nosetip Erosion Evaluation Tests in Track G-Data Reduction and Results. TR-82-11/ATD, Acurex Corp., 1982. 9. Stultz, J. W.; and Williams, R. R.: Nose-Tip and Heat-Shield Tests in the Hip Arc Heater Facility. MDC-00608 (Contract No. N60921-75-C-0250), Nov. 1976. (Available from DTIC as AD BO16 213L.) 10. Dirling, R. B., Jr.; Kratsch, K. M.; and Jortner, J.: Ablation/Erosion Evaluation of Reentry Vehicle Materials-Volume 1: Nosetip Materials. AFML-TR-76-2, vol. 1, U.S. Air Force, Feb. 1976. (Available from DTIC as AD BO12 034L.) 11. Wolf, C. J.; Nardo, C. T.; and Dahm, T. J.: Znterim Report Passive Nosetip Technology Program, Volume 22. Coupled ErosionlAblation of Re-entry Materials. SAMSO-TR-74-86, Acurex Corporation, 1975.

221

C C Materials and Composites 12. McAllister, L. E.; and Taverna, A. R.: Development and Evaluation of Mod 3 CarbonICarbon Composites. Materials Review for '72, Volume 17 of National SAMPE Symposium and Exhibition, SOC. of Aerospace Material and Process Engineers, 1972, pp. 111-A-THREE-1-111-A-THREE-7. 13. McAllister, L. E.; and Taverna, A. R.: A Study of Composition-Construction Variations in 3D CarbonKarbon Composites. Paper Presented at Pacific Coast Regional Meeting of the American Ceramic Society, Los Angeles, 1974. 14. McAllister, L. E.; and Taverna, A. R.: A Study of Composition-Construction Variations in 3D CarbonKarbon Composites. Proceedings of the International Conference on Composite Materials, Volume 1, Metallurgical SOC.of AIME, 1976, pp. 307-326. 15. Warren, J. W.; and Williams, R. M.: Isothermal CVD Processing. Non-Metallic Materials Selection, Processing and Environmental Behavior, Volume 4 of National SAMPE Technical Conference Series, SOC. of Aerospace Material and Process Engineers, 1973, pp. 623-633. 16. Wingard, B. L.: Testing and Evaluation of Missile Materials-Task VI: Carbon-Carbon Materials for Space Structures. SRI-MME-89-596-59 12-7-XI, Southern Research Inst., 1989.

222

Chapter 8 High-Temperature Coatings on Carbon Fibers and Carbon-Carbon Composites James E. Sheehan General Atomics San Diego, California

Abstract 224 Introduction 225 High-Temperature Coatings on Carbon Fibers 226 Carbon Fibers 226 Fiber Coating Methods 229 Chemical Vapor Deposition 230 Physical Vapor Deposition 23 1 Electroplating 234 Liquid Precursor Methods 235 Liquid Metal Transfer 236 Coated Fiber Properties 237 High-Temperature Coatings on CC Composites 239 Carbon-Carbon Composites 239 External Coatings 241 Space Shuttle Orbiter Thermal Protection 241 Structural Applications 242 Performance Issues 245 Internal Coatings 249 Liquid Precursor Method 250 Matrix Chemical Modifications 250 Chemical Vapor Deposition 25 1 Advantages and Limitations 252 Temperature Limitations 253

223

CC Materials and Composites Current Coatings 253 Ultrahigh-Temperature Coatings 255 References 256

Abstract The use of coatings to protect carbon fibers and carbon fiber-carbon matrix (CC) composites from chemical attack at elevated temperatures is important for the development of lightweight structural materials for use in advanced aircraft and aerospace applications. Carbon fibers are the optimum reinforcement for many composite systems and are the constituent of choice for hot structures because of their unique retention of mechanical properties at extreme temperatures. Methods that have proved to be particularly effective for producing thin, adherent inorganic coatings on micron-sized carbon fibers within yarns are chemical vapor deposition (CVD), physical vapor deposition (PVD) processes, electroplating, and liquid precursor methods. Ceramic coatings made by CVD have been shown to substantially retard the oxidation of carbon fibers; coatings made by several techniques also have been effective in preventing fiber attack during the fabrication of metal-matrix composites. Although the mechanical properties of high-performance carbon fibers can be significantly reduced by coating, submicron coatings deposited under appropriate conditions produce minimal detrimental effects. The primary reason that coatings are applied to CC composites is to provide oxidation or erosion protection. Except for limited-life rapid heating and cooling applications, external coatings on CC composites are usually applied as multilayer systems. The most successful external coating systems use a hard ceramic as the primary oxygen barrier with a glaze or glass that can flow to accommodate thermal mismatch strains and can seal defects in the hard ceramic coating. A coating system of this type prevents oxidation of the CC that is used to provide reusable thermal protection for the Space Shuttle orbiter vehicles. This general approach is also being employed to develop structural oxidation-protected CC composites for airbreathing engine and hypersonic vehicle airframe applications. The coatings and coating methods now being used for external protection of structural CC composites are S i c and Si3N4 outer coatings made by CVD, inner glass-former coatings of boron compounds made by slurry coating, CVD, and carbide conversion of the CC surface, and bond layers formed by the conversion of the CC surface to Sic. Performance issues associated with the current multilayer external coating systems are coating spallation due to thermal expansion mismatch with the CC, corrosion of the outer coating by the borate glass sealants, moisture sensitivity of borate glasses, and high oxygen permeability of borate glasses. These problems are being addressed by chemically modifying the inner coatings and

224

High- Temperature Coatings on Carbon Fibers and CC Composites developing coating arrangements that limit glass formation. Although the high oxygen permeability of borate glasses is a fundamental limitation, optimization of the current external coating approach and the use of coatings on pore and fiber surfaces within the CC should allow hundreds of hours of component performance. Internal coatings have been made by impregnation with liquid precursors, chemical modifications of the carbon matrix, and chemical vapor infiltration (CVI). Adding borate glass-forming powders to chemically modify matrices is currently the most widely used internal coating method because of the simplicity of the process, the effectiveness of borate glasses, and the achievement of high concentrations of active materials. The principal drawbacks of the powder method are nonuniform glass-former distribution, reductions in composite mechanical properties associated with increased ply spacings in fabric laminates, and fiber damage during composite consolidation. These problems are being addressed by the use of fiber coatings, by chemical modifications at the molecular level, and by the optimization of powder particle sizes. The use of borate glasses limits, even for very short times, the present oxidation-protected CC materials to temperatures below 1550OC. Certain shortterm applications do not require the use of borate glasses; here the limits are in the 170OoC to 1800°C range under optimum conditions and are set by the oxidative instability of the S i c and Si3N4 outer coatings. The most attractive oxygen barrier coating materials for temperatures above this range are iridium and Si02. Iridium has the advantages of ultralow oxygen and carbon permeabilities, and an excellent chemical compatibility with carbon. On the other hand, the high coefficient of thermal expansion (CTE) of indium is a serious drawback for using the material as an external coating on CC. In contrast, vitreous Si02 has a very low CTE and can flow to accommodate mismatched strains; however, Si02 must be isolated from carbon by the use of a system of oxide and carbide layers of questionable thermochemical stability. Ultrahigh-temperature coatings that employ iridium and Si02 as essential constituents are currently under investigation.

Introduction The purpose of this chapter is to review work that has been published in the open literature concerning coatings on carbon fibers and carbon fibercarbon matrix (CC) composites. The chapter focuses on inorganic coatings that are appropriate for providing protection from chemical attack at elevated temperatures. This subject is important because a key factor in improving the capabilities of high-performance aircraft and aerospace vehicles is the development of lightweight structural materials that can operate at very high temperatures in reactive environments. Carbon fibers are among the strongest and stiffest known fibers, surpassing all others in strength retention and creep resistance at elevated temperatures. Superior

225

CC Materials and Composites structural properties combined with a very low density make carbon fibers the constituent of choice for many current composite structures. Carbon fibers have the greatest potential for reinforcing new composites with vastly improved hightemperature capabilities. Although not attacked by many chemicals that are corrosive to other materials, carbon is particularly susceptible to oxidation at elevated temperatures. This is because the oxides of carbon are gases and, therefore, are not protective. Because the most important flight-related applications for high-temperature materials are as air-breathing engine components or structures that must withstand aerodynamic heating, effective oxidation protection is essential to utilize the full potential of carbon fiber composites in advanced aircraft and aerospace vehicles. The application of discrete reaction-barrier coatings to carbon fibers and carbon fiber composites is the most effective method for protecting the materials from chemical degradation. In general, a coating must be both physically and chemically compatible with the underlying material and must resist degradation and permeation by the corrosive species. Furthermore, the coating must adhere strongly to the substrate, but its presence must not detract significantly from the important properties of the fibers or composite. Coatings that meet these requirements enable the development of advanced composites and composite structures for a broad range of high-temperature applications.

High-Temperature Coatings on Carbon Fibers Carbon Fibers At ambient pressures, carbon crystallizes as closely packed atoms in hexagonal arrays. The covalent bonds that exist between carbon atoms in the arrays are the strongest bonds known (ref. 1). When continuity within arrays and alignment of the arrays are achieved, the resulting oriented layered structures can have extremely high strengths as well as elastic moduli in directions parallel to the alignment. Structures of this type are also highly anisotropic because the bonding is weak between atoms in adjacent layers. The strong covalent bonds that are responsible for the strength and stiffness of carbon also make it unique in terms of property retention at high temperatures. The strong bonds inhibit thermally activated atom movements to such a degree that mechanical properties are retained at temperatures up to 220OOC. Other consequences of the strong bonding include very low thermal expansion parallel to the atom layers, and vapor pressures that become appreciable only at temperatures over 275OOC. These mechanical and thermal attributes, along with densities of less than 2.2 g/cc, make carbon materials ideal for high-temperature, flight-related applications.

226

High- Temperature Coatings o n Carbon Fibers and CC Composites The utility of carbon for rocket propulsion and atmospheric reentry thermal protection was recognized early in the late 1950’s and early 1960’s in the development of long-range strategic missiles (refs. 2 and 3) where synthetic graphites were used for thermal shock and thermal erosion resistance. These molded and pyrolytic graphite components could not be used as structures, however, because of strength and elastic modulus limitations. Pyrolytic graphites have a highly aligned layered structure, but growth defects limit tensile strengths and moduli to approximately 140 MPa and 28 GPa in the preferred directions (ref. 3). Carbon is a brittle material, so even comparatively small flaws that are unavoidable in bulk processing prevent realization of the very high strength and modulus potential of aligned carbon structures. The extreme anisotropy of such structures greatly enhances the detrimental effects of crystal boundaries and regions of abrupt misalignment because local thermal expansion differences produce microcracks and residual stresses on heating and cooling due to constraints of the surrounding material (refs. 4 and 5). These difficulties are very pronounced in bulk material but can be minimized when highly oriented carbon is made in the form of flexible fine-diameter fibers. The development of modem-day carbon fibers started more than 30 years ago. Methods for large-scale production of highly oriented micron-sized fibers with minimal defects were identified in the mid-1960’s. Although the defect size that can influence strength is limited by the fiber diameter, eliminating pores, inclusions, surface defects, kinks, and abrupt changes in orientation was found to be essential for producing carbon fibers with consistently high mechanical properties. All of the commercial high-performance carbon fibers currently available are made by the pyrolysis of either petroleum pitch or polyacrylonitrile (PAN) precursors. Highly oriented fibers made from reconstituted cellulose (rayon) were once available but are no longer a commercial product. Inexpensive, low-modulus carbon fibers used for a variety of filler, insulation, and nonstructural applications are now made from rayon and pitch precursors. Commercial carbon fibers are produced by spinning the precursor fibers, carbonizing the precursor fibers by heating in an inert environment to at least 80O0C, and heat-treating the carbonized fibers in inert gas at temperatures between 1200OC and 2700OC. Pitch and PAN are thermoplastics that require low-temperature oxidative cross-linking before carbonization. The fabrication of high-performance carbon fibers is achieved by the alignment of the planar arrays of carbon atoms along the axis of each fiber. This alignment is accomplished in different ways, depending on the carbon precursor. The previously mentioned high-performance rayon precursor fibers were aligned by stretching the carbonized fibers at high temperatures (ref. 6). Currently, a strong preferred orientation is achieved in pitch- and PAN-derived fibers by aligning the precursors.

227

CC Materials and Composites Spinning mesophase pitch aligns the liquid crystal plates that subsequently evolve into planar arrays of carbon atoms during pyrolysis (ref. 7). The PAN fibers are stretched before or during oxidative cross-linking. This stretching produces axial orientation of the polymer structure, which is later converted to carbon (ref. 8). Increasing heat treatment temperatures allows for developing higher degrees of preferred orientation. Commercial carbon fibers that are twisted or untwisted are available in a variety of yam sizes ranging from 1000 to over 10 000 filaments per yarn. Because most high-performance fibers are used in resin matrix composites, all but a few special types are given an oxidative surface treatment to improve resin wetting and bonding. Manufacturers normally apply a protective organic coating, or sizing, to the fibers as the final step in processing in order to ease handling and avoid fiber damage during weaving. Carbon fibers are generally classified according to their elastic moduli, which correlate directly with the degree of alignment. The four basic classifications include low, intermediate, high, and ultrahigh modulus. Table 1 shows the properties of commercial fibers representative of the four classifications that have been reported by the manufacturers. Decreases in the thermal expansion coefficient and increases in the fiber density with an increasing elastic modulus are consequences of the development of more highly aligned and compact atomic structures. The reduction of fiber tensile strength with increasing elastic modulus in the PAN-precursor fibers is due to defects that develop between an increasingly aligned skin and less oriented core (ref. 8). Pitch-precursor fibers are generally not as strong as PAN-derived fibers at the intermediate and high-modulus levels. This is due primarily to inclusions present in the natural product that are incorporated into the fibers and act as strength-limiting flaws (ref. 7).

In the past several years, PAN-precursor fibers exhibiting unprecedented strengths at the intermediate and high-modulus levels have become available. Recent data from the manufacturers show tensile strengths of 5000 MPa to 5650 MPa at modulus values of approximately 275 GPa and strengths of 4100 MPa to 5500 MPa at modulus values of more than 340 GPa. Because all of these fibers have diameters between 5 pm and 6 pm, it is logical to assume that the strength improvements are the result of more complete alignment of the fibers, with smaller defects due to a more uniform skin-to-core orientation. Carbon fibers are sold as yarns wound on spools. Configuring and consolidating the yarns into preforms or directly into useful shapes can be accomplished by a wide variety of textile processes, including fabric weaving, filament winding, braiding, and multiaxial weaving. Carbon fibers are generally not as flexible as common

228

Type

Filament diameter, Density, Precursor pm g.cc-'

Elastic modulus, GPa/Msi

Tensile strength, MPaksi

Axial thermal expansion coefficient,* (10-6 O c - 1 )

1030-1 380/150-200

3.5

Low modulus

Rayon

9-10

1.4-1.5

Low modulus

Pitch

1 6 11

1.8-1.9

140-170/20-25

1030-1380/15CL200

-

Intermediate

PAN

7-8

1.7-1.8

210-240/30-35

3100-3790/45CL550

Unstable above

High modulus

PAN

6-8

1.8-1.9

340-390/50-57

2210-2760/320-400

1.8

High modulus

Pitch

9-10

1.9-2.0

340-380/50-55

172CL2070/250-300

1.8

34-55/5-8

120O0c

modulus

Ultrahigh modulus

PAN

8-9

1.9-2.0

480-520/70-75

1520-1860/220-270

-

Ultrahigh modulus

Pitch

9-10

2.0-2.1

480-520/7CL75

1720-2070/250-300

1.5

Ultrahigh modulus

Pitch

9-10

2.1-2.2

69CL724/106105 1720-2070/250-300

*RT-20OO0C

-

Coating method Chemical vapor deposition, CVD

Coatings TiB, Tic, TIN, Sic, BN, Si, Ta, C

Thicknesses, Pm <0.1 to 1.0

References (14), (17)-(23)

Sputtering

Sic

0.05 to 0.5

(15)

Ion plating

AI

2.5 to 4.0

(27)

Ni, Co, Cu

0.2 to 0.6

(141, (24)-(26)

Electroplating Liquid precursor Liquid metal transfer

Si02 Nb2C, Ta2C, TiC-TiqSNaCz, ZrC-ZrqSN2C2

0.07 to 0.15 0.05 to 2.0

(16) ( 2 0 (29)

High- Temperature Coatings o n Carbon Fibers and CC Composites keeping the fibers separated. Lowering the temperature also produces the desired effect of a decrease in the reaction rate relative to the diffusion rate. However, low temperatures may compromise the temperature uniformity across the yarn and can result in porous, poorly adherent coatings. A method to ensure thermal uniformity and to improve the quality of CVD coatings made at low pressures and temperatures is to perform the process in a radiofrequency (RF) or microwaveinduced plasma (refs. 32 and 33). Plasma-assisted CVD (PACVD) is currently an important technique for the fabrication of thin films in the microelectronics industry and has been applied to the continuous coating of carbon fibers (ref. 20). Equipment for the continuous CVD coating of carbon fiber yams is described in the literature (refs. 15 and 18). Figure 1 shows a schematic and a photograph of a CVD apparatus at General Atomics which is currently being used for fiber coating. The schematic shows the apparatus for conventional CVD, where an induction-heated graphite susceptor tube is used to heat the yam. The photograph shows the apparatus modified for PACVD operation. In the PACVD mode, a stable glow discharge is established by RF coupling to the process gas mixture at reduced pressures. A scanning electron micrograph of carbon fibers coated with silicon nitride by PACVD at General Atomics is shown in figure 2.

Physical Vapor Deposition The three PVD processes are sputtering, ion plating, and evaporation. All involve vaporizing the coating material in a chamber to produce a flux of atoms or molecules which condenses on the material to be coated. These are essentially line-of-site processes in which little penetration of multifiber yams is normally achieved. For this reason PVD fiber coating is conducted most often by spreading the yams into layers only several fibers thick. In sputtering, a low-pressure inert gas glow discharge is produced by establishing a large electrical potential between a cathode of the coating material and an anode. The substrate to be coated can be the anode or can be situated adjacent to the cathode and anode. Positive inert gas ions bombard the cathode target, dislodging groups of atoms to form a vapor that deposits on the substrate. In evaporation and ion plating, the vapor is produced by directly heating the coating material at low pressures. Ion plating is different from evaporation because the vapor passes through a gaseous glow discharge that ionizes some of the atoms. The glow discharge is created by introducing a gas such as argon at low pressures and negatively biasing the material to be coated. Thus, positive ions are accelerated toward the substrate receiving the coating. All three PVD processes can also be conducted in a reactive mode. This means that the vaporized material can be made to chemically react with an active gas that

231

CC Materials and Composites Gas inlet \

Diffusion -Vacuum Pump

(a) Schematic showing conventional CVD arrangement.

(b) Photograph showing apparatus for PACVD operation. Figure I , Continuous CVD fiber coating apparatus.

232

High- Temperature Coatings o n Carbon Fibers and CC Composites

20 p m

Figure 2. Carbon fibers coated with silicon nitride by PACVD. 1000 X .

has been introduced to the coating chamber. In this way, coatings of compounds that are not easily sputtered or do not evaporate stoichiometrically can be produced. The microstructure of PVD coatings depends on the nature of the substrate and a number of process parameters including substrate temperature. Heating of the substrate is often used to obtain coatings that are free of metastable phases and residual stresses. A general description of the PVD methods and coating characteristics has been given by Bunshah (ref. 34). As indicated in table 2, successful S i c coating of carbon fibers within yarns by sputtering has been described recently in the literature (ref. 15). The coating process was performed continuously with the yarn leaving a supply spool, then moving between an anode plate and a cathode target plate. A take-up spool was used to rewind the coated yarn during the process. Two coating runs were made for each yarn. After the first run, the yarn was turned over so the side that originally faced the anode then faced the cathode, which was the source of the S i c vapor. Process parameters were varied, demonstrating that dense, adherent, amorphous S i c coatings of uniform thickness could be deposited in this manner. Coatings between 0.05 pm and 0.50 pm in thickness were produced.

233

CC Materials and Composites Ion plating also was used to coat carbon fibers (ref. 27). In this work, aluminum coatings were made on individual fibers within yarns containing 6000 fibers. Batch processing was used in which 80-mm lengths of yarn were spread into a layer four or five fibers thick. Uniform and adherent aluminum coatings were reported. Recently, continuous carbon fibers have become available with aluminum, copper, and magnesium coatings deposited by ion plating.* Figure 3 shows a scanning electron micrograph of carbon fibers that have been ion plated with magnesium.

Figure 3. Carbon fibers coated with magnesium by ion plating. 1000 X . (Courtesy of Cordec Corporation.) Electroplating This process is used in many industries for depositing protective or decorative metal coatings. In electroplating, the material to be coated is immersed in a solution containing ions of the coating metal. Electroplating is carried out by making the material to be coated the cathode in a direct-current circuit. Current is then passed between the cathode and an anode of the coating material through the electrolyte solution. This process creates metal ions from the anode and reduces the ions in order to deposit the metal on the cathode. Metals that are routinely electroplated include nickel, copper, tin, zinc, chromium, cadmium, gold, silver, platinum, and *Weher, R. J., Cordec Corporation, personal communication, 1989.

234

High- Temperature Coatings o n Carbon Fibers and CC Composites brass alloys. The quality and adherence of the coatings are strong functions of the processed parameters, the chemical characteristics of the electrolyte, and the nature of the substrate. Equipment that has been developed for the continuous electroplating of the individual fibers in carbon yarns is described in the literature (ref. 25). Nickel- and cobalt-electroplatedcarbon fibers were originally evaluated for use in metal-matrix composites, but this work was discontinued because of chemical incompatibility between the fibers and metals at high temperatures (ref. 24). However, nickeland copper-plated carbon fibers are sold commercially for resin matrix composite applications in which the high electrical conductivity of the fibers is beneficial (ref. 26). Thin electroplated nickel coatings do not detract significantly from the mechanical properties of the fibers, and dense coatings of uniform thickness can be made on individual fibers within the yarns, as shown in figure 4.

20 p m

Figure 4. Carbon fibers coated with nickel by electroplating. IO00 X. (Courtesy of American Cyanamid Company.) Liquid Precursor Methods Liquid precursors to inorganic materials include a wide range of compounds in which organic groups are bound to metal atoms. For coating purposes, the most useful liquids of this type are those that can be gelled to form coherent and adherent 235

CC Materials and Composites solid or semisolid polymeric films. Decomposition of the polymer by slow heating produces the desired inorganic coating. Precursors that do not contain oxygen as a major element are known as organometallics and yield nonoxides (refs. 35 to 38). Metal alkoxides are the principal source of oxides, but reducing conditions also can be used to produce nonoxides from these materials (refs. 39 and 40). A second method is the use of an inorganic solution of the metal or a colloidal suspension of inorganic particles in a liquid. Massive precipitation of solids out of the solution or drying of the colloid allows the solid particles to link up and form a gel. Heating sinters the particles to yield a coherent solid. This method involves aqueous processing and usually yields oxide materials, although carbides and nitrides have been made by adding reducing agents and processing in reducing environments (ref. 40). Individual carbon fibers within yarns were continuously coated with Si02 using tetraethoxysilane (TEOS) as the liquid precursor (ref. 16). The process involved heating the fibers to remove the sizing, running the fibers through an ultrasonically agitated bath of diluted TEOS solution, gelling the TEOS by hydration in steam to stabilize the coatings, and, finally, heating the fibers in argon at 550OC to decompose the TEOS and form the Si02 coatings. A sequence of two runs produced Si02 coatings 0.07 pm to 0.15 pm thick. Carbon coatings were also made on ultrahigh-modulus carbon fibers by passing the fibers through a toluene solution of petroleum pitch, vaporizing the solvent, and then decomposing the pitch in a series of increasing-temperature furnaces. This process was used to improve the adherence of the Si02 coatings. Several organometallic compounds suitable for the development of nonoxide coatings such as Sic, B4C, Si3N4, AlN, and BN have been identified in the past few years; a wide range of metal alkoxides are commercially available for producing oxide coatings (refs. 35 to 40). A scanning electron micrograph of carbon fibers coated with Y2O3 by the alkoxide method at General Atomics is shown in figure 5.

Liquid Metal Transfer This method has been used to deposit carbide coatings on carbon substrates by immersing the material to be coated in a bath of molten tin that contains one to several weight percents of a Group IV, V, or VI metal in solution (ref. 28). The metal in solution reacts with the carbon substrate at temperatures between 90OoC and 150OOC to produce a carbide coating. Carbide coatings, 0.05 pm to 2.0 pm in thickness, have been made on individual carbon fibers within yams using the liquid metal transfer technique (refs. 28 and 29). The coatings appear to be dense and uniform, with time and temperature determining the thickness. So far, only batch processing of short strands has been demonstrated. An apparent drawback to the process is that the tin alloy is retained between the fibers and must be removed by chemical dissolution.

236

High- Temperature Coatings o n Carbon Fibers and CC Composites

20 pm

Figure 5. Carbon fibers coated with Y2O3 by liquid alkoxide coating, gelation, and decomposition. 1000 X . Coated Fiber Properties As stated earlier, the impetus for most of the carbon fiber coating development has been to provide a matrix or to protect the fibers and to enhance wetting by molten metals during the fabrication of metal-matrix composites. Carbon-fiberreinforced aluminum, magnesium, and copper have been the composites of primary interest. Fiber coatings of CVD titanium boride and alkoxide-derived Si02 have proved successful for fabricating these materials (refs. 14 and 16). The ability of nickel coatings made by electroplating and electroless deposition to protect carbon fibers from attack by molten aluminum has also been evaluated. However, the formation of embrittling and nonprotective A13Ni compromised the composite properties (ref. 14). Although carbon fibers are unstable when in contact with nickel at temperatures over 6OO0C, because of the high solubility of carbon in nickel, carbon fibers coated with nickel by electroplating are being produced for low-temperature applications in which improved electrical and thermal conductivities are important (refs. 24 and 26). Copper coatings made by electroplating are significantly more effective than nickel coatings for providing high electrical and thermal conductivities; copper is chemically compatible with the fibers at high

237

CC Materials and Composites temperatures, so the coated fibers can be used to fabricate copper matrix composites (ref. 41). In the past 10 years, substantial effort has been made to develop carbonfiber-reinforced carbon and ceramic matrix composites that can be used for high-temperature structural applications in oxidizing environments. Although the development of external coatings to protect the composites has been vigorously pursued, little work has been conducted in the United States on fiber coatings as a method of internal oxidation protection. However, studies reported by Soviet investigators on CVD Sic- and ZrC-coated carbon fibers have been encouraging (ref. 42). This work has shown an order-of-magnitude decrease in the oxidation rate at 600°C and 800°C in air compared with the uncoated fibers. Increasing the thickness of the SIC coatings from 0.02 pm to 0.10 pm resulted in an additional fourfold decrease in the oxidation rate. Tests in C02 revealed very low oxidation rates for the coated fibers, even at 1000°C. It is clear that the Soviet tests were conducted so that oxidation from the ends of the fibers did not occur; similar tests have demonstrated rapid oxidation of Sic-coated fibers at temperatures above 5OO0C because of attack from the unprotected cut ends (ref. 20). The effects of coatings on fiber mechanical properties are well documented. Generally, studies show that without major strength degradation, adherent coatings change the elastic modulus according to the simple rule of mixtures (refs. 15, 26, and 43). The tensile strengths of high-performance carbon fibers are rarely enhanced by the presence of coatings, although electroplated ductile metal coatings have relatively little effect on fiber strength because of the high failure strain of the coatings (refs. 24, 26, and 43). In contrast, brittle coatings that are well bonded to the fibers tend to decrease fiber strengths because cracks that propagate in the coatings at low strains extend into the fibers and result in premature failure (refs. 15, 17, and 18). This physical effect combined with possible chemical damage to the fiber surfaces during coating can produce substantial fiber strength losses (refs. 19 and 21). Strength reductions can be minimal, however, if ceramic coatings are kept thin and are made under conditions that reduce the tendency for chemical attack. Maximum coating defect sizes decrease with decreasing coating thickness, so coating failure strains increase. Thin coatings produce a more flexible yarn and are less likely to join fibers together, so fiber damage during weaving and handling is less of a factor (ref. 20). Coating conditions such as lower temperatures and carbonrich environments can be anticipated to reduce the effects of aggressive chemicals. Under the best conditions, strength reductions of 10 percent to 15 percent can be expected when 0.1-pm- to 0.2-pm-thick brittle coatings are deposited on 7-pm- to 8-pm-diameter, high-performance carbon fibers (refs. 14, 15, 17, 18, and 21).

238

High- Temperature Coatings on Carbon Fibers and C C Composites

High-Temperature Coatings on CC Composites Carbon-Carbon Composites In the late 1950’s, converting carbon fiber-resin matrix composites into carbon fiber-carbon matrix composites by slow pyrolysis of the resin was found to produce useful high-temperature materials (ref. 44). Densification of these materials by repeated cycles of resin or pitch impregnation and pyrolysis or by infiltration and decomposition of hydrocarbon gases produced bodies with thermal shock and fracture toughness properties far superior to those of earlier synthetic graphite materials (ref. 45). To this day, CC composites are principally known as lightweight, veryhigh-temperature materials with superior thermal shock, toughness, and ablation properties. These attributes and specific tribological properties are emphasized in present aerospace and military applications for CC composites such as rocket nozzles, heat shields, reentry vehicle nosetips, and aircraft brakes (refs. 46 to 50). It is not because of a lack of structural capability that current applications do not involve long-term operation under substantial primary loads. When made with high-performance carbon fibers, CC composites can have strength and stiffness characteristics comparable to those of many metal alloys. Because a variety of fiber weaves are available, properties can be tailored for specific component applications. Carbon-carbon composites with superior mechanical properties contain highly oriented fibers derived from mesophase pitch or PAN resin. Fibers of this type have been thermally stabilized for normal CC processing and have strengths in the 2000 MPa to 2400 MPa range and elastic moduli between 350 GPa and 400 GPa. Composite elastic moduli in the principal fiber directions usually reflect the fiber modulus according to the rule of mixtures. Strength utilization of the fibers is 25 percent to 50 percent of rule-of-mixtures predictions, depending on the fiber architecture and processing specifics (refs. 5 1 to 53). Similar to many fiber composites, shear and cross-fiber tensile properties are often a major issue and frequently have a strong influence on design and materials selection. In general, the shear strengths and moduli of CC composites are low compared with those of more conventional materials. The same limitations exist for tensile properties in directions normal to the fiber directions because of processing that creates less than optimum bonding between the matrix and fibers. One of the reasons CC composites are attractive is because they can exhibit relatively high values of fracture toughness. These materials belong to a special group of composites in which the failure strain of the matrix is much lower than that of the fibers. As a result, strong bonding between the matrix and fibers produces composites that are both brittle and low in strength because strong bonds promote 239

CC Materials and Composites failure of the fibers at the low strains where the matrix fails. Conversely, relatively weak bonds allow matrix cracking to occur without crack propagation throughout the fibers. Matrix cracking can occur during CC processing as a result of matrix shrinkage during pyrolysis, the thermal expansion coefficient (CTE) mismatch between the fibers and matrix, and when the composite is loaded in service. Strong bonding in carbonization can result in fiber damage and, combined with strong bonding that persists in the final product, can result in weak and brittle composites. In contrast, weak bonds leave fibers undamaged and allow the matrix to crack in service; however, they continue to transfer load to the fibers so the composite can utilize the high strength of the fibers. Shear displacement of the matrix relative to the fibers and eventual fracture of the fibers at points away from the primary zone of matrix failure create the fibrous pullout fracture surfaces indicative of toughness.

A good method of comparing the fracture toughness of two materials is the notched beam work-of-fracture test. This simple test measures the energy required to move a tensile crack through the material in a controlled fashion. Work-offracture tests on strong, fully densified CC specimens containing high-performance fibers showed fracture energy values of approximately 2 x lo4 J.mP2 for a unidirectional material and 3 x lo3 J.mP2 for a fabric laminate when cracks were propagated across the fibers (refs. 54 and 55). By comparison, plastics give values of 1O3 to 1O4 J.m-2, technical-grade ceramics and premium-grade graphites show values less than lo2 J.mP2, wood exhibits values of about lo4 J.mP2, and ductile metals have values of more than lo5 J.mP2 (ref. 56). Carbon materials are unique in their retention of mechanical properties at high temperatures. High-performance fibers exhibit modest tensile strength increases as temperatures increase to about 22OO0C, while the tensile elastic modulus gradually decreases to about 50 percent of the room temperature value at 220OoC (ref. 57). Carbon-carbon composites show similar behavior in the principal fiber directions. The matrix and matrix bonding to the fibers, however, play a large role in determining the effect of temperature on the shear, cross-fiber tensile, and compressive strengths. One reason for matrix cracking is the CTE mismatch with the highly oriented fibers. Cracking is extensive at room temperature. Cracks close as the temperature increases, providing a more coherent and, perhaps, adherent matrix. This phenomenon accounts for increases in shear and crossfiber tensile strengths that are observed with increasing temperature and, together with increasing fiber strengths, is consistent with observed increases in compressive strength. Moduli either decrease or increase with temperature, depending on the influence of matrix cracking relative to the basic thermal reductions in modulus of the constituents. This change appears to depend on fiber architecture and specifics of the CC processing.

240

High- Temperature Coatings on Carbon Fibers and CC Composites The high-temperature time-dependent deformation of carbon fibers and CC composites is currently an active area of investigation (refs. 58 to 60). Although comprehensive scientific and engineering studies are needed to define creep mechanisms and the effects of microstructure, fabrication processing, and thermal and mechanical history on deformation behavior, the available data clearly demonstrate the superior structural capability of CC composites at very high temperatures. For example, conservative interpretation of the data projects creep rates on the order of lop3 percent per hour in the 20OO0C to 220OOC range at stresses of 70 MPa to 140 MPa for unidirectionally reinforced CC composites.

External Coatings The primary reason coatings are applied to CC composites is to provide oxidation or erosion protection at high temperatures. Until recently, CC composites were considered to be nonstructural thermal protection materials based on their excellent performance in a number of rocket propulsion and atmospheric reentry applications. These applications involve extremely high temperatures but require very short lifetimes and take advantage of the unmatched ablative properties of the all-carbon composite. Although ceramic and refractory metal external coatings were considered and evaluated for these applications, they are generally found to be of little or no benefit. The first real need for a coated CC composite was identified during the development of the Space Shuttle orbiter vehicles. Again, CC was used as a thermal protection material, but in this case, the temperatures were much lower than those for earlier applications, and the Space Shuttle material had to be reusable. The success that was achieved with coated CC in the Space Shuttle program provided the impetus for the current development of oxidation-protected CC composites for structural applications in aircraft and limited life of turbine engines, and as airframe components for multiuse hypersonic vehicles. A renewed interest also exists in the use of coated CC composites for rocket propulsion and missile aerodynamic heating applications because of the requirements of advanced higher performance systems.

Space Shuttle Orbiter Thermal Protection The first coated CC was developed in the early 1970’s for use on the National Aeronautics and Space Administration (NASA) Space Shuttle orbiter vehicles. Termed reusable carbon carbon (RCC), this CC is a resin-densified fabric laminate made from low-modulus, rayon-precursor carbon fibers (ref. 44). The RCC material is protected from oxidation by a two-layer coating system consisting of a porous S i c inner layer sealed with Si02 and an outer coating of alkali silicate glass filled with S i c particulate. A pack process is used to convert the CC surface to S i c (refs. 44, 61, and 62), followed by repeated cycles of impregnating the porous converted surface with acid-activated tetraethoxysilane (TEOS) liquid that is gelled and then heated to produce Si02 (ref. 63). The outer coating on the carbon structure is made

241

CC Materials and Composites from a paste of commercial alkali silicate bonding liquid filled with S i c powder. The coating is applied to the structure by brushing and then it is heat-cured (ref. 63). The RCC has been used to construct the nose cap and leading edge components for all of the Space Shuttle vehicles. It has performed well as a reusable thermal protection material through multiple rapid heating and cooling cycles, with shortterm exposures at temperatures of more than 130OOC (ref. 44). The durability of the RCC coating system can be attributed to several factors. First, the RCC has a relatively high CTE and a low elastic modulus in the plane of the fibers. Second, the inner S i c conversion layer is a strongly bonded, integral part of the CC surface. Finally, the outer coating is a fluxed Si02 glaze that can plastically deform at high temperatures to accommodate differential strains and is fluid enough at temperatures of more than 9OO0C to be self-healing. A problem with the RCC outer coating is that the rate of alkali volatilization from Si02 becomes significant in the 120OoC to 130OoCrange, allowing the loss of the fluxing agent over periods of hours to compromise the compliancy and sealing characteristics of the glaze. By the late 1970's, the encouraging test results that had been obtained with the RCC material were well known. NASA's success with CC, combined with disappointments in a Department of Defense (DoD) sponsored high-performance monolithic ceramics program, generated strong DoD support for the development of oxidation-protected CC materials for very-high-temperature structures in advanced military systems. Although it was recognized that major differences exist between RCC and structural CC materials, it is clear that the amount of development required to meet a number of very optimistic design and performance goals has only recently been fully appreciated.

Structural Applications Coated RCC has very modest mechanical properties and a relatively high CTE in the plane of the fibers. A tensile strength of 62.5 MPa, a tensile elastic modulus of 14.5 GPa, and a CTE of 2.6 x lop6 OC-' are average RCC in-plane properties (ref. 44). In contrast, the projected structural applications required tensile strengths in the 207-MPa to 345-MPa range. This dictated the use of high-performance carbon fibers that produce composites with elastic moduli in the 90-GPa to 117GPa range and CTE values of approximately 1.8 x 1O-6 OC-'. Early systematic oxidation testing of the first structural oxidation-protected CC was conducted by two U.S. Air Force contractors in 1981. The CC was a twodimensional laminate made from high-performance PAN-precursor fibers and was designated advanced carbon carbon (ACC). The best results were obtained with the RCC coating system in which the S i c conversion layer was modified with boron (ref. 64). Burner rig testing that involved rapid heating and cooling between 15OoC and 1371OC with half-hour holds at the maximum temperature resulted in several early failures; this testing also resulted in a number of specimens that

242

High- Temperature Coatings o n Carbon Fibers and CC Composites performed for periods of 25 h to 100 h with no weight loss. However, consistently poor performance was shown in static oxidation tests below 76OoC because of obviously inadequate sealing by the glaze overlay coating and the boronated S i c conversion. The initial tests of the ACC structural CC with the boron-modified RCC coating were encouraging for limited-life applications in which the material experiences rapid heating and is required to perform for several hours at temperatures in the 1000°C to 14OO0C range. On the other hand, extended-life applications require hundreds of hours of exposure from 649OC to 1371OC. Oxidation at lower temperatures was a problem with the modified RCC coating system, and a major development effort was clearly needed. Furthermore, it was recognized that the basic RCC coating approach had features that posed serious problems for many of the new structural applications. First, the outer glaze coating was susceptible to flow and particulate erosion. The glaze was also susceptible to alkali loss at high temperatures and acted as an adhesive that prevented movement of mating parts without damage to the coating. A second problem was that the proposed structural parts were often 2.5 mm to 5 mm thick and the airfoil trailing edges could be as thin as 1 mm. Very often the coating approached 1 mm in thickness and the mechanically poor S i c conversion layer constituted a large portion of the part. The outer glaze and coating thickness problems associated with the RCC type coatings have led to the current concept of using a hard, dense outer coating for both limited-life and extended-life applications. The CVD process is used to produce the outer coatings that act as the primary barriers to oxygen ingress. The materials S i c and Si3N4 have received the most attention for outer coatings because of their relatively low CTE values and excellent oxidation resistance up to at least 17OO0C. Recent work for limited-life applications has shown that under conditions of rapid heating to above the coating temperature, in which thermal expansion mismatch cracks in the coating are closed, CVD outer layers can provide excellent CC protection for several hours up to 1750OC (ref. 65). Extended-life applications and even limited-life use at temperatures below that in which the coating cracks are open require a glass sealant to fill the cracks (refs. 65 and 66). Figure 6 shows the current concept of using a boron-rich inner layer to provide the glass for sealing the cracks in the outer coating. Elemental boron, boron carbide, and several configurations of mixed boron compounds with and without S i c and elemental silicon are now under evaluation as inner layers (refs. 65 to 70). Inner layers are being made by particulate slurry coating, CVD, and carbide conversion of the CC surface (refs. 61 to 70). The main purpose of the inner layer is to form a borate glass by oxidation through the cracks in the outer coating. Oxidation of the portion of the inner layer beneath the cracks to form a glass produces a 200-percent to 2.50-percent volume increase that forces the glass into the cracks. In addition to the outer coating and inner glass-forming layer, a

243

CC Materials and Composites

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Figure 6. Current CC composite coating concept. 100 X . (Courtesy of Chromalloy Research and Technology.) base Sic conversion layer that improves bonding to the CC surface will be used often if the glass-forming layer is not a conversion layer. In addition to acting as the primary oxygen barrier, outer S i c and Si3N4 coatings provide hard erosion-resistant bearing surfaces that cover the inner layers and inhibit vaporization of the borate glass sealants. The inner layers provide the sealant glasses but they also must establish and maintain strong bonding with the outer coating and CC to inhibit coating spallation. Conversion layers that are integral to the CC surface and have been densified by some pore-filling technique may be optimum for bonding. The use of borate and phosphate glasses to protect carbon bodies from oxidation has a long history (refs. 71 to 73). Boric oxide is particularly attractive because it melts at approximately 45OoC and has viscosity, wetting, and thermal stability characteristics that make it an effective sealant over a wide range of temperatures. In figure 7, the viscosity of B2O3 as a function of temperature is compared with the viscosities of other glasses; the figure shows that the behavior of B2O3 is unique in that the viscosity is low and does not decrease or increase rapidly with changes in temperature (ref. 74). Figure 8 shows that the surface tension of B2O3 is also low compared with that of other glasses, and this translates into excellent wetting on most materials (ref. 74). The ability of B2O3 to wet and flow over a broad temperature range results in the formation and maintenance of protective glass films within coating cracks and surface pores of the CC substrate. Phosphate glasses have wetting 244

High- Temperature Coatangs o n Carbon Fibers and CC Composates

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Performance Issues Performance issues associated with the present generation of CC coating systems are coating spallation due to CTE mismatch, borate glass corrosion of

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247

CC Materials and Composites

Spalled region

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exposure because of the degradation of borate glass that had formed beneath the S i c coating. Boric acid can be seen in the outer coating cracks. The relatively high rate at which oxygen permeates borate glasses is not only the cause of outer coating corrosion and the generally negative effects associated with excessive glass formation, but also represents an inherent limitation of the current approach of using borate glasses to seal cracks in outer coatings. The oxygen permeability of B2O3 as a function of temperature is compared with the permeabilities of other materials in figure 11 (refs. 81 to 87). Using this type of information in conjunction with a model of a cracked outer coating with B2O3 filling the cracks, an analysis showed that under ideal conditions even the smallest cracks can allow significant oxidation of the CC substrate over the hundreds of hours of operation required for many important applications (ref. 79). It is known that chemical modifications of the borate glasses can improve moisture resistance (refs. 74 and 80). Furthermore, it may be possible to vary the physical arrangement and chemistry of the coatings to prevent excessive glass formation. Limiting the glass formation would benefit both the moisture and

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CC Materials and Composites produce external glass layers when the glasses are present in high concentrations near the surface of the CC. Coatings have been made on the internal surfaces of CC composites by three techniques: liquid precursor impregnation, chemical modification of the carbon matrix, and CVD. Liquid precursor impregnation is the most chemically comprehensive of the three because with this method a wide range of inorganic coatings can be made. Chemical modifications, however, result in coatings only when a mechanism is present for producing glass layers. Finally, the CVD process is most appropriate for producing nonoxides in the presence of carbon; it can be used to make internal coatings, chemically modify the carbon matrix, or completely replace the matrix.

Liquid Precursor Method The liquid precursor method involves impregnating the porous CC substrate with organometallics, metal alkoxides, metal salt solutions, or colloidal suspensions. Drying, gelation, or chemical conversion produces solid layers on internal surfaces that can then be decomposed or reacted by heating to form ceramic coatings. Multiple impregnation and heat treatment cycles are needed to produce continuous coatings of significant thickness because of the normally low yield of the precursors. This method has been the subject of numerous patents for improving the oxidation resistance of porous carbon bodies in general and recently was applied to CC composites (refs. 77 and 90 to 94). The recent CC work has shown that by virtue of B2O3 glass formation, internal coatings rich in boron are particularly effective in enhancing oxidation performance (ref. 77). As described earlier, the liquid precursor method is used to coat carbon fibers and can be effective either before or after the fibers have been configured into a composite.

Matrix Chemical Modifications Chemical modifications are made to the carbon matrix either by adding the elements or compounds as a powder to the resin or pitch matrix precursor, or by altering the chemistry of the matrix precursor on the molecular level (refs. 68 to 70 and 9.5 to 100). The powder method was originally developed to improve the oxidation resistance of synthetic graphites (refs. 101 to 10.5). Chemical modifications are most effective and produce coherent coatings only when a lowviscosity, wetting glass is formed initially or when exposed to oxygen. This modification is illustrated in figure 12, which shows a region of borate glass formed by the oxidation of powders that were added to the matrix of a CC composite. The glass was produced near the surface of the composite because of oxygen permeation through a crack in the external CVD S i c coating. The glass forms at the expense of matrix oxidation, then coats fibers, fills pores, and partitions off regions of attack to prevent rapid catastrophic oxidation. Glass that originates in the composite directly below the external coating can also migrate to fill the CTE mismatch cracks in the coating. Chemical modification by glass-former powder additions is currently the

250

High- Temperature Coatings on Carbon Fibers and C C Composites

Glass region

200 pm

Figure 12. Borate glass produced within CC composite by oxidation of particulate glass formers. 100 X . most widely used method for forming internal coatings (refs. 68 to 70, 96, 98, and 99). Chemical Vapor Deposition The CVD process within a porous body is termed chemical vapor infiltration (CVI). This process involves the diffusion of gases into the body and the decomposition or reaction of the gases in order to produce a solid product that is deposited on the pore walls and fibers. Chemical vapor infiltration is one of the principal densification techniques used in the fabrication of CC composites (ref. 45). It can also be used to chemically modify the carbon matrix and to produce discrete internal coatings on the carbon constituents (refs. 45, 106, and 107). Unless the carbon fibers are previously coated for protection, the CVI deposit must be a nonoxide to prevent fiber degradation during processing. Improvements in the oxidation resistance of carbon fiber composites were demonstrated when the carbon matrix was replaced almost entirely with CVI matrix materials such as Sic, TIC, and BN (refs. 108 to 111). From the standpoint of oxidation resistance, it clearly is logical to fully densify with the ceramic if the composite alterations resulting from the replacement of the carbon matrix with a ceramic can be accommodated. The microstructure of a fabric laminate composite with carbon fibers in a CVI S i c matrix that was fabricated at General Atomics is shown in figure 13. Chemical vapor deposition is previously described as a 251

CC Materials and Composites

80 pm

Figure 13. Microstructure of carbon fiber-Sic matrix composite fabricated by CVI. 250 X . useful fiber coating method; figure 13 shows that CVI densification actually occurs by the buildup of coatings on the fibers. This type of composite has exhibited excellent short-term oxidation resistance at temperatures up to 15OO0C (ref. 108). Over longer periods of time, oxygen permeation through pores and cracks in the matrix results in significant oxidation of the fibers, which demonstrates the need for effective external coatings.

Advantages and Limitations The advantage of the liquid precursor method for producing internal coatings is the direct application of the coating to areas that are most susceptible to oxidative attack. Disadvantages of this method include the need for multiple processing cycles because of the low yield of most precursors, the need for compatibility between the carbon surfaces and liquid for thorough wetting to yield continuous and adherent coatings, and the need to control the high, potentially disruptive shrinkage that accompanies conversion of the liquid to a solid.

252

High- Temperature Coatings o n Carbon Fibers and CC Composites When used to incorporate glass formers, the matrix chemical modification technique has the advantage of producing increasing quantities of protective material as oxygen intrusion proceeds. This process results in the filling of surface pores and the formation of a coherent glass layer beneath the external coating. Particulate additions can be made simply and in relatively high concentrations. However, powders have the disadvantage of not penetrating the fiber yam bundles and may mechanically damage fibers during composite consolidation. Segregation of the particulate between the fabric plies of CC laminates also increases the ply spacing and results in reduced inplane mechanical properties. These problems currently are being addressed by the use of fiber coatings and optimization of powder particle sizes. Although molecular alterations to the resin or pitch matrix precursors have the advantage of penetrating the yams to produce more uniform protection, the concentrations of active species that can be incorporated by this technique are much lower than can be achieved by powder loading. For this reason, the particulate loading and molecular alteration techniques may be most effective in combination with one another. The CVI method for providing internal oxidation protection is extremely versatile; it can be used to chemically modify the carbon matrix, to coat the internal surfaces, or to completely replace the carbon matrix with another material. This method is limited to nonoxides when carbon is present and is a time-consuming and inefficient process as normally practiced when used to produce a composite matrix. However, recent work has demonstrated that satisfactory densification of fiber preforms can be accomplished in periods of 24 hours or less if the reactive gases are forced through the fibers and a thermal gradient is maintained to create a front of preferred deposition within the composite (refs. 112 to 114). Figure 14 shows the type of forced-flow thermal gradient arrangement currently being used for rapid CVI within deposition chambers at General Atomics.

Temperature Limitations Current Coatings Although actual temperature limitations on the coatings currently being used for CC oxidation protection will vary depending on specific operating conditions and component performance requirements, general guidelines can be established based on experimental observations and the known thermochemical properties of the coating constituents. For example, unprotected B2O3 glass layers were shown to be effective in retarding the oxidation of CC substrates for extended periods at temperatures up to about 1000°C, but volatilization of the glass is limiting at higher temperatures (ref. 77). When B2O3 was used as a crack sealant for hard outer coatings or as a binder in coatings composed of hard oxide particles, oxidation protection was extended to temperatures of more than 12OO0C (ref. 78). Tests have

253

CC Materials and Composites

Figure 14. Fixturing for rapid CVI.

demonstrated that coating systems of the type shown in figure 6, in which a boronrich inner layer provides glass to seal cracks in a CVD S i c outer layer, can provide hundreds of hours of oxidation protection under thermal cycling conditions with maximum temperatures close to 140OoC (ref. 65). The maximum temperature at which B2O3 can be used in contact with carbon at atmospheric pressure is approximately 1575OC because of the disruption of the glass by equilibrium CO reaction product pressures greater than 10-1 MPa (1 atm) at higher temperatures (ref. 79). The vapor pressure of B2O3 at this temperature also is projected to be about lop3 MPa under dry conditions and close to 10-1 MPa in moist environments (refs. 74 and 77). These factors, combined with the tendency previously described for B2O3 to corrode the siliconbased outer coatings, suggest maximum temperatures for the borate glass internal and glass-sealed external coatings of 150OOC to 155OoC even for short periods of performance. Certain applications do not require glass-sealed outer coatings because the component is rapidly heated to and cooled from temperatures high enough to close the coating coefficient of thermal expansion (CTE) mismatch cracks. Relative to this situation, pure CVD S i c and SijNq coatings are viable in the 170OoC to 1800°C range under strongly oxidizing conditions (refs. 65 and 115). This temperature limit exists because at higher temperatures the Si02 layers that normally protect these materials are disrupted by CO and N2 interfacial reaction product pressures that become greater than 10-1 MPa, thus causing the coatings to erode by uncontrolled oxidation. Under less oxidizing conditions, the temperature limit can be reduced to the 1500°C to 160OoCrange by active oxidation (refs. 116

254

High- Temperature Coatings on Carbon Fibers and C C Composites and 117). Gaseous S i 0 is then formed instead of the protective Si02, and coatings can fail by rapid loss of material.

Ultrahigh-Temperature Coatings Once the temperature limit for S i c and Si3N4 coatings is exceeded, identifying coatings with acceptably low oxygen permeabilities becomes the issue. For an external 100-pm-thick coating, estimated coating permeabilities of approximately lo-' and lo-'' g.cm-l.s-l or less would be required to provide acceptable shortand long-term oxidation protection of CC composites (ref. 79). Figure 11 shows that at temperatures above 1800"C, the materials expected to meet these criteria are Si02, Al2O3, and the high-temperature noble metals, rhodium and iridium. At such high temperatures, materials that form refractory oxides such as ZrOz and HfO2 oxidize very rapidly because the oxides are inherently permeable to oxygen. Based on previous work and the considerations discussed here, current ultrahigh-temperature coating development, for all but very short-term applications, is focused on concepts that use iridium and Si02 as essential constituents. More than 20 years ago, extensive work was conducted on iridium and iridium alloy coatings to be used to protect synthetic graphite bodies from oxidation at very high temperatures (ref. 87). Iridium is attractive because of its 2440OC melting point, extremely low oxygen and carbon permeabilities, and inertness to carbon. Coatings were made by a variety of techniques including CVD, electroplating, and a combination of plasma spraying and hot-isostatic pressing. Although excellent protection was demonstrated for short times in the 2000°C to 2100°C range, difficulties in fabricating high-quality coatings and the cost and availability of iridium were deterrents to further development. Iridium and rhodium are also susceptible to erosion by the formation of volatile oxides and have CTE values that are extremely high compared with those of CC composite materials (ref. 118). Coating erosion may or may not be an issue depending on oxygen partial pressure, gas flow conditions, and required time of performance. Oxide overcoatings have been proposed to inhibit erosion, but the CTE mismatch problem is profound and might very well prohibit the use of these materials as external coatings on CC composites. This same problem exists for A1203 and most other oxides. As an external coating, the advantages of Si02 are a very low CTE and an ability to flow at high temperatures to accommodate differential strains. The viscosity of Si02 is about lo7 dPa.s at 180O0C (ref. 74). Then, the glass will adhere on contact, begin to flow under its own weight, and can be considered to behave as a sealant. Because Si02 is reduced by carbon and is unstable in contact with carbides at very high temperatures, Si02 glass must be separated from the CC substrate by a hard oxide or some other compatible material. An outer coating must be used for most applications to prevent flow erosion and excessive vaporization. The vapor pressure of Si02 under strongly oxidizing conditions is projected to

255

CC Materials and Composites be somewhat less than lO-5 MPa at 20OO0C (ref. 119). Low oxygen pressures increase the volatility of Si02 by promoting S i 0 formation. Inner and outer coatings are required to envelop and protect the Si02, which invariably leads to the consideration of materials that are unattractive from the CTE standpoint. It is interesting to note that several refractory oxides are known to exhibit very low CTE values. This phenomenon is a result of lowtemperature microcracking, which produces extreme crystallographic anisotropy. The recombination of the microcracks produces significantly increased CTE values at higher temperatures (refs. 4 and 120). Furthermore, any appropriate oxide in contact with carbon at very high temperatures will form an interfacial carbide. This reaction raises the issue of continued conversion of the oxide to the carbide by transport of carbon through the carbide. The presence of the carbide also reintroduces the CTE problem. The most refractory of the low-CTE oxide is HfTiO4, which melts incongruently at about 1980OC (ref. 121). Materials selection for internal coatings to perform at very high temperatures is even more limited than the selection for external oxygen and carbon diffusion barriers. The coatings need to be chemically stable when in contact with carbon, and it is likely that this stability must result from thermodynamic compatibility rather than transport limitations because of the very short diffusion distances. Because the oxides are reduced by carbon at the temperatures in question and because rhodium forms a eutectic melt with carbon at about 17OO0C, the only material that clearly meets the above criteria is iridium (ref. 118). Totally replacing the carbon matrix with a ceramic would prevent matrix gasification, but experience has shown that this replacement would do little over extended periods of time to protect the carbon fibers from oxidation (ref. 108). Fiber protection would require thin, adherent barrier coatings that are chemically compatible with the fibers and matrix. Again, iridium appears to be the preferred material.

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108. Christin, F.; and Naslain, R.: A Thermodynamic and Experimental Approach to S i c CVD Infiltration of Porous Carbon-Carbon Composites. Chemical Vapor Deposition 1979, T. 0. Sedgwick and H. Lydtin, eds., Electrochemical SOC., 1979, pp. 499-514. 109. Christin, F.; Heraud, L.; Choury, J. J.; Naslain, R.; and Hagenmuller, P.: In-Depth CVD of S i c Within Porous Carbon Materials. Chemical Vapor Deposition 1980, H. E. Hintermann, ed., Lab. Suisse de Recherches Horlogeres (Switzerland), 1980, pp. 154-161. 110. Rossignol, J. Y.; Naslain, R.; Hagenmuller, P.; Heraud, L.; and Choury, J. J.: Carbon-Carbon Titanium Carbide Composite Materials Obtained by CVI of Porous Carbon-Carbon Substrates. Chemical Vapor Deposition 1980, H. E. Hintermann, ed., Lab. Suisse de Recherches Horlogeres (Switzerland), 1980, pp. 162-168. 111. Hannache, H.; Quenisset, J. M.; Naslain, R.; and Heraud, L.: Composite Materials Made From a Porous 2D-Carbon-Carbon Preform Densified With Boron Nitride by Chemical Vapour Infiltration. Mater. Sci., vol. 19, no. 1, 1984, pp. 202-212.

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CC Materials and Composites 112. Stinton, D. P.; Caputo, A. J.; and Lowden, R. A.: Synthesis of Fiber-Reinforced S i c Composites by Chemical Vapor Infiltration. American Ceram. SOC.Bull., vol. 65, Feb. 1986, pp. 347-350. 113. Caputo, Anthony J.; Stinton, David P.; Lowden, Richard A.; and Besmann, Theodore M.: Fiber-Reinforced S i c Composites With Improved Mechanical Properties. American Ceram. Soc. Bull., vol. 66, Feb. 1987, pp. 368-372. 114. Stinton, David P.; Besmann, Theodore M.; and Lowden, Richard A.: Advanced Ceramics by Chemical Vapor Deposition Techniques. American Ceram. SOC. Bull., vol. 67, Feb. 1988, pp. 350-355. 115. Chown, J.; Deacon, R. F.; Singer, N.; and White, A. E. S.: Refractory Coatings on Graphite, With Some Comments on the Ultimate Oxidation Resistance of Coated Graphite. Special Ceramics 1962, P. Popper, ed., Academic Press, Inc., 1963, pp. 81-115. 116. Hinze, J. W.; and Graham, H. C.: The Active Oxidation of Si and S i c in the Viscous Gas-Flow Regime. J . Electrochem. SOC.,vol. 123, no. 7, July 1976, pp. 1066-1073. 117. Singhal, S. C.: Thermodynamic Analysis of the High-Temperature Stability of Silicon Nitride and Silicon Carbide. Cerumurg. Int., vol. 2, July-Sept. 1976, pp. 123-130. 118. Darling, A. S.: Some Properties and Applications of the Platinum-Group Metals. Int. Metall. Reviews, Sept. 1973, pp. 91-122. 119. Schick, Harold L.: A Thermodynamic Analysis of the High-Temperature Vaporization Properties of Silica. Chem. Reviews, vol. 60, no. 4, Aug. 1960, pp. 331-362. 120. Holcombe, C. E.; Morrow, M. K.; Smith, D. D.; and Carpenter, D. A.: Survey Study of Low-Expanding, High-Melting Mixed Oxides. Rep. Y-1913, Oak Ridge National Lab., 1974. 121. Coutures, Jean Pierre; and Coutures, Juliette: The System HfO2-TiO2. J . American Ceram. SOC.,vol. 70, no. 6, June 1987, pp. 383-387.

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Chapter 9 Applications of Carbon-Carbon Louis Rubin* The Aerospace Corporation El Segundo, California

Approximately 30' years ago, carbon-carbon (CC) was developed to meet the anticipated needs of the emerging space programs for materials that were resistant to high temperatures and were able to maintain structural integrity while experiencing the thermal stresses of reentry from space. The utility of this material was first demonstrated in a major Space Shuttle application where it performed on the wing leading edge and nose cap thermal protection system. Carbon-carbon technology has matured considerably since the first Space Shuttle application. Although more advanced versions continue to perform well on the Space Shuttle, CC has evolved as a versatile material for a wide variety of new applications. The key to many of the new uses of CC is in the development of improved oxidation-resistant systems for atmospheric use at high temperatures, of new high-modulus carbon fibers that provide dimensional rigidity and low thermal expansion for structural applications, and of newer matrices based on pitch or advanced high-char yielding resins that result in greater composite integrity and reduced processing time. The newest approach to rapid matrix processing is related to chemical vapor infiltration (CVI), but it is based on a liquid hydrocarbon precursor that is used instead of a gas (ref. 1). One of the more visible applications of CC is on the wing leading edge and nose cap of the Space Shuttle orbiter (refs. 2 and 3). These components, which LTV Corporation manufactured, are exposed to temperatures up to 2800°F (1538°C) during orbiter entry into the atmosphere; in addition, they must provide thermal protection and maintain structural integrity over multiple missions. Each orbiter wing contains 22 leadingedge CC airfoil panels and 22 sealing strips of CC. The nose cap is 4; f t in diameter and consists of the primary cap and eight circumferential seal *Currently with Research Opportunities, Inc., Torrance, California.

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C C Materials and Composites

segments that allow for thermal growth and provide an interface with the adjacent tiles. The key to successful use of CC in this application is with the development of an effective oxidation-resistant coating system. Silicon is diffused into the outer layers of the CC fabric reinforcement in an argon atmosphere at a temperature of 3000'F (1649'C). This diffusion yields an oxygen barrier that is further enhanced by a sealant consisting of tetraethylorthosilicate (TEOS), sodium silicate, and graphite fibers. Details of these CC applications are illustrated in figures 1-3.

Leading Edge Structural Subsystem

Figure 1. CC areas on Space Shuttle orbiter.

In a related but more demanding application, the National Aero-Space Plane (NASP) Program is considering the use of CC in thermally critical areas. The anticipated temperatures of 5000'F (2760'C) on the nose cap and approximately 3500'F (1930°C) on the wing- and tail-leading edges, coupled with minimum weight and high-strength requirements, make CC one of the more attractive candidate materials (ref. 4). A specific NASP application under development at LTV Corporation involves the fabrication of a CC elevon control surface (ref. 5 ) . This control surface will be based on an advanced version of the CC used on the Space Shuttle, and it will also use a silicon carbide-type of coating for oxidation resistance to 3000'F (1649'C). NASA Langley Research Center is sponsoring this effort, with the deliverables scheduled for 1990. The European spaceplane 268

Applications of Carbon-Carbon

Figure 2 . Segment of CC orbiter wing leading edge.

Figure 3 . Space Shuttle orbiter CC nose cap.

Hermes is also using CC, with appropriate oxidation-resistant protection for the nose cap, leading edges, and underwing areas (ref. 6 ) .

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CC Materials and Composites The fabrication, by The B.F. Goodrich Company, Super-Temp, of a 10 ft x 4 ft, unidirectionally, rib-stiffened curved panel to simulate a structural component suitable for use on a hypervelocity vehicle illustrated the application of CC to large-area, structural panels. The panel was based on a two-dimensional (2-D) layup of a balanced-weave, T-300 graphite fabric with a phenolic matrix precursor. The panel (fig. 4) was designed to achieve a uniform fiber volume across the width with strength ranging between 50 ksi to 60 ksi and a tensile modulus of 16 msi in the warp direction (ref. 7). A General Dynamics X-30 vehicle design exemplified the potential use of this structural panel, suggesting that a primary CC structure be used for the entire vehicle. This CC design offered a significant savings in structural weight (ref. 8).

Figure 4. CC rib-stiffened structural panel (The B.F. Goodrich Company, Super-Temp).

Using CC in jet-engine rotors and stators to permit high-combustion temperature operating conditions holds the promise of reduced engine size, weight, and fuel consumption. Higher temperature capabilities reduce the need for engine bypass cooling, thereby increasing engine efficiency. Carbon-carbon offers this potential capability plus a significant weight

270

Applications of Carbon-Carbon

advantage over current high-temperature alloys. Operating temperatures that are well in excess of 1000°F hotter than temperatures used in conventional engines are being sought. The LTV Missiles and Electronics Group is developing CC jet-engine components under the Air Force Extended Long-Range Integrated Technology Evaluation (ELITE) program. This includes turbine wheels, combustion chambers, and exhaust nozzles. A CC turbine wheel that is 14.4 in. in diameter and weighs only 7.5 lb was successfully spun at 28000 r/min at a temperature of 3200°F (1760°C) (ref. 9). The pretest turbine rotor is shown in figure 5.

Figure 5. CC turbine rotor (LTV Missiles and Electronics Group).

An innovation in CC turbine rotors and disks was also demonstrated by the LTV Missiles Division through the use of a polar-weave-based CC in test rotors. Fill fibers radiated from the rotor hub plies in this arrangement with hoop or warp fibers oriented in concentric circles. Spin tests on polar-weave CC rotors achieved approximately 40 000 r/min before failure (ref. lo). Potential applications for CC components are found in conventional high-performance jet engines. Afterburner flaps and seals used to provide thrust control have had prototypes designed in CC. The F-100 jet-engine 271

CC Materials and Composites afterburner nozzle uses coated CC (ref. 11). These CC applications offer weight reduction and higher temperature capabilities than those of conventional metallic materials. A CC prototype jet aircraft turbine engine flap fabricated by HITCO* is illustrated in figure 6.

Figure 6. CC prototype jet aircraft turbine engine flap (HITCO).

Piston-driven engines such as gasoline and diesel engines could operate at increased efficiency and reduced weight if cooling requirements were minimized through the use of high-temperature capability materials. Carbon-carbon has been proposed and is under evaluation for many applications in this area. A CC piston has been tested successfully at NASA Langley Research Center (ref. 12), and a U.S. patent has been issued covering this composite piston concept (ref. 13). Carbon fiber reinforcements of four-dimensional (4-D) angle interlock and a knitted multilayer warp sock were evaluated with resin infiltration and chemical vapor deposition (CVD) carbon densification to complete the CC processing. The CC piston has a lower density than the conventional aluminum lightweight piston. This reduced weight improves internal combustion engine efficiency. However, the very low thermal expansion of CC, relative to aluminum, and the high-strength retention of CC at elevated temperatures allow a pistonto-cylinder wall clearance that is small enough to eliminate the need for piston rings. Lightweight CC pistons permit the use of lighter weight reciprocating components, allowing possible higher engine speeds and improved *Now called BP Chemicals (HITCO).

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Applications of Carbon-Carbon engine performance. Issues such as oxidation life, oil, and fuel infiltration effects and cyclic endurance require continued study, but the prognosis for CC pistons is very promising. An illustration of a CC piston together with an aluminum counterpart is shown in figure 7. Carbon-carbon pistons for experimental motorcycle engines were also developed for Harley-Davidson vehicles with the CC fabrication carried out at HITCO.

Figure 7. CC automotive piston.

Irr anothrr vtigiiic application, CC is 1)ciilg cvill1latd for iisc on rol a ~ y valvcs in two- anti four-strokc-cyclc sinall riigiiicx Thc CC rot,iiry valvc. tlrivcs off f,hc crankshaft and, as a rc~placcincrit for thc eoiive11t,ioiia1 poppct valvc, it oflcrs wcight savings, qiiict,cr ciigiiic op(wtioi1, arid a rcdiiccd t,orqiir load. As in tlic CC pistmi, 1,tic aclii(waI)l(~ low t t i ( ~ i i r a 1 expansion a i d good st,rciigt,h rrtcntiori a t high tcinpcriiOiircs allow 1 hc closc tolrranccs nccdcd for t21icrot,ary valvc fiiuetion. Drvc~lolmicwtwork in this area is 1)cing condnctcd at thv Ncrtint,cc Corrip:iuy ill Wcstland, Michigan. A very promising area for CC engine components is in diesel engines. Potential benefits that could be derived from CC components have been identified in both adiabatic and low-heat rejection designs. The adiabatic

273

CC Materials and Composites diesel engine is based upon the adiabatic or no-heat-loss process. Essentially, this process involves an insulated combustion chamber from which the high-temperature exhaust gases are passed, first through an insulated duct system and then through turbine wheels, to extract additional energy before finally being expelled. The turbines power the compressor and the crankshaft. A basic diesel engine has approximately 36 percent thermal efficiency and the adiabatic diesel has approximately 48 percent thermal efficiency; this thermal measurement is equivalent to a turbine inlet temperature of over 4000°F (2205°C) for a gas turbine engine. The turbocharger and pistons are CC components. The maximum temperature anticipated, with advanced materials, is approximately 2000°F (1093°C) compared with current maximum diesel temperatures of approximately 600°F (316°C). High-temperature resistant materials such as CC would permit reduced airflow and reduced engine volume. Thermal shockresistant material is required and, because reduced heat transfer is a key consideration, low thermal conductivity CC or shield materials also may be required. Other diesel-engine designers have considered CC to be a viable material for use on advanced engines but have emphasized a low-heatrejection diesel design rather than an adiabatic diesel design. A low-heatrejection diesel engine retains heat within the engine and is a passive means for improving efficiency. The temperatures of interest are on the order of 1600°F (871°C). The low-heat-rejection engine would offer a thermal efficiency of 45 percent; this efficiency is achieved with less complexity than would be required for the adiabatic diesel engine to achieve an anticipated 48 percent thermal efficiency. Components that would benefit from the use of CC would be valves, piston crowns, and, possibly, cylinder liners. Another interesting application would be on the foredeck of the diesel engine, which is currently cast iron. The CC composition and constructions would have to be carefully designed because of the necessity for a low thermal expansion material. The chief benefits of CC are the weight savings and the efficiency accrual that are the result of high-temperature operations. The majority of the uses of CC to date involve relatively low-volume production (i.e., Space Shuttle and rocket motor components) in which performance requirements outweigh production efficiencies. One wellknown use of CC which involves high production, however, is in aircraft wheel brake sets. The B.F. Goodrich Company and HITCO are two of the major producers of these disk-like components. The CC brakes consist of either laminated fabric constructions or chopped carbon fabric moldings. The CC brake rotors weigh 20 percent less than comparable 274

Applications of Carbon-Carbon

steel parts and, because of the higher temperature capabilities of CC (2.5 times the heat capacity of steel), durability is increased, which permits up to 3000 jet aircraft landings compared with approximately 1500 landings for metal rotors (ref. 14). On the Boeing 767 aircraft, using CC yields a weight savings of 871 lb over steel brake systems (ref. 15). The market for CC brakes has also expanded and includes automotive brakes for highperformance cars such as racers as well as CC clutch assemblies. A CC clutch assembly and brake components, which HITCO fabricated, are illustrated in figure 8. The potential for significant growth in these areas is further reinforced by the interest shown by domestic and foreign passenger car manufacturers in utilizing CC in their advanced automotive braking systems.

Figure 8. CC clutch assembly and brakesfor racing cars (HITCO).

The aerospace field continues to be one of the primary areas for use of CC. In addition to the'Space Shuttle that was discussed previously, CC has been successfully used in solid-propellant rocket motor nozzles and exit cones, and as ablative nosetips and heat shields for reentry vehicles. Solid rocket motor gas temperatures on the order of 5400'F (3O0O0C), coupled with near-isotropic thermomechanical loads, have made 3-D CC an attractive material for nozzle throat and exit areas. While 2-D CC exit 275

C C Materials and Composites cone constructions are still being used (fig. 9), 3-D weaving technology, as developed by Aerospatiale and Brochier, S. A., in France, has been licensed and transferred to U S . firms because 3-D CC components are beginning to supplant the weaker 2-D materials. An example of this technology use is the automatically woven, 3-D CC exit cone with an exit diameter of 48 in. which was supplied to Hercules by Aerospatiale and successfully test-fired in 1985 (ref. 6).

Figure 9. Rocket motor 2 - 0 CC exit cone assembly (HITCO).

A relatively new application of CC that is currently being explored is its use as a structural material for advanced spacecraft. Items for CC applications might include thin-wall tubes, angles, and panels for use as booms, trusses, equipment support mounts, and thermal management components. This new interest in CC use is based on the demonstrated fabrication of thin-ply CC shapes and the use of the newer high-modulus graphite (E = 100 msi) fibers as reinforcement materials. Additionally, the high-temperature capabilities of CC coupled with the dimensional stability achievable through the use of high-modulus, negative-thermal expansion fibers make it one of the most resistant materials to high thermal pulse environments. For many space applications, CC does not outgas;

+

276

Applications of Carbon-Carbon therefore, CC does not present a contamination problem for sensitive optical surfaces (ref. 16). A related application was demonstrated, based on achievable dimensional stability and high electrical conductivity. Through a cooperative effort, the Aerospace Corporation, Ford Aerospace Corporation,* and HITCO designed and fabricated a parabolic radio frequency (RF) antenna reflector for satellite communications. This component (fig. 10) was shown at the 33rd International Symposium of the Society for the Advancement of Material and Process Engineering.

Figure 10. CC parabolic RF antenna rejector (HITCO).

Key properties of CC, such as good strength retention at temperatures approaching 5000°F (276OoC), low density, dimensional stability, and complex shape fabricability, make CC an appropriate material for a variety of evolving applications. The use of RF heating in the Tokomak Fusion Reactor requires RF limiters to act as plasma shields to protect the RF launchers. Earlier limiters were made of Poco or ATJ graphite tiles fastened to a water-cooled Inconel structure. Substitution of CC limiters *Now called Space Systems/Loral.

277

CC Materials and Composites because of good structural properties and toughness allows a thin shell design and passive cooling by thermal radiation. The limiters consist of 24 U-shaped segments in a 5- to 6-ft-diameter toroidal arrangement with the base of the U-shaped segments facing inward. The CC limiter is based on a 0-go", 2-D composite using staple fiber PAN yarn as a knit cloth. The width of the U-shaped segment is 21 cm and the leg height is approximately 6 cm with a wall thickness of 1 cm. The CC limiters must function at approximately 3990°F with short-duration spikes of 5000°F to 6000°F (ref. 17). Carbon-carbon is being considered for containers in which nuclear wastes are stored because of the high temperature that might be generated. Good thermal stability also makes CC attractive for use on laser shields to protect space-based satellite systems from the heat of high-powered laser beams in a space defense scenario. In a completely different environment, the compatibility of carbon with body tissues makes CC an interesting bone replacement in areas such as the hip instead of the currently used stainless steel (ref. 14). Carbon-carbon is used in fuel cells in a commercial application that is related to electric power generation. The basic fuel cell system consists of a fuel processor that converts raw material fuel into hydrogen-rich gas, fuel cells that directly convert chemical energy into electrical energy, and a power conditioner to convert the fuel cell dc current into ac current. In practice, many individual fuel cells are connected in series to form fuel cell stacks; the power generator consists of many modules of these stacks. Currently, the most developed system is based on a phosphoric acid electrolyte with CC electrodes and other structural components. The cells operate at 400°F and generate from 200 kW to 11 MW of electrical energy. The CC functions well because of its relatively high thermal and electrical conductivity and its resistance to the fuel cell environment (ref. 18). Carbon-carbon use in glass container forming machines as an asbestos replacement for hot-end glass contact applications illustrates its potential commercial growth. A typical machine arrangement is shown in figure 11, where K-Karb 2-D CC from Kaiser Aerotech is used as pushout pads, stacker bars, ware transfer pads, and machine conveyor ware guides (ref. 19). The CC material showed wear characteristics from 100 to 300 times greater than asbestos for these applications, and because it does not get wet by molten glass and does not require external cooling or frequent replacement, it is a cost-effective replacement for asbestos. Additional commercial applications of K-Karb CC include vanes for rotary vane compressors and vacuum pumps, in which the CC replaces other composite and graphite parts to increase service life; nuts, bolts, and fittings to assemble major graphite elements in vacuum furnaces and increase working life over formerly used amorphic graphite parts; and 2 78

Applications of Carbon-Carbon

Figure 11. Section of glass container forming machine showing CC applications (Kaiser Aerotech).

flat and cylindrical heating elements for hot isostatic presses to increase work life up to 10 times over amorphic graphite, provide consistent resistance numbers to minimize power source adjustments, and allow increased working temperatures to 4000°F (2200°C). Other commercial applications include sintering trays for carbonizing and carbiding furnaces in which toughness of the CC tolerates rough shop usage better than brittle graphites; clutch disks for racing and other high-performance cars to provide high-temperature stability; and planar bearings in which temperatures are too high for Teflon or similar materials and in which graphite suffers from catastrophic failures (R. Jensen, Kaiser Aerotech, personal communication, August 1989). The applications and ultimate markets for CC materials are continually developing for both military and commercial use. New domestic and foreign suppliers have contributed to the acceptance of this unique and versatile material through the use of improved high-modulus and highstrength carbon fibers, the development of high-char organic polymers, improved pitch matrix precursors, new rapid processing CVI technology, a better understanding of fiber-matrix interface phenomenology, and ever-increasing production capabilities. The applications discussed here 279

CC Materials and Composites indicate the variety of roles that CC can fulfill; it is hoped that they will lead to additional future uses.

References 1. Carroll, T. J.; and Connors, D. F.: Final Report for Rapid Densification of 2D Carbon-Carbon Preforms. Textron Specialty Materials Report, Dec. 1989.

2. Curry, Donald M.; Latchem, John W.; and Whisenhunt, Garland B.: Space Shuttle Orbiter Leading Edge Structural Subsystem Development. AIAA-83-0483, Jan. 1983. 3. Curry, Donald M.: Carbon-Carbon Materials Development and Flight Certification Experience From Space Shuttle. Oxidation-Resistant Carbon Carbon Composites for Hypersonic Vehicle Applications, Howard G. Maahs, ed., NASA CP-2051, 1988, pp. 29-50. 4. Martin, Jim: Creating the Platform of the Future NASA-The National Aerospace Plane. Def Sci., vol. 7, no. 9, Sept. 1988, pp. 55-61. 5. Industry Observer Brief. Aviation Week & Space Technol., Aug. 15, 1988, p. 11.

6. Hordonneau, A.; and Grenie, Y.: Multi-Direction in Carbon-Carbon. Aerosp. Compos. & Mater., Autumn 1988, p. 17. 7. Dixon, G. A.: Large Rib-Stiffened Composite Panel. News Release, B.F. Goodrich, May 1989.

Super-Temp

8. Covault, Craig: X-30 Technology Advancing Despite Management Rift. Aviation Week & Space Technol., Mar. 7 , 1988, pp. 36-43. 9. Sheehey, P.: LTV Turbine Rotor Exceeds Test Objectives. Release M88-21, LTV Corp., Aug. 23, 1988.

News

10. Processing Materials and Fabrication Improvements Puts CarbonCarbon Technology Ahead. Adv. Mater., vol. 10, no. 18, Oct. 24, 1988. 11. Marsh, G.: Braving the Heat. Aerosp. Compos. & Mater., vol. 1, no. 4, Summer 1989, pp. 28-32.

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Applications of Carbon-Carbon 12. Taylor, Allan: Fabrication and Performance of Advanced CarbonCarbon Piston Structures. Fiber-Tex 1988, John D. Buckley, ed., NASA CP-3038, 1989, pp. 375-395.

13. Taylor, Allan H.: Lightweight Piston. U S . Patent 4,683,809, Aug. 1987. 14. Holusha, J.: Withstanding 3,000-Degree Heat. The New York Times, NOV. 23, 1988, p. C-6. 15. West, P.: Weight Saving Carbon-Carbon Brakes Gaining Favor on Commercial Aircraft. Adv. Mater., vol. 10, no. 20, Nov. 28, 1988. 16. DeMario, William F.: New World for Aerospace Composites. Aerosp. America, vol. 24, Oct. 1985, pp. 36-40, 42. 17. Labik, G. W.; Bialek, J.; Owens, T. K.; Ritter, R.; and Ulrickson, M.: TFTR Carbon-Carbon Composite RF Limiters. Plasma Physics Lab., Princeton Univ., 1987, pp. 125-128. 18. Rastler, D.: Electric Power Research Institute Fuel Cell Program. EPRZ Tech. Brief, Electric Power Research Inst., Mar. 29, 1988. 19. Report on the Economies Achieved by Replacing Asbestos With K-Karb Carbon-Carbon. Kaiser Aerotech Tech. Bull., 1980.

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