Phase Diagrams in Ceramics

PHASE DIAG~~WS IN ADVANCED CERAMICS

A VOLUME OF THE TREATISE ON MATERIALS SCIENCE AND TECHNOLOGY

TREATISE EDITORS

HERBERT HERMAN Department of Materials Science

and Engineering State University of New York

at Stony Brook Stony Brook, New York

GERNOT KOSTORZ Institut fiir Angewandte Physik ETH-Honggerberg Ziirich, Switzerland

ADVISORY BOARD

M.E. FINE Department of Materials Science Northwestern University Evanston, Illinois

P.B. HIRSCH, FRS Metallurgy and Metal Science

Department Oxford University Oxford, England

A.N. GOLAND DePartment of Physics Brookhaven National Laboratories Upton, New York

J.B. WACHTMAN Department of Ceramic Science/

Engineering Rutgers-The State University New Brunswick, New Jersey

PHASE DIAGRAMS IN

ADVANCED CERAMICS

Edi ted by

A L L E N M. A L P E R

0 SRA M S YL VA NIA, In c. Chemical and Metallurgical Products Towanda, Pennsylvania

J

ACADEMIC PRESS San Diego New York Boston London Sydney Tokyo Toronto

This book is printed on acid-free paper. @

Copyright �9 1995 by ACADEMIC PRESS, INC.

All Rights Reserved.

No part of this publication may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopy, recording, or any information storage and retrieval system, without permission in writing from the publisher.

Academic Press, Inc. A Division of Harcourt Brace & Company 525 B Street, Suite 1900, San Diego, California 92101-4495

United Kingdom Edition published by Academic Press Limited 24-28 Oval Road, London NW1 7DX

Library of Congress Cataloging-in-Publication Data

Phase diagrams in advanced ceramics / edited by Allen M. Alper. p. cm. -- (Treatise on materials science and technology)

Includes bibliographical references and index. ISBN 0-12-341834-8 1. Ceramics. 2. Ceramic materials. 3. Phase diagrams.

I. Alper, Allen., date. II. Series: Treatise on materials science and technology (unnumbered) TP810.5.P48 1995 666-dc20 94-28600

CIP

PRINTED IN THE UNITED STATES OF AMERICA 94 95 96 97 98 99 QW 9 8 7 6 5 4 3 2 1

This book is dedicated to Vincent A. St. Onge who has been a mentor and coach in my business growth and development

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Contents

CONTRIBUTORS PREFACE

Phase Chemistry in the Development of Transparent Polycrystalline Oxides

W. H. Rhodes

I. Introduction II. Yttria

III. Alumina IV. Magnesium Aluminate V. Aluminum Oxynitride Spinel

VI. Lead-Lanthanum-Zirconium-Ti tanate VII. Conclusions

References

1 2

15 23 28 32 38 39

The Use of Phase Diagrams to Predict Alkali Oxide Corrosion of Ceramics

Jesse J. Brown, Jr.

I. Introduction II. Literature Review

III. Predictions of Alkali Corrosion of Alumino-Silicate Refractories Using Phase Diagrams

IV. Comparison of Predicted Na20 Corrosion of Alumino-Silicate Refractories with Experimental Results

V. Conclusions References

43 44

74

79 80 81

vii

viii C O N T E N T S

Application of Phase Diagrams to the Production of Advanced Composites

William B. Johnson and Alan S. Nagelberg

I. Introduction II. Review of the Directed Metal Oxidation Process

III. Thermochemical Considerations of Matrix Formation IV. Phase Equilibria in Carbide/Boride Systems V. Conclusion

References

85 88 95

112 121 122

Use of Phase Diagrams in the Study of Silicon Nitride Ceramics

Tseng-Ying Tien

I. Introduction II. Representation of Silicon Nitride-Metal Oxides Systems

III. Si3Na-Metal Oxide Systems IV. Alloy Design V. Conclusion

References

127 128 131 140 154 155

The Use of Phase Studies in the Development of Whiskers and Whisker-Reinforced Ceramics

Aleksander J. Pyzik and Alan M. Hart

I. Introduction II. Whiskers and Ceramic Matrix Whisker-Reinforced Composites

III. Ceramics Toughened by in Situ Synthesized Whiskers IV. Self-Reinforced Ceramics V. Conclusion

References

157 160 195 199 221 222

Index 227

Contributors

Numbers in parentheses indicate the pages on which the authors' contributions begin.

JEssE J. BRowN, JR. (43), Department of Materials Science and Engineering, Virginia Polytechnic Institute and State University, Blacksburg, Virginia 24061

ALAN M. HART (157), Dow Chemical Company, Central Research and Development, Advanced Ceramics Laboratory, Midland, Michigan 48674

WILLIAM B. JOHNSON (85), Lanxide Corporation, Newark, Delaware 19714

ALAN S. NAGELBERG (85), Lanxide Corporation, Newark, Delaware 19714

ALEKSANDER J. PYZIK (157), Dow Chemical Company, Central Research and Development, Advanced Ceramics Laboratory, Midland, Michigan 48674

W. H. RHODES (1), OSRAM SYLVANIA, Inc., Danvers, Massachusetts 01923

TSENG-YIN6 TIEN (127), Department of Materials Science and Engineering, University of Michigan, Ann Arbor, Michigan 48109

Preface

Since the first four volumes of "Phase Diagrams: Materials Science and Technology ''1 (Alper, Ed.)were published by Academic Press approximately twenty years ago, there has been a revolutionary change in the development and understanding of ceramics materials. Significant ad- vances in the sintering and control of optical properties of ceramics have taken place. An understanding of complex alkali oxides-aluminia-siIica phase relations has been gained, and knowledge of multicomponent silicon nitride-meta| oxides-nitride-carbide systems has been greatly advanced. The ways in which phase diagrams are used to make unusual composites of refractory oxides and nonoxides have undergone revolutionary changes.

This book is composed of some of the leading work in this field.

ALLEN M. ALPER

1Volume I: The Use of Phase Diagrams in Ceramic, Glass, and Metal Technology Volume II: The Use of Phase Diagrams in Metal, Refractory, Ceramic, and Cement

Technology Volume III: The Use of Phase Diagrams in Electronic Materials and Glass Technology Volume IV: The Use of Phase Diagrams in Technical Materials Volume V: Crystal Chemistry, Stoichiometry, Spinodal Decomposition, Properties of

Inorganic Phases

Phase Chemistry in the Development of Transparent Polycrystalline Oxides

W. H. R H O D E S

OSRAM SYLVANIA, Inc. Danvers, Massachusetts 01923

I. II .

I I I .

IV .

V~ V I .

V I I .

I n t r o d u c t i o n . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1

Y t t r i a . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2

A . Y 2 0 3 - L a n t h a n i d e A d d i t i v e s . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4

B. Y 2 0 3 - G r o u p - F o u r A d d i t i v e s . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8

C. Y 2 0 3 - A I 2 0 3 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11

A l u m i n a . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15

A . A I 2 0 3 - M g O . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15

B. A 1 2 0 3 - M g O - Y 2 0 3 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20

M a g n e s i u m A l u m i n a t e . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 23

A . S t o i c h i o m e t r y V a r i a t i o n s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 24

B. M g A 1 2 O a - L i F . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 26

A l u m i n u m O x y n i t r i d e S p i n e l . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 28

L e a d - L a n t h a n u m - Z i r c o n i u m - T i t a n a t e . . . . . . . . . . . . . . . . . . . . . . . . 32

A . P h a s e R e l a t i o n s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 32

B. L i q u i d - P h a s e S i n t e r i n g . . . . . . . . . . . . . . . . . . . . . . . . . . . , . . . 36

C o n c l u s i o n s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 38

R e f e r e n c e s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39

I. Introduction

Polycrystalline oxides have been available for optical applications since the early 1960s when Coble [15] invented translucent A1203. Before this, it was thought that porosity, with its inherent light scattering property, was a necessary consequence of the ceramic fabrication process. Translucent A1203 is a key element in high-pressure sodium lamps manufactured all over the world. Although it is not clear that Coble made extensive use of phase diagrams in his original development, phase relations are important to the success of translucent A1203, and subsequent researchers seeking to sinter other oxides to transparency have found phase relations important to their success in achieving their ultimate goal in sintering: elimina-

Copyright �9 1995 by Academic Press, Inc. All rights of reproduction in any form reserved.

2 W.H. RHODES

tion of residual porosity. For example, in the La203-Y20 3 system, Rhodes [54] used a two-phase field to control grain growth during the pore- removal period of sintering and then shifted to a single-phase field to anneal to transparency.

A number of oxides have been developed for optical applications in which glass, because of refractoriness, chemical compatibility, or limitations on bandwidth, typically cannot compete. These oxides are A1203, MgAl204, ALON, and 5120 3. Although other oxides have been sintered or hot pressed to transparency, these four are considered to be the prime candidates to extend optical applications.

Another important category of optical application uses the electro-optic switching character of perovskites. The system (Pb, La)(Zr, Ti)O 3 has been hot pressed and sintered to transparency by Haertling [26] and Snow [65], respectively. Phase relations are critical to successful fabrication where Pb volatility makes it difficult to retain compositional limits, and compositional variations make possible memory, linear, or quadratic applications.

II. Yttria

Yttria is of interest for optical applications principally because of its high melting point (2464~ and capacity for wideband transmittance (0.23 to 9.5 ~m). This combination of properties makes possible a material that will perform at high temperatures with low emittance and minimal migra- tion of the phonon edge into the 3- to 5-~m band of interest for many infrared applications. Early development of transparent Y203 was directed toward the lamp envelope application in competition with A120 3. Interest in this application fell mostly because of the high cost of powder, but it remains viable for applications requiring the thermodynamic properties of Y203. The infrared window application has received considerable attention throughout the world in the last 10 years. Also, unique x-ray scintillators based on transparent YzO3-Gd203 developed by Greskovich et al. [25] are extremely successful commercially.

The early work of Brissette et al. [12] and Lefever and Matsko [39] demonstrated that pure undoped Y203 could be fabricated to transparency by employing a combination of fine active powders and high-temperature press forging. More recently, Hartnett et al. [34] and Shibatta et al. [63] have taken a similar powder approach with hot isostatic pressing to fabricate transparent complex geometries. Success in these approaches relies not only on powder properties and the application of pressure to

1 TRANSPARENT POLYCRYSTALLINE OXIDES 3

enhance densification kinetics, but also on high purity to prevent precipitation of second phases or coloration from transition or rare-earth ions.

Phase relations become critical to success when pressureless sintering is the fabrication mode chosen to attain transparency. A number of sintering mechanism options are available to ceramists in their quest to affect the densification and grain-growth control necessary to eliminate porosity and achieve transparency. These options include liquid-phase sintering, tran: sient second-phase sintering, and doped solid-state sintering. All require accurate knowledge of solid solution limits, eutectics, and other phase relations governing microstructure development.

o

2400

2200

1800

1400

1000

600

- (7 - - ' - - - - -

, f . _ - -

- i

L a 2 0 3 I

57

C

I

.iZ/ d

Nd203

59

/ r

c

q

H L ' �9

-1

I I Y

B I

/

Ee 0 3 Tb: 93 Ho 2 13 :

s. o, I 62 63 64 65 66 67 68 70

Oxide R 2 0 3

Atomic Number of the Element

FIG. 1. Phase field diagram for the rare-earth oxides. After Foex and Traverse [18]. Reprinted by permission of Societie Francaise de Mineralogie et al Crystallographie.

4 W.H. RHODES

The phase field diagram of Foex and Traverse [18] is useful in evaluat- ing possible sintering aids for Y20 3. Figure 1 shows that the rare-earth oxides are generally found in one of three structures, depending on cation size. The cubic, or C, structure is a distorted fluorite structure having octahedral cation coordination with 32 MO1. 5 groups per unit cell and a large vacancy in the center. Y203 has this structure since it has the same ionic radius (0.892 A) as Ho (0.894), atomic number 67. The lower atomic numbers have larger ionic radii, which result in a monoclinic, or B, structure having either six or seven cationic coordination. Still larger ions form the seven coordinated hexagonal, or A, structure. One can see that additions of larger cations to Y203 would develop extremely complex, interesting, and potentially useful phase diagrams.

A. Y2 0 3-Lanthanide Additives

Rhodes [54] was looking for a noncoloring aliovalent sintering aid for Y203 . Among the rare-earth ions, this restricts additions to La or Gd, because these ions have no electronic transitions in the visible or infrared frequencies. Rhodes's early sintering experiments were compared with the

FIG. 2. Uncontrolled La203-Y203 microstructure resulting from lack of phase diagram knowledge.


i I ' , i i . , , i i , , I | ' ' i |

2 4 0 0 ~ , ~ Liquid

2200 ~,~~,~:~~~ "~ 2o0o . , . r.-'.-:;L-;:

~ H s s . " . , ' " I , , .. " " #sr -

Yss "~ +P P+L Lss 16oo Pi I II@ o .I -

1400 I"30Op'~PI: ~4]:/l ~ --

/ , . I P, Y+P P 1200 I-- I , i , i , . i : - , . , .

Y203 10 30 50 70

Mole %

! 90 La20 3

o

2400

2200

2001]

1800

1600

1400

La203-Y203 L+H

_ (

Liquids

, .t ., Xss ,,F

" " % - ' ~ ~ - ~ / - - , :+.~ r I ~ . . + . - _ ~1,~_ 4,. 4" H+C [

/~4 /"+PII P+C 1 1 i i I I I I I i I ~,J

20 40 60 80 100

Mole% Y203

FIG. 3. Phase diagrams for the La203-Y203 system by (top) Coutures and Foex [17] and (bottom) Mizuno et al. [47]. Reprinted by permission of Academic Press and the Ceramic Society of Japan.

6 W.H. RHODES

incorrect phase diagram of Cassedanne and Forestier [13]. This greatly retarded progress, because early experiments with 10 mole % La203 additions did not show evidence for the predicted liquid phase, but exhibited highly unusual microstructures (such as the one shown in Fig. 2) that were difficult to understand. Finally, a model emerged that can be explained with the aid of the phase diagram of Contures and Foex [17], shown in Fig. 3 (top), which coincidentally was being established at the same time. Mizuno et al. [47] were also working on the LazO3-Y203 system, and they found some differences (Fig. 3, bottom).

A composition of 9 mole % La203 was typically prepared by coprecipi- tating the oxalates. After calcining to form the oxides and milling to reduce agglomerate size, isostatic pressing was used to form near-net-shape

FIG. 4. A 0.09 La20 3 �9 0.91 Y203 pore-free microstructure achieved by sintering in the two-phase (21500(;) and annealing in the single-phase (1900~ region of the phase diagram.


compacts. The compacts were heated at a rate that was rapid enough to prevent pore entrapment into the two-phase C solid solution plus H solid solution field. A sintering temperature, typically 2150~ was chosen to assure approximately 25 vol % H phase. Grain growth was markedly retarded by the second phase and followed a one-third time constant kinetics predicted for grain growth limited by particles of a second phase coalescing by lattice diffusion. Pores were thought to remain attached to the slow-moving grain boundaries and annihilate by a classic vacancy solid-state diffusion process. The temperature was lowered to approximately 1900~ which is well inside the C solid solution field. Holding at this temperature dissolved the H phase by a slow diffusion process. Fortunately, the phase change was not displacive nor accompanied by a volume change that generated porosity. Instead, the pore-free microstructure shown in Fig. 4 was achieved.

Ceramic scintillators developed by Greskovich et al. [25] contained large concentrations of lanthanides based on Y203-Gd203 solid solutions

2300

2200 -

o

2100

[... 2000 -- Cubic Solid Solution [

/ I Cs s + Bss

1900 - I I I I I I

Y203 10 20 30 40 50 60

Mole % Gd20 3

FIG. 5. Tentative Y203-Gd203 phase diagram.

8 W.H. RHODES

and one or more rare-earth activator oxides, such as Eu203, Nb203, Yb203, DY203, Pr203, and Tb203. The Gd203 is essential in absorbing the x rays, while the activators undergo electronic transitions that convert the energy to visible wavelengths for photodiode detection. This polycrystalline ceramic may also contain a group-four sintering aid, which is discussed in the next section. The concentrations of lanthanides added to Y203 appear to adhere deliberately to the general trends suggested by Fig. 1 for C solid solution limits. In previously unpublished work, Rhodes defined the Y203-rich end of the Y203-Gd203 diagram (Fig. 5). The work of Greskovich et al. [25] shows a cubic/monoclinic phase boundary at 50 mole % Gd203, in agreement with Figure 5. It is interesting to note, based on similar phase diagrams, that the transient second solid-phase sintering mechanism discussed for La203 might also be applicable for Gd203 additions. Greskovich [23] reported that the heavy doping in scintillator compositions can act as grain-growth retardants, allowing transparency to be achieved without group-four sintering aids. The sintering mechanism operative in this case must be solute drag, since these aliovalent substitutions would not create defects affecting diffusion rates.

B. Ye O 3-Gr~176 Additives

The classic work of Anderson [2] in developing Yttralox (General Electric, Schenectady, N.Y.) originated from research in the Y203-ZrO2 system to develop better ionic conductors for fuel cells. After exploring the ZrO2-rich end of the system that resulted in the good 0 -2 conductors of the fluorite structure, he decided to check the fluorite structure at the Y203-rich end. Quite unexpectedly, a number of the very first samples in this series sintered to transparency. This led to a research concentration in transparent Y203, and to the discovery that group-four oxides, HfO 2 and ThO 2, also behaved as effective sintering aids. ThO 2 was the most fully studied and characterized additive.

The ThO2-Y20 3 phase diagram after Sibieude and Foex [64], of Fig. 6 shows a solid solution range up to 15.5 mole % at 2000~ Jorgensen and Anderson [37] performed precision lattice parameter measurements in this system in an attempt to elucidate the sintering mechanism. Their results, shown in Fig. 7, agree quite well with the published phase diagram at 2000~ The lower temperature results suggest a less steep temperature dependence on the solubility limit. A new line has been added to Fig. 6, incorporating the Jorgensen and Anderson measurements and the results of Greskovich and Chernoch [24]. Transparent ceramics were obtained with as little as 5 mole % ThO2; additions of 7 to 10 mole % were


3200

2 8 0 0

o.., 2 4 0 0

2000

1600

I I I I

- ",~,- " ~ H + Liq. T h O 2 (ss) ~ ~ m . . . . . . . . . _ ~

ThO 2 (ss) + H

Y205 (ss)

ThO 2

I I I I 20 40 60 80

Mole %

Y203

FIG. 6. Y203-ThO2 phase diagram. After Sibieude and Foex [64]. The dashed line incorporates the results of Jorgensen and Anderson [37] and Greskovich and Chernoc [24]. Reprinted by permission of Eisevier Science Publishers BV.

preferred. These compositions were clearly in the solid solution range, so the research centered on delineating the sintering mechanism. Various methods, including autoradiography and microhardness profiles, were explored in search of a possible solute grain-boundary segregation. Micro- hardness profiles consisted of a series of 5-g load indentations across a grain boundary in a protected environment to prevent stress corrosion effects. The results showed a dramatic hardness increase across the grain boundary (Fig. 8). They also showed that the bulk underwent solid solution hardening, which led them to conclude that ThO 2 segregated to the grain boundary. The autoradiography also supported the solute segregation mechanism for grain-growth retardation and the resulting achievement of theoretically dense sintered ThOz-doped Y203.

Greskovich and Chernoch (24)followed this work with a program to develop polycrystalline NdzO3-doped lasers based on a ThOz-Y203 matrix. This was a very ambitious program, as no one had previously at-

10 w . H . R H O D E S

10.68

10.67

10.66,

10.65

~. 10.64

~ 10.63

10.62

10.61

I I l l l l l l l l I

I I 92ooo~ �9 _ ~ 1 9 0 0 o c _

, f ' _ 1800~ _

10.60 i I ! ,, I I I I I I I I I ~-, 0 2 4 6 8 10 12 14 16 18 20 22 24

ThO 2 (mole %)

FIG. 7. Lattice parameters a function of ThO 2 in Y203 . After Jorgensen and Anderson

[37]. Repr in ted by permission of the American Ceramic Society.

r

=

700 -t

650 -

6 0 0 -

550 -

500 -

450 -

400

350

I I I I

! I i I I 60 50 40 30 20

i ' i i i i i I

I I I I I l ~ 10 0 10 20 30 40 50 60

Distance from B o u n d a r y (llm)

FIG. 8. Y203-ThO2 grain-boundary microhardness plot. After Jorgensen and Anderson

[37]. Repr in ted by permission of the American Ceramic Society.


FIG. 9. High-density microstructure achieved with 10 mole % ThO2-doped Y203 . After Greskovich and Chernoch [24].

tempted to fabricate a polycrystalline material of this optical quality. The result, however, was laser efficiencies that were ~ 20% that of Nd glass lasers. Extremely low residual porosities (< 10-4%)were demonstrated, and the excellent microstructure shown in Fig. 9 is evidence of that fact. The principal scattering defect remaining, which affected laser efficiencies, was thought to be minor refractive index variations resulting from compositional gradients within the solid solution. These could result from the segregation effects noted earlier. If this were true, the very chemical effects permitting the attainment of low porosities would in fact result in the defect inhibiting application as a polycrystalline laser. Segregation and compositional inhomogeneities are probably quite common in most polycrystalline optical ceramics.

C. Y203-AI203

The early attempts of Rhodes and Reid [55] to sinter transparent lamp envelopes began with an attempt to sinter pure Y203 . It had been learned

12 W.H. RHODES

that group-four sintering aids lowered Y203's stability in low oxygen partial pressure environments such as those present in high-pressure sodium lamps, so the approach centered on controlling pure Y203's powder particle and agglomerate size through synthesis and deagglomeration techniques. One deagglomeration procedure was to dry ball mill using the highest purity AI2O 3 jar mill and balls available. In contrast to results of earlier work, these sintered samples possessed a distinct grain-boundary phase. Furthermore, the samples had a higher sintered density than had heretofore been achieved. Microprobe results confirmed that they had unintentionally "contaminated" the powder with an excellent sintering aid. An extensive program was launched to take advantage of this acciden- tal finding.

The most recent Y203-A1203 phase diagram by Noguchi and Mizuno [50] Fig. 10, shows a solid solution region up to 4 mole % A120 3 in Y203 at 1940~ This portion of the diagram is dashed to indicate some uncertainty. The initial work to use A1203 deliberately as a sintering aid was based on the theory that it should be possible to sinter in the liquid

2 4 0 0 I I I I I

I Liquid 'I

2200 I,

1Y203 s~ ! A120 3 ss + Liq. �9 Liq. [ + ~ _

2000 3:5 + Li . 1:1 + Liq. 2:1

Id ?..l ;..l r [ ~ " A120 3 ss

1600 [ 4 ~e~

AI20 3 ss + 3:5

1400

I

A120 3 20 40

3:5 + 2:1 2:1 + Y203 ss

I

1:1 60 80 Y203

Mole %

FIG. 10. Y203-A1203 phase diagram. After Noguchi and Mizuno [50]. Reprinted by permission of the American Ceramic Society.

1 T R A N S P A R E N T P O L Y C R Y S T A L L I N E OXIDES 13

plus Y203 field and then drive the A1203 back into solid solution by annealing at a lower temperature, say, 1940~ If successful, this would qualify as a transient liquid-phase sintering technique.

An extensive research program was carried out with A1203 additions from 0.02 to 10.0 mole % and sintering temperatures from 1800 to 2200~ in a H 2 atmosphere. Numerous attempts were made to incorporate three temperature holds. Samples were held at temperatures ranging from 1700 to 1900~ to allow diffusion enough time to form the solid solution; then sintered above the 1940~ eutectic over a range from 2000 to 2200~ and annealed for 1 to 8 h in the range of 1800 to 1940~ to drive the eutectic

FIG. 11. Translucent microstructure achieved with 0.5 mole % AI203 and three-temperature sintering cycle in an attempt to minimize the second phase.

14 W.H. RHODES

phase back into solution. There was never any evidence that the two holds on either end of the sintering cycle formed a solid solution. The only effect of the lower temperature hold at the end of the cycle was to crystallize the grain-boundary liquid phase. Figure 11 shows the microstructure of a sample given a three-temperature sintering cycle. The residual grain- boundary phase is evident even though the sample's starting composition contained 0.5 mole % A120 3. The effectiveness of liquid-phase sintering in pore elimination is also shown. Extensive studies on the effect of powder properties on sintered optical properties were discussed by Palilla et al. [52] in conjunction with the development of this material as a lamp envelope.

The lamp envelope development work demonstrated that the solubility limit may be on the order of 100 ppm, supporting the earlier AlzO3-Y203 phase diagram of Toropov et al. [67], which is shown in Fig. 12. The YzO3-rich end of the diagram shows no solubility for AI203. The cubic- to-hexagonal transition at ~ 2200~ for pure Y20 3 is not accounted for in Fig. 12. Another major difference between Fig. 12 and Fig. 10 is whether

Liquid 2020~ L + 3:5

2400 ~ ~ ~ , , , . . ~ ~ ,

- \ Liquid 12~,1~�9 5 ~ I- . . . . "! . . . . . 1865~

• Liq. + \ \2020s~ ~- - .... A

2000--" " \ ~ L i q . + l : l ~ , . . . . / [ �9 q_, 1. ~ i 193~C LIq. + ,5:~ / I

~194~176176 / . . . . A [ ['~ 1 tJ 2:1 + 1:1 ~ " " ~ ~ / " Llq. + (X AI2U 3 I

I ! ' , 1835~ ,,"! 1' "~ / " ]

l 2:1 r ~[ 3:5 I 1600,. - "1 +

} , I , I I I I ,

Y203 20 40 60 80 AI20 3

Mole %

FIG. 12. Y203-AI203 phase diagram. After Toropov et al. [67]. Reprinted by permission of the Consultants Bureau.


or not the 1:1 compound melts congruently. The A1203-rich end of Fig. 10 shows extensive solubility of Y203 in A1203. This is discussed more fully in Section III, but solubility is clearly in the range of 100 ppm, as is the case for the Y203 end, again supporting the diagram shown in Fig. 12.

III. Alumina

The prospect and potential for the entire field of polycrystalline oxides for optical applications was opened by Coble [15] with his invention of translucent A1203. Background to this important industrial discovery was the fact that low-pressure Na discharge lamps required special two-ply borate-rich glass envelopes to prevent the Na from attacking SiO 2 in any glass at a life limiting rate. However, high-pressure Na discharge experiments in sapphire tubing showed considerable resistance to chemical attack, and the improved color emission from the high-pressure Na discharge was attractive to the lamp industry. Consequently, there was a strong motivation for developing an economical high-pressure sodium discharge lamp envelope. With the acievement of the translucent A1203 envelope and compatible sealing technology, a new discharge lamp type was born. These lamps are manufactured throughout the world at an annual production estimated to be 3 x 107 units. A new application for translucent A1203 is "clear" orthodontic braces, which are more aestheti- cally pleasing than metallic braces.

A. A12 0 3-MgO

The story is told that the sintering aid MgO was discovered inadver- tently in the course of sintering studies on A1203. One particular powder lot was noted to have unusual sintering and grain-growth behavior. When the supplier was queried about the difference in this lot, he replied, "Oh yes, the powder lot was accidentally contaminated by MgO, but we won't let this happen again." Coble, however, had the right training and percep- tion to realize he was onto an important discovery. Careful sintering and grain-growth kinetics studies by Coble [15] revealed that the presence of MgO enhanced sintering by 1% to 2%. Meanwhile, grain-growth kinetics were unaffected prior to the onset of discontinuous grain growth, which occurred at about 99% density in the MgO-free specimens, but did not occur in the MgO-doped A1203 . Remarkably, MgO allowed normal grain growth to proceed right up to the achievement of theoretical density. Rhodes and Wei [56] showed (Fig. 13) the microstructure that is typical

FIG. 13. Sintered microstructure of AI203. (a) Low-porosity high-transmittance resulting from MgO sintering aid and start-of-the-art sintering process. (b) Pore entrapment resulting in opaque A1203 when MgO is omitted. Reprinted with permission from Rhodes and Wei [56]. Pergamon Press Ltd.

16


for the translucent A1203 employed in lamp envelopes, and compared that with the AI203 microstructures having entrapped porosity typical in all ceramics prior to 1961.

Explaining the mechanistic role of MgO in the sintering of A120 has been a popular and controversial research topic for the last 30 years. Bennison and Harmer [8] reviewed this topic and found that there was general agreement that MgO decreased and grain-boundary mobility by a solid solution pinning mechanism. This reduced the pore-boundary tendency for separation (pore entrapment within grains, as shown in Fig. 13(b)) by a factor of 25. MgO also increases the densification rate by a factor of 3 and the grain-growth rate by a factor of 2.5. Further, there is evidence that MgO increases surface diffusivity, thereby aiding in pore mobility, which keeps the pores on the boundaries until they are annihilated by solid-state diffusion. Sung et al. [66] have recently found direct evidence for MgO segregation to the surface of pores. Overall, MgO is a microstructure stabilizer against green density variations, which might otherwise lead to porosity or other inhomogeneities. It appears that MgO is unique among sintering aids in that it performs multiple roles, all

o

11

Im

/ "~'.'- .- . . . . . Liquid [

2 0 0 0 - 14 Spinel + Liquid "" - .~ " - . '~1

�9 12~)rundum + // 1 8 0 0 - " [] -

S p ~ e ~ " Liquid

1 6 0 0

1400 - [ ]

o 1 2 0 0 - ,

{ Spinel + Corundum I 0 0 0 -~ a

#

8 0 0 [] [ ]

6 0 0

I I I I o37o-14

MgO.A120 3 60 70 80 90 A i 2 0 3

M o l e % A 1 2 0 3

FIG. 14. MgAI204-AI203 phase diagram. After Roy et al. [59]. Reprinted by permission of American Journal of Science.

18 W. H. RHODES

Temperature (~

O

t_.

,q.

•

~

~ J . ,m

J O

100

10

1.0 I 4.6

1830 1730 1630 1530 I I I I

I o3~13

4.8 5.0 5.2 5.4 5.6

1/T x 104 (K)

0.1

0.01

0.001

, I =

~

.e

O

, J

FIG. 15. Solid solubility limit for Mg/A1 [61] with conversion to weight percent MgO added to A120 3.

favorable for the achievement of theoretically dense translucent A120 3 articles.

The A l z O 3 - M g A l z O 4 phase diagram by Roy et al. [59] shows essentially no solid solution of MgO in Al20 3 (Fig. 14). This is appropriate since the solubility has been reported by Roy and Coble [61] to be in the parts per million range. Roy and CoNe reported values in terms of Mg/A! ratios measured in vacuum by equilibrating A120 3 with 0.15 MgO. 0.85 A120 3 through a vapor transport mechanism. Their results are reported in Fig. 15 along with a conversion to weight percent. This conversion should prove useful since most formulations in the literature are reported in this unit. They also measured TiO x and MgTiO 3 solubility and found that both were higher than MgO alone. They rationalized that the Ti +4 charge compensates for the Mg +2 in the A120 3 lattice and increases its solubility.


FIG. 16. Transmission electron micrograph of grain-boundary precipitate in translucent A1203 with EDS, energy dispersive spectroscopy, identification of Mg, A1, and O and convergent beam electron diffraction pattern of the [111] zone axis showing presence of MgA1204.

This principle was invoked by Charles et al. [14] in their patent on ZrO 2 + MgO additions to A120 3. where increased solubility was claimed.

The solid solution limit is an issue in the development of translucent A120 3 because MgAI204 precipitates act as light scattering centers due to their index of refraction difference with A120 3, and detract from both total and spectral transmission. (Pores with an index of 1.0 are much greater scattering centers.) Figure 16 illustrates a typical precipitate of MgAI204 in A120 3 containing 0.08 wt % of MgO in the starting composition followed by sintering at 1900~ in H 2. According to Fig. 15, this composition should have been in the solid solution field, and precipitation should be absent. The cooling rate on this sample was 35~ which is quite rapid, but the possibility exists that MgAI204 precipitated during cooling. The other possibility is that the data in Fig. 15 are incorrect, and the solubility limits are much lower than oridinally believed. This is particularly intriguing because Roy and Coble [61] saw a lattice parameter

20 W.H. RHODES

shift in samples heat treated in H2, suggesting that the solubility could be larger in H 2 than vacuum, the environment for their experiments.

B. AI203-MgO-Y203

Y203 has been used in combination with MgO as a sintering aid for translucent A1203 by a number of manufacturers. The principal advantage of incorporating Y203 is that sintering temperatures and times are lowered. The main disadvantage is that grain growth is more difficult to control. Rossi and Burke [58] studied the sintering and grain-growth behavior of A1203 co-doped with either MgO or Y203 and with MgO and Y203 together. Y203 alone was shown to yield a 97% dense microstructure with 60- to 80-~m pore clusters in the center of 150-/zm grains together with large pores at grain boundaries. The microstructure evolved first by normal and then by rapid secondary grain growth sweeping past

FIG. 17. One of 20 ternary eutectic islands in the AI203-Y203-MgO system used to estimate composition. After Bosomworth et al. [10].

1 T R A N S P A R E N T POLYCRYSTALLINE OXIDES 21

Ai20 3

19195~00,//"

11850~ ,~,~1 oc, . / " \ ne,

Y A ~

FIG. 18. Ternary eutectic diagram in the A1203-Y203-MgO system. After Bosomworth et al. [10].

porosity. This was followed by a period of normal growth assisted by the liquid phase, but at a rate that was compatible for pore elimination and the collection of porosity at grain boundaries. When MgO was included along with Y203, an altogether different microstructure was achieved. Porosity was nearly eliminated, but the grain size was bimodal and much larger than when MgO alone was employed.

Bosomworth et al. [10] studied the MgO and Y203 additive system. By studying compositions much higher in additive than normally employed in

FIG. 19. Microstructure of Y203 and MgO-doped A1203 showing bimodal grain size distribution as a result of the liquid phase present during sintering.

22 W.H. RHODES

practice, they found that the liquid phase was present as a ternary eutectic liquid, as shown in Fig. 17. By analyzing a number of these eutectic islands, they determined that the eutectic composition was 65.5 wt % Y3A15O12, 30.3 wt % A1203, and 4.2 wt % MgAI204. The ternary diagram is shown in Fig. 18. They performed a series of melting point tests as well as high-temperature DTA, differential thermal analysis, runs where the eutectic composition was blended with 70 wt % A120 3. They found a eutectic temperature of 1794~ in the AlzO3-Y20 3 system and 1761~ in the A12 O 3-Y2 O 3-MgO system.

Figure 19 shows a microstructure of A120 3 sintered with 0.05 wt % Y203 and 0.05 wt % MgO above the ternary eutectic temperature. The lack of entrapped porosity reported by Rossi and Burke [58] is confirmed. The microstructure has a distinct shift in grain size distribution if one compares this structure with Fig. 13. It is hypothesized that MgO remains an effective grain growth inhibitor in the presence of the ternary eutectic

AI

Y

eV

Y f~

FIG. 20. Transmission electron micrograph of grain-boundary precipitate in translucent AI20 3 with EDS identification of Y, AI, and O and convergent beam electron diffraction pattern of the [001] zone axis showing the presence of Y3AIsO12.


liquid phase. This condition remains throughout the important pore- removal process. As sintering time continues, MgO can no longer retard grain growth against the rapid liquid-phase assisted grain-boundary diffusion. Secondary grain growth occurs as isolated incidents. These isolated grains continue to grow until they intersect, leaving only small patches of the original first-generation grains. The effective sintering temperature is lowered about 50~ over compositions using MgO alone. Sintering times are also shorter.

The discussion in Section II.C on the YzO3-rich end of the Y203-A1203 diagram demonstrated that there is ~ 100 ppm solid solubility of A1203 in Y203. Figure 10 suggests there is ~ 10 mole % solid solubility of Y203 in A1203, while Fig. 12 shows no solubility. Analytical transmission electron microscopy detected Y3A150~2 precipitates such as those shown in Fig. 20 in samples having a starting composition of 100 ppm Y203 or more. Precipitation could have occurred during the 35~ cooling, but this suggests that the solid solubility limit is near 75 ppm in the ~ 1800~ sintering range.

Translucent A1203 from the People's Republic of China contains La203 in addition to MgO and Y203. A further depression in the eutectic temperature is expected, which probably lowers the sintering temperature even further. The grain size distribution is approximately the same as that shown in Fig. 19, but a higher concentration of grain-boundary phases is present.

IV. Magnesium Aluminate

Transparent magnesium aluminate spinel has been examined for potential applications as lamp envelopes, watch lenses, bullet-proof windows, pressure vessel windows, waveguides, optical components, and infrared windows. The hardness of spinel is 16.1 GPa, which is only 18% less than sapphire. This results in excellent resistance to scratching, dust, and high-velocity water droplets, and assists in defeating projectiles. Spinel has a cubic crystal structure, which means single crystals have no optical anisotropy, and single-phase pore-free polycrystals do not scatter light at grain boundaries.

The MgO-AI203 phase diagram was first reported by Rankin and Merwin [53] in their study of the CaO-MgO-A120 3 system. They showed no spinel stability beyond stoichiometry. Subsequently, Roy et al. [59] and Alper et al. [1] examined the MgAlzO4-A12O 3 system and the MgO-MgAI204 system, respectively. Their diagrams, reproduced in

24 W.H. RHODES

2825

2600

2400

2200 Spinel s.s. + Liquid

2105~ 2000

" 7 " "

1600 f ~1500~ Periclase + Spinel

1400 I I ' ' ..... ' 100 90 80 Weight % MgO 50 40 30

0 10 20 Weight % Ai203 50 60 70

FIG. 21. MgO-MgAl204 phase diagram. After Alper et al. [1]. Reprinted by permission of the American Ceramic Society.

Figs 14 and 21, show a substantial range of stability beyond stoichiometry. Nonstoichiometry on the AlzO3-rich side is more substantial, and slight variations among investigators on the range of stability and the pressure dependence have been compared by Mysen [48].

A. Stoichiometry Variations

Bagley [4] studied the sintering of transparent spinel in a hydrogen atmosphere on the MgO-rich side of stoichiometry. Sintering was conducted at 1600 and 1850~ for 7 h, with the results shown in Fig. 22. Optimum transparencies were obtained with only 0.8 wt % excess MgO. As shown in Fig. 22, compositions out to 28.9 wt % MgO at 1600~ and 34.5 wt % MgO at 1850~ should have been within the solid solution field. It is interesting to note that the peak transparency occurs at the same composition for both temperatures. This suggests that the operative sintering mechanism is not related to solid solution limits, but rather a synergis- tic role for MgO similar to that played in the sintering of translucent A1203, as discussed in Section III. Another possibility is that the solubility of the MgO-rich side is not as wide as that shown in Fig. 21.

Sintering studies on the AlzO3-rich side of stoichiometric spinel were carried out by Bailey and Russell [6]. Transparent ceramics were not the objective of this study, but several findings were noteworthy. Compositions from 71.67 wt % A120 3 (stoichiometric spinel) to 93 wt % A120 3 were


50

4O

30

[...

2O

i �9 i �9

, ,, 1850~ ' Solid "', ' Solution "-. t mit ;

%�9

-x,

1600oc ".

Solid Solution Limit

" l

10

i v Stoichiometric MgAI204

I ~ r i J i I I I t j o + 28 29 30 31 32 33 34 35 36

Wt. % MgO in MgO'AI20 3 Compositions

Fic. 22. Transmittance of sintered MgAI204 versus wt % MgO showing maxima that do not correlate with phase boundary. After Bagley [4] .

sintered in the 1640 to 1685~ range. Comparisons of these compositions with Fig. 14 shows that this study cuts across the solid solution field into the two-phase MgA1204-A120 3 field. It was noted that the sintering temperature to reach a given density increased monotonically across the solid solution field to the phase boundary and then decreased again in the two-phase field. The highest densities (98.2%)were achieved with the highest A1203 compositions. Figure 23 shows the grain size versus composition curve from this study. This curve shows that the largest grain sizes were realized right at the phase boundary. The low density at this composition is probably due to exaggerated grain growth entrapping porosity within grains. Bagley and Bowen [5] studied grain growth in spinel at the solid solution limits. They observed rapid growth at the phase

26 W.H. RHODES

100

80

4o

20

70 75 80 85 90 95

A!203 (%)

FIG. 23. Average grain size versus Al20 3 content showing marked grain size decrease in the Al20 3 + MgA120 4 field. After Bailey and Russell [6]. Reprinted by permission of the Institute of Ceramics.

boundaries and referred to this phenomena as intrinsic grain growth. This same type of growth was observed by Rhodes [54] in the La20 3 system (see Section II and Fig. 2). The high densities in the two-phase MgAlzO4-AI20 3 phase field may have been due to the grain growth retarding effect of the Al20 3 phase allowing porosity to be annihilated at grain boundaries. Note in Fig. 23 that the grain sizes are significantly smaller in the two-phase samples than in the smallest single-phase sam- pies. It is clear based on the evidence presented that control of grain growth is difficult and key to the achievement of transparent spinel.

B. MgA1204-LiF

Atlas [3] reported the effectiveness of certain lithium compounds, particularly halides, in promoting densification of MgO. This finding was exploited by Benecke et al. [7], where LiF was the sintering aid used to hot press transparent MgO. This same technology has been applied to spinel by Sellers and Roy [62] and has continued at Coors Ceramics with a comprehensive patent by Roy and Hastert [60]. The consolidation technology still centers on hot pressing with or without a subsequent hot isostatic pressing step. In the case of MgO, LiF remains as a grain boundary phase after pressing, and the product is translucent. The sample is heated slowly in air to about 1100~ and held for many hours to affect


FIG. 24. Microstructure of hot-pressed transparent MgAI204. After Roy and Hastert [60].

evaporation of LiF and render the product transparent. The discussion on the use of LiF with MgAI20 4 does not mention a postpressing anneal. Presumably, the starting LiF concentration and pressing cycle are carefully controlled to end the pressing cycle with little or no second phase. The other possibility is that there is sufficient solid solubility of LiF in the spinel system to avoid second-phase scattering. The author believes the former argument is more likely. Figure 24 shows the microstructure of transparent spinel fabricated by this process.

Oda et al. [51] and later Matsui and Takahashi [43] describe sintering with the aid of LiF. The level of transparency is not as high as the

28 W.H. RHODES

hot-pressed product, but nevertheless may be sufficient for applications such as arc tubes for discharge lamps. They prefer to operate on the A1203-rich side of stoichiometry. They state that exaggerated grain growth inhibits the achievement of transparency with compositions lower than 52 mole % A120 3. Compositions greater than 70 mole % A120 3 were in danger of forming A120 3 precipitates, depending on the sintering temperatures. Based on their description, it is clear that evaporation of LiF is a key element of their process. It would appear that liquid-phase sintering is the operative sintering mechanism, but control over the concentration of LiF present during the various stages of the cycle is essential. A typical starting composition was 0.2 wt %, with only 0.05 wt % remaining after completion. In a manner similar to that described earlier for hot pressing, enough LiF was vaporized during sintering to limit light scattering from residual grain-boundary phases.

Numerous additional studies have been performed to achieve transparent spinel, including vapor-phase transport by Navias [49], hot pressing by Ho [36] arc melting by Gatti and Noone [20], press forging with SiO 2 and Li20 additives by Rhodes et al. [57], sintering with CaO additive by Bratton [11], sintering pure or Y203 doped by Hing [35], fusion casting by Gentilman [21], forging by Maguire and Gentilman [41], hot pressing by Shibata et al. [63], and hot isostatic pressing with or without MgO additions by Boch et al. [9].

V. Aluminum Oxynitride Spinel

Aluminum oxynitride spinel has been given the acronym ALON. This material has an inverse spinel structure with vacancies on the cation site. As with other high-purity, single-phase cubic polycrystalline oxides, pore- free bodies are transparent. The physical properties of ALON are very similar to magnesium aluminate spinel except for a higher hardness (19.1 GPa) and slightly lower thermal expansion coefficient. The principal application to date is for infrared windows, although many of the other applications discussed for this class of materials may be possible. The high-temperature stability in a low oxygen pressure situation does not appear to be sufficient for the discharge lamp envelope application, however.

ALON is a solid solution region on the binary line of A1N-AI20 3, which is in reality a pseudo binary in the A1-O-N system. The first phase diagram to include ALON was reported by Lejus [40], whereas the first


report of the existence of this cubic spinel was by Yamaguchi and Yanagida [69]. A major effort to synthesize, sinter, and understand the system was conducted by McCauley and Corbin [45] and has been recently reviewed by Corbin [16]. A phase diagram was drawn for this system by McCauley and Corbin [46], with the assumption that the atmosphere was flowing N 2. Figure 25 shows this interesting diagram, which contains vapor above solid A1N and liquid above solid A120 3. ALON is shown to melt congruently at 35.7 mole % A1N and 2165~ McCauley [44] derived a constant anion model for describing the spinel structure centered around the composition 5 A1N �9 9 A120 3 rather than the stoichiometric A1303N composition. This model assumes that A1 +3 occupies both octahedral and tetrahedral sites, and A1 vacancies form to preserve charge neutrality as the O / N ratio

3000 3000

2500 por 2500 ~.~

2400 2400

2,s 2 1 R + V ~ / / 1. 2300 - AIN ; ; + : c V / + - 2300

+ -2225 ~ L ALON 2200 - 27R 21 + L " ~~216S~176 - 2200 O ~

12H + L ~ " - ] "~X~ ~ I'= ~'x - - 2 o s s o c r / ,,'~ f f ~ -

' 12H + A L O N r , f ~ |

- 21R + ALON [ |-*-1 / t + =,%03 -

27R + ALON I [ L + ff.,-Al203

.--I - - -

AIN + ALON ALON + ~-AI203

27D~ 21R 12H 1 ~ AIN + 0~'Ai203 ALON (V) ~' I I I I + I i ~ I . . . . .

2900 2900

2 8 0 0 2 8 0 0

2700 2700

2600 2600

2100 2 1 0 0

2000 2000

1900 - 1900

1800 - 1800

1700 - 1700

100 90 80 70 60 50 40 30 20 10 0

AIN M o l e % AIN A120 3

FIG. 25. A I N - A I 2 0 3 p h a s e d i a g r a m . A f t e r M c C a u l e y and C o r b i n [46]. R e p r i n t e d by

p e r m i s s i o n of M a r t i n u s Ni jhof f Pub l i she r s .

30 W. H. RHODES

changes within the solid solution field. The anion lattice remains constant, with 32 anions per unit cell. The solid solution limits are not well defined according to the workers in this field.

McCauley and Corbin [45] reactively sintered A1N and A120 3 milled powders forming ALON in situ. They explored many combinations of processing parameters. As shown in Fig. 25, liquid phases are found on either side of a fairly narrow solid solution field at high temperatures. Many of the microstructures reported in their papers show second phases, which they believe were liquid at the processing temperature. Single-phase structures were obtained by careful control of impurities and composition. Sintering at 1975~ resulted in considerable porosity entrapped within grains, but a slightly higher temperature of 2025~ resulted in much less porosity and the first transparent ALON.

Later, Hartnett et al. [32] teamed up with Corbin and McCauley to perfect the processing of transparent ALON. One key was to prereact the starting powders, forming an ALON powder with a narrow > 0.5 < 0.5-/xm particle size powder. All workers in this field have found control of agglomerates to be critical to achieving homogeneous high-transparency ceramics. Because high-purity A1N powders are difficult to obtain and high-purity A120 3 is readily available, investigators such as Martin and Cales [42] have obtained their starting powder by the reaction

A12O3(s) + C(s ) + N2(g ) >_ ALON(s ) + CO~g)

They discussed the use of 7 2 0 3 as a sintering aid in the 0.05 to 0.5 wt % range. Low Y203 contents resulted in heterogeneous microstructures with entrapped porosity, whereas the higher concentrations promoted higher densities and mainly intragrain porosity. Sintering was not successful in achieving transparent samples, but hot pressing followed by long anneals at 1940~ did show some promise. There was no discussion of the role of Y203 in sintering. Based on the experience of this additive with A120 3 (see Section III.B), one would speculate that a liquid phase forms, promoting densification by liquid phase sintering. On cooling, this phase would probably form the Y3A15012 phase. The heterogeneous structures discussed by these workers could result from low concentrations of liquid phase due to incomplete grain-boundary wetting.

Hartnett and Gentilman [31] have perfected the processing of transparent ALON, achieving the microstructure shown in Fig. 26. Little residue porosity remains. The microstructure appears equiaxed, with grains on the order of 2 /xm, suggesting that hot pressing or hot isostatic pressing was employed. Y203 (and La20 3) plus B sintering aids, which apparently function as transient liquid phase sintering aids, are discussed by

1 T R A N S P A R E N T P O L Y C R Y S T A L L I N E O X I D E S 31

FIG. 26. Microstructure of ALON. (a) Small equiaxed-grain, pore-free structure. (b) Stacking fault type defect found occasionally within grains.

32 W.H. RHODES

Hartnett et al. [33]. Little or no residual second phase remains after sinter although stacking fault defects were detected in ~ 5% of the grains (Figure 26b). Large domes and plates have been achieved with a very high degree of optical perfection.

VI. Lead-Lanthanum-Zirconium-Titanate

Lead-lanthanum-zirconium-ti tanate (PLZT) is a leading electro-optic ceramic developed by Haertling and Land [29]. Electro-optic materials are useful in devices where we desire to alter the optical properties of a material with an applied voltage or where we wish to transform electrical information into optical information. Applications are wide-aperture electronic shutters, modulators, linear gate arrays, color filters, stereoviewing devices, image storage devices, segmented displays, etc. Single crystals such as LiNbO 3 and liquid crystals compete effectively for many electro- optic applications, but the polycrystalline material PLZT has certain property advantages, such as a high electro-optic coefficient, fast response time, memory capability, etc., that make it attractive to designers [28]. Transparent ceramics were difficult to achieve in this system because not only was it necessary to eliminate the usual porosity, second phases, and grain anisotropy, but it was also necessary to control internal grain twinning and ferroelectric domains.

A. Phase Relations

The subsolid phase diagram for PbO-TiO2-ZrO 2 investigated by Webster et al. [68] is shown in Fig. 27. A complete solid solution series exists across the PbTiO3-PbZrO 3 tie line. The diagram shows that the tetragonal, rhombohedral, and orthorhombic phases exist in this solid solution series. These are room-temperature phases and convert to cubic, a nonferroelectric phase, at low temperatures and are the stable phase at the 1100~ isothermal as illustrated. The tetragonal and rhombohedral phases are ferroelectric, whereas the orthorhombic phase is antiferroelec- tric. The pseudobinary of Fushimi and Ikeda [19] for PbTiO3-PbTiO 3 shown in Fig. 28 illustrates the rising solid-liquidus boundary from 1300 to 1550~ going from the Ti- to the Zr-rich side. This boundary governs the upper temperature limit for processing in PLZT, although we do not have information concerning the effect of La20 3 on the liquidus. Another phase diagram that is important in processing shows the effect of free PbO


P b O

l l O 0 ~

V b T i O ~ g ~ ~ .................. ~r ............ ,~. VbZrO 3

/ / \\ ~ ",, ,~ \

A ~ / / T s s + ZT + 0PT-PZ)ss ~ Zss + ZT + 0 p T _ ~ ' " ~

mT~

(8%) 20 40 60 s o (86%)

TiO 2 TiZrO 4 ZrO 2

M o l e %

FIG. 27. PbO-TiO2-ZrO 2 phase diagram. After Webster et al. [68]. Reprinted by permission of the Canadian Ceramic Society.

1500

o 1400

1300

I I*

L ; L+Z

p [ ,

�9 []

- ~ ~ ~

L+S S �9 : L+S : L+S+Z

[] : L+Z

1200 I I I I I I I I I o3~ 0.1 0 .2 0.3 0 .4 0.5 0 .6 0.7 0.8 0 .9

PbTiO3 PbZrO 3 X

FIG. 28. PbTiO3-PbZrO 3 phase diagram. After Fushimi and Ikeda [19]. Reprinted by permission of the American Ceramic Society.

34 W.H. RHODES

3O0O

L ...__ . . . . . t

r , ~ t s

o / ~, L + Z 1,) I

I.. 2000 '

~ L + P Z I /

1000 PZ + z

P ~

P b O PZ ZrO 2

FIG. 29. PbO-ZrO z phase diagram. After Fushimi and Ikeda [19]. Reprinted by permission of the American Ceramic Society.

on the liquidus. This is illustrated for ZrO2, the most refractory side of the system, in Fig. 29 [19]. PbO melts at 888~ The processing t empera tu res for PLZT are in the 1200~ range, so any free PbO would result in a liquid phase in the presence of solid.

Ferroelectr ic materials in this family have the perovskite crystal structure, where Ti+4(Zr+4) is in the B site at the center of the unit cell within an octrahedral coordination of O-2 , and Pb + 2 (La + 3) occupies the A site at cube corners. Considerable discussion and experimentat ion have occurred to decide what charge-compensat ing defects are created by the

FIG. 30. PbTiO3-PbZrO3-La phase diagram with insets illustrating hystereses behavior and letters denoting major commercial compositions. Reprinted from Haertling [28] by courtesy of Marcel Dekker, Inc.


addition of La +3 to the system. The weight of the evidence according to Haertling [28] favors vacancies on both the A and B site. La203 has a significant solubility in this system, as shown with the ambient temperature phase diagram of Fig. 30. Also shown on this diagram are the major commercial compositions and their respectively hysteresis loops. This diagram illustrates the critical importance of understanding and controlling phase relations for the successful production of these materials. La +3 has 4 atom % solubility on the Zr-rich side and 32 atom % on the Ti-rich side. The addition of La+3 reduces the anisotropy of the ferroelectric phases, which favors the ability to fabricate transparent materials. The

0.8

0.6

0.4

0.2

o.o

O

~, 4

• ~ 3 II

~ 2

o

FERh Io~/~1~p.4.,9[ ' FET~ and Slin~ Loop

/ ~ , ~ PLZpmT#6nS/3sLe

! I 6

12

lO

8

6

4

2

I ~'1~ I I I o37o-o 0

8 10 12 14

~q

w.r • w.r

Atom % La-X

b @

o.,- 6t X, PLZT,

~ x 0.6 ~ 4

0.5 ca 2

,~ [ ' F E m . ~ - ~ I FETet

0.4 ~ 0 ~ .... X - � 9 66 64 62 60 58 56 54 52 50

y . � 9 34 36 38 40 42 44 46 48 50

12

10

8

6 •

4 7

FIG. 31. Selected electrical and optical properties as functions of (a) La content for 65/35

Z r / T i ratio and (b) Z r / T i ratio for La content of 7 atom %. After Haertling and Land [29]. Reprinted by permission of the American Ceramic Society.

36 W.H. RHODES

electro-optic compositions usually have greater than 6 atom % La for this reason. Further, they lie near the rhombohedral/tetragonal phase boundary, where properties are maximized. The influence of phase chemistry on properties as determined by Haertling and Land [29] is illustrated in Fig. 31. The indistinct change in properties across the boundary implies that the boundary is "smeared" as though it reflects second-order or higher transitions, similar to those observed by Goodman [22] in barium zirconium metaniobate. There is an uncertainty of about 4 atom % in the tetragonal/rhombohedral boundary [29]. The slashed zone between the ferroelectric phases and the cubic and orthorhomic phase in Fig. 30 is graphical representation of this same phenomena. Excess La20 3 forms LazTi20 7 or LazZr20 7 plus the solid solution. Pb4La20 7 has also been identified by Land et al. [38] in compositions with excess PbO in the formulation.

B. Liquid-Phase Sintering

Transparent PLZT was first fabricated by hot pressing [26]. Later Snow [65] developed an atmospheric sintering technique to achieve transparency. Sinter-HIPing was successfully employed by Hardtl [30]. The sintering mechanism appears to be similar in all of these cases. Excess PbO must be provided in the original batching formula. If the starting ingredients are formulated according to an A site vacancy model, excess PbO is not provided and transparency cannot be achieved. The excess PbO forms a liquid phase (Fig. 29) at the 1200 to 1350~ processing temperatures. Liquid-phase sintering promotes particle rearrangement and rapid densification through dissolution and reprecipitation from high to low chemical potential sites on the grains. Workers in this field felt that grain growth was retarded due to the longer diffusion distances through the liquid phase. The final grain sizes achieved were in the 2- to 20-txm range, which certainly supports this contention. Extremely long processing times of 18 to 60 h were employed. Shorter times showed that densification was rapid, but left residual PbO phase, which gave the body an amber or red color. Acetic acid leaching experiments carried out by Snow [65] demonstrated that the PbO phase was interconnected since it could be completely leached out. It appears that the long process times were necessary to volatilize the last remnants of the liquid phase. Since it was interconnected, it could diffuse to the surface and vaporize. In this respect, the densification process was very similar to that discussed in Section IV.B for LiF-doped MgAI20 4. Grain rearrangement also occurred during the volatilization step, and accomplishing this without pore genera-

FIG. 32. Microstructures of PLZT. (a) Thermal etch showing grain structure. (b) Chemi- cal etch showing lack of relationship of domain structure to grain structure. After Haertling [27].

37

38 W.H. RHODES

tion required small grain sizes. It was also important that the PbO did not become pinched off and remain a trapped second phase. A shift in the UV cutoff edge with increased sample thickness was interpreted as being due to residual PbO [27].

The role of La20 3 in achieving transparency is less clear than the influence of PbO. The high La20 3 solubility in the Pb(Ti, Zr)O 3 lattice appears to be important. It is known that substitutions such as Ba and Sn for Pb are less effective [28]. La20 3 has a much stronger effect (37~ %) on lowering the Curie temperature than the other substitutions. Thus the crystal structure is closer to cubic in the > 6 atom % La range, where transparent ceramics are achieved. There is the possibility that the role of La20 3 may be similar to the role of ThO 2 in sintering Y203, as discussed in Section II.B, where large solid solution regions were also essential in reaching transparency.

Haertling [27] showed that interrelationship between grain and domain structure to be that illustrated in Fig. 32. These micrographs show that the domains transcend grain boundaries, indicating a lack of grain-boundary domain pinning phases and good bonding. This grain size material is desired when scattering effects are being optimized for both memory and nonmemory applications. Fine (~ 2-/~m) grain sizes are required for memory material making use of the birefringent effect. This is typically accomplished by holding the hot pressing time and pressure constant, and lowering the temperature from ~ 1250 to l l00~ This implies, of course, that grain boundaries have significant influence on domain structure and ease of switching. The domain boundaries and optical anisotropy influence optical scattering, as shown by the fact that raising the temperature above the Curie point into the cubic region of the phase diagram improves transmission [28]. It has also been shown that transparency decreases with increased grain size (> 3/~m), presumably due to an increase in domain boundaries that form to relieve internal strain [38]. This suggests that domain boundaries are more effective scatterers than the optical anisotropy effect associated with traversing grain boundaries.

VII. Conclusions

This review has shown the role of phase diagrams and phase relations in fabricating five polycrystalline oxides to optical transparency. Most of the examples cited had a clear connection between phase relations and the first successes in reaching full density and optical transparency. A120 3 was the only case history where there was not a clear record of the importance


of phase relations in the early development. Even here, later workers have been concerned with the limited solubility of the sintering aids. The LazO3-Y20 3 example was unique because the lack of a phase diagram held back early progress, but once the phase relations were established, they were manipulated advantageously to reach full density. Many other parameters, such as powder properties, green microstructure, temperature, time, pressure, heating rate, gas evolution, etc., play essential roles in the success of a fabrication effort, but without a clear understanding of phase relations, all is lost.

References

10.

11. 12.

13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27.

A. M. Alper, R. N. McNally, P. H. Ribbe, and R. C. Doman, J. Am. Ceram. Soc. 45, 263-268 (1962). R. C. Anderson, in "High Temperature Oxides Part II" (A. M. Alper, ed.), pp. 1-40. Academic Press, New York, 1970. L. M. Atlas, J. Am. Ceram. Soc. 40, 196-199 (1957). R. D. Bagley, U.S. Pat. 3,531,308 (1970). R. D. Bagley and H. K. Bowen, Am. Ceram. Soc. Bull. 56, 732 (1977). J. T. Bailey and R. Russell, Jr., Trans. Br. Ceram. Soc. 68, 159-164 (1969). M. W. Benecke, N. E. Olson and J. A. Pask J. Am. Ceram. Soc. 50, 365-368 (1967). S. J. Bennison and M. P. Harmer, Ceram. Trans. 7, 13-49 (1990). P. Boch, P. Bergez, J. Charpin, P. Raynal, and M. Boncoeur, in "Ceramics Today-Tomorrow's Ceramics" (P. Vincenzini, ed.), Part D, pp. 2343-2358. Elsevier, Amsterdam, 1991. P. A. Bosomworth, M. P. Harmer, H. M. Chan, and W. H. Rhodes, Bull. Am. Ceram. Soc. 65, 499 (1986). R. J. Bratton, J. Am. Ceram. Soc. 57, 283-286 (1974). L. A. Brissette, P. L. Burnett, R. M. Spriggs, and T. Vasilos, J. Am. Ceram. Soc. 49, 165-166 (1966). J. Cassedanne and H. Forestier, C. R. Hebd. Seances Acad Sci. 253, 2953-2955 (1961). R. J. Charles, S. Prochazka, and C. E. Scott, U.S. Pat. 4,285,732 (1980). R. L. Coble, J. Appl. Phys. 32, 787-799 (1961). N. D. Corbin, J. Eur. Ceram. Soc. 5, 143-154 (1989). J. Coutures and M. Foex, J. Solid State Chem. 11, 294-300 (1974). M. Foex and J.-P. Traverse, Bull. Soc. Fr. Mineral. Cristallogr. 89, 184-205 (1966). S. Fushimi and T. Ikeda, J. Am. Ceram. Soc. 50, 129-132 (1967). A. Gatti and M. J. Noone, Final Rep. AMMRC-CR-70-8 (1970). R. L. Gentilman, Am. Ceram. Soc. Bull. 60, 906-909 (1981). G. Goodman, J. Am. Ceram. Soc. 43, 105-113 (1960). C. Greskovich, personal communication (1991). C. Greskovich and J. P. Chernoch, J. Appl. Phys. 44, 4599-4606 (1973). C. D. Greskovich, D. A. Cusano, and F. A. DiBianca, U.S. Pat. 4,466,930 (1984). G. H. Haertling, Bull. Am. Ceram. Soc. 49, 564-567 (1970). G. H. Haertling, J. Am. Ceram. Soc. 54, 303-309 (1971).

40 W . H . R H O D E S

28. G. H. Haertling, in "Electronic Ceramics" (L. M. Levinson, ed.), pp. 371-492. Dekker, New York, 1988.

29. G. H. Haertling and C. E. Land, J. Am. Ceram. Soc. 54, 1-11 (1971). 30. K. H. Hardtl, Bull. Am. Ceram. Soc. 54, 201-205 (1975). 31. T. M. Hartnett and R. L. Gentilman, Proc. SPIE--Int . Soc. Opt. Eng. 505, 15-22 (1984). 32. T. M. Hartnett, E. A. Maguire, R. L. Gentilman, N. D. Corbin, and J. W. McCauley

Ceram. Eng. Sci. Proc. 3, 67-76 (1982). 33. T. M. Hartnett, T. M. Gentilman, and E. A. Maguire, U.S. Pat. 4,481,300 (1984). 34. T. M. Hartnett, M. Greenberg, and R. L. Gentilman U.S. Pat. 4,761,390 (1988). 35. P. Hing, J. Mater. Sci. 11, 1919-1926 (1976). 36. S. M. Ho, U.S. Pat. 3,530,209 (1970). 37. P. J. Jorgensen and R. C. Anderson, J. Am. Ceram. Soc. 50, 553-558 (1967). 38. C. E. Land, P. D. Thacher, and G. H. Haertling, Appl. Solid State Sci. 4, 137-233. (1974). 39. R. A. Lefever and J. Matsko, J. Mater. Res. Bull. 2, 865-869 (1967). 40. A. M. Lejus, Bull. Soc. Chem. Fr., pp. 2123-2126 (1962). 41. E. A. Maguire and R. L. Gentilman, U.S. Pat. 4,347,210 (1982). 42. C. Martin and B. Cales, Proc. SPIEmlnt . Soc. Opt. Eng. 1112, 20-24 (1989). 43. M. Matsui and T. Takahashi, U.S. pat. 4,543,346 (1985). 44. J. W. McCauley, J. Am. Ceram. Soc. 61, 372-373 (1978). 45. J. W. McCauley and N. D. Corbin J. Am. Ceram. Soc. 62, 476-479 (1979). 46. J. W. McCauley and N. D. Corbin, in "Progress in Nitrogen Ceramics" (F. L. Riley, ed.),

pp. 111-118. Martinus Nijhoff, Boston, 1983. 47. M. Mizuno, A. Rouanet, T. Yamada, and T. Noguchi, Yogyo Kyokaishi 84, 342-348

(1976). 48. B. O. Mysen, in "Phase Diagrams for Ceramists," Vol. 8, pp. 5-6. Am. Ceram. Soc.,

Westerville, OH, 1990. 49. L. Navias, J. Am. Ceram. Soc. 44, 434-446 (1961). 50. T. Noguchi and M. Mizuno, (1967). Kogyo Kagaku Zasshi 70, 839-846 (1967); modified in

"Phase Diagrams for Ceramists, 1975 Supplement" (E. M. Levin and H. F. McMurdie, eds.), p. 132. Am. Ceram. Soc., Westerville, OH, 1975.

51. I. Oda, M. Kaneno, and I. Hayakawa, U.S. Pat. 4,273,587 (1981). 52. F. C. Pallila, W. H. Rhodes, and C. S. Pitt, in "Materials Characterization for Systems

Performance and Reliability" (J. W. McCauley and V. Weiss, eds.), pp. 149-187, Plenum, New York, 1986.

53. G. A. Rankin and H. E. Merwin, J. Am. Ceram. Soc. 38, 568-588 (1916). 54. W. H. Rhodes, J. Am. Ceram. Soc. 64, 13-19 (1981). 55. W. H. Rhodes and F. J. Reid, U.S. Pat. 4,174,973 (1978). 56. W. H. Rhodes and G. C. Wei, in "Concise Encyclopedia of Advanced Ceramics" (R. J.

Brook, ed.), pp. 273-276. Pergamon, New York, 1991. 57. W. H. Rhodes, P. L. Berneburg, and J. E. Niesse, AMMRC-CR-70-19 (1970). 58. G. Rossi and J. E. Burke, J. Am. Ceram. Soc. 56, 654-659 (1973). 59. D. M. Roy, R. Roy, and E. F. Osborn, Am. J. Sci. 251, 337-361 (1953). 60. D. W. Roy and J. L. Hastert, U.S. Pat. 5,001,093 (1991). 61. S. K. Roy and R. L. Coble, J. Am. Ceram. Soc. 51, 1-6 (1968). 62. D. J. Sellers and D. W. Roy, U.S. Pat. 3,768,990 (1973). 63. K. Shibata, N. Nakamura, and A. Fujii, Proc. SPIE--Int . Soc. Opt. Eng. 1326, 48-53.

(1990). 64. F. Sibieude and M. Foex, J. Nucl. Mater. 56, 229 (1975); modified in "Phase Diagrams

for Ceramists" (R. Roth et al., eds), Vol. 4, p. 140. Am. Ceram. Soc., Westerville, OH, 1981.


65. G. S. Snow, J. Am. Ceram. Soc. 56, 91-96 (1973). 66. C. M. Sung, G. C. Wei, K. J. Ostreicher, and W. H. Rhodes, J. Am. Ceram. Soc. 75,

1796-1800 (1992). 67. N. A. Toropov, I. A. Bondar, F. Ya. Galakhov, X. S. Nikogosyan, and N. V. Vinogradova,

Izu. Akad. Nauk SSSR, Ser. Khim. 7, 1162-1166 (1964); modified in "Phase Diagrams for Ceramists, 1969 Supplment" (E. M. Levin et al., eds), p. 96. Am. Ceram. Soc., Westerville, OH, 1969.

68. A. H. Webster, R. C. MacDonald and W. S. Bowman, J. Can. Ceram. Soc. 34, 97-102 (1965); modified in "Phase Diagrams for Ceramists, 1975 Supplement" (E. M. Levin and H. F. McMurdie, eds.), p. 250. Am. Ceram. Soc., Westerville, OH, 1975.

69. G. Yamaguchi and H. Yanagida, Bull. Chem. Soc. Jpn. 32, 1264-1268 (1959).

The Use of Phase Diagrams to Predict Alkafi Oxide Corrosion of Ceramics

JESSE J. BROWN, JR.

Department of Materials Science and Engineering Virginia Polytechnic Institute and State University

Blacksburg, Virginia 24061

I. In t roduc t ion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 43

II. L i t e r a t u r e Rev iew . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 44

A. Phase E q u i l i b r i u m Re la t ionsh ips . . . . . . . . . . . . . . . . . . . . . . . . . . 45

B. Reac t ions with Sod ium C a r b o n a t e Vapor s . . . . . . . . . . . . . . . . . . . . . 59

C. Reac t ions wi th Po ta s s ium C a r b o n a t e V a p o r s . . . . . . . . . . . . . . . . . . . 69

III. P red ic t ions of Alkal i Cor ros ion of Alumino-S i l i ca te Ref rac to r i e s

Us ing Phase D i a g r a m s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 74

IV. C o m p a r i s o n of P red ic t ed Na2 ~ Corros ion of Alumino-S i l i ca te Ref rac to r i e s with

E x p e r i m e n t a l Resu l t s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 79

V. Conc lus ions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 80

R e f e r e n c e s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 81

I. Introduct ion

High-temperature corrosion by alkali oxides has always been a serious problem that affects the service life of oxide materials. Refractories used in blast furnaces, ceramic kilns, glass tanks, and cement kilns experience aggressive alkali attack, and refractory selection for these applications routinely involves consideration of high-temperature alkali corrosion. Ap- plications of advanced ceramic materials such as in coal gasification, hot-gas cleanup of fossil fuel combustion, heat exchangers, and gas tur- bines for both power generation and transportation are also prone to alkali attack. Advanced ceramics applications are especially prone to alkali corrosion if coal combustion is involved because of the high temperatures present and the significant alkali content of most coals.

Generally alkali corrosion of ceramics involves the salts or oxides of sodium and/or potassium. These alkalis are corrosive in the solid, liquid,

43 Copyright �9 1995 by Academic Press, Inc.

All rights of reproduction in any form reserved.

44 JESSE J. BROWN, JR.

and vapor states with the vapor phase being the most aggressive. The reaction rates of the corrosion increase with temperature; however, significant corrosion at 800~ is not unusual.

Of the common oxides, the alkalis are somewhat unique in that they are very potent fluxes. This fact has been used to advantage for many years in the glass and ceramics industries in order to reduce the melting/maturing temperatures. On the other hand, ceramics designated for extended use at high temperatures are frequently damaged by the fluxing action of alkali. This is an especially serious problem in any high-temperature application that involves fossil fuel combustion. Here not ofily are the temperatures high but the alkalis are frequently present as vapor phases. As the alkalis react with the low thermally expanding high thermal shock resistant ceramic, either a fluid, melt (sometimes water soluble) forms or a very high thermally expanding crystalline phase forms. Both of these result rapidly reduce the integrity of the ceramic and eventually result in material failure and frequently system failure.

Phase diagrams are extremely useful in determining the reactions that occur when alkali oxides react with many common ceramics. Most scientists and engineers are easily able to evaluate binary phase diagrams that correspond to an alkali reaction with single oxide ceramics; however, when multi-oxide ceramics such as mullite are involved, multi-component phase diagrams are not fully used and extremely time-consuming experimentation is unnecessarily conducted.

II. Literature Review

A huge amount of information related to the alkali corrosion of ceramics is available in existing phase equilibrium diagrams. Too often this information is not being used; instead, expensive experimentation is being conducted to reestablish information already known and generally available. In addition to an extensive literature review of the alkali corrosion of refractories, it is the purpose of this chapter to illustrate the value of using multi-component phase equilibrium diagrams to evaluate high- temperature reactions. Alkali oxide (Na20 and K20) corrosions of alumino-silicate ceramics are used as examples.

Finally, the inexperienced reader should note that phase equilibrium diagrams only tell what will eventually happenmat equilibrium. These diagrams do not contain information relating to the kinetics of the reactions. In general, however, reactions involving alkali oxides are relatively fast for ceramic systems. If, for example, the phase diagram reveals liquid

2 P R E D I C T I O N O F A L K A L I O X I D E C O R R O S I O N 45

2 . -

LIQUID 1900 ~ COR.

-SiO,+LIQ. J / l " J723 o / / " I P"..

- I ',' / I CORUNDUM v-. I I + CRISTOBALITE + MULLITE

15001.1470 o [- _ _ ~ [ MULLITE,, l TRIDYMITE +MULLITE, , II

SiOz 20 I 40 60 '8~0 AI20• WEIGHT PERCENT ALUMINA

, I I SEMISILICA I MULLITE CORUNDUM

SILICA FIRECLAY H IGH ALUMINA

(3 o.., 1800 lad n."

~ 1700

W

FIG. 1. The phase diagram for AI203-SiO 2 showing common refractory compositions [1]. Reprinted by permission of the American Ceramic Society.

formation at the temperature of interest, expect this to occur relatively fast in days or weeks, not years, providing sufficient alkali is present.

A. Phase Equilibrium Relationships

1. THE A 1 2 0 3 - S i O 2 SYSTEM

The A1203-SiO 2 system contains the compositions of a wide range of alumino-silicate refractories. These include silica, semisilica, fireclay, low alumina, mullite, high alumina, and corundum. Much of the high temperature behavior of these refractories can be attributed to relationships expressed in the phase diagram (Fig. 1) [1]. 1 The diagram is characterized by its two end members, A120 3 and SiO2, and an intermediate compound mullite, 3A1203.2SiO 2. There are two eutectics in the system: one between cristobalite and mullite at 1595~ and the other at 1840~ between mullite and corundum. As can be seen from the diagram, mullite takes a limited amount of alumina into solid solution.

lFor the most current discussion of the AI203-SiO 2 phase diagram see R. F. Davis and J. A. Pask, J. Am. Ceram. Soc. 55(10), 525 (1972); I. A. Aksay and J. A. Pask, ibid. 58(11-12), 507 (1975); and S. H. Risbud and J. A. Pask, ibid. 60(9-10), 418 (1977).


With mixtures containing less A1203 than mullite, liquid occurs above 1590~ The addition of 5 wt % A120 3 lowers the melting point of SiO 2 from 1725 to 1595~ Thus, A120 3 acts as a weak flux for SiO 2. Neverthe- less, silica and semi-silica compositions (Al20 3 content of less than 25%) can be used at temperatures near the melting point since the liquid that is formed is extremely viscous. However, even small amounts of alkalies can render the liquid more fluid.

As the alumina content is increased from 5% to 72%, the melting points of compositions in this range increase as shown in Fig. 1. As mentioned earlier, maximum use temperature of fireclay (25% to 45% A120 3) and low alumina (45% to 65% A120 3) compositions is shown approximately by the dashed line in the diagram. Compositions in this range can tolerate limited amounts of liquid during service due to the viscous nature of the melt.

The phase diagram shows that the most refractory mixtures contain 72 wt % to 100 wt % A120 3 (high-alumina compositions). A liquid phase develops in these mixtures (if pure) at a temperature above 1840~ Refractories in this composition range can be used to temperatures near 1840~

The phase diagram for the NazO-AI203-SiO 2 system illustrates the thermochemical relationships between these three components [2-4]. Knowledge of the equilibrium relations in this system is pertinent to the attack of alumina-silica ceramics by alkali vapors.

The phase relations in the NazO-Al203-SiO 2 and KzO-AI203-SiO 2 systems were investigated extensively by Schairer and Bowen [4] and their complete results were published in 1956. Schairer and Bowen prepared glasses of the desired composition and employed conventional quenching techniques for phase analysis. Their study clarified most of the equilibrium relations in these systems. However, due to experimental difficulties associated with the necessary high temperatures of melting and probabilities of compositional changes by alkali loss, they were unable to extend their investigation to the high-alumina portion of the system. Consequently, no reliable phase-equilibrium data were available on the NazO-AI20 3 system until DePablon-Galan and Foster [5] conducted their experiments in 1959. The primary concern of their study was establishing beta-alumina as a discrete compound rather than a mere metastable polymorph of corundum. Their findings justify the insertion of a tie-like line linking beta- alumina to carnegieite and even indicate a large primary field of beta- alumina. In 1960 Osborn and Muan [6] revised the NazO-AI203-SiO 2 system, which is shown in Fig. 2.

The system contains three sodium silicates, 2NaO �9 SiO2, Na20 �9 SiO2, and Na20 �9 2SIO2; two sodium aluminates, Na20 �9 A120 3 and N a 2 0 .

2 PREDICTION OF ALKALI OXIDE CORROSION 47

SiOz 1723" Crystalline Phases

Notation Oxide Formula Cristobolite ] Tridymite Si02 Quartz Beta-A lumina No=O. IIAI=O 3 Corundum AI203 Mull i te 3AIzO 3. 2SiOz Albite No20 �9 AIz03.6SiO~. Nepheline Carnecjieite t NozO. AltO 3 �9 2 SiO~.

N~

2NoeO" SiOe

- - --'-'~-.."~i \ "" %""~'~ "%~ "~4~" .!,.. :!, , ,840.

to-Alumina k ",,.

V V V V v ~ ~ v v ', ~o Na,O-A~ N~..A~ A~6,

�9 ,,,, L:~20 �9

Fie. 2. Phase diagram of the N a 2 0 - A l 2 0 3 - S i O 2 system [2]. + + + + Inferred maximum extent of solid solution in carnegieite phase. Temperatures up to approximately 1550~ are on the Geophysical Laboratory Scale; those above 1550~ are on the 1948 International Scale. Reprinted by permission of The American Ceramic Society.

11A1203; and three ternary compounds, albite (Na20" A1203 �9 6SiO 2) and nepheline and its high-temperature modification carnegieite (Na20 �9 A1203 �9 2SIO2). Some of these solid phases are of variable composition as a result of solid solution.

2. ThE N a 2 0 - S i O 2 SYSTEM

The N a 2 0 - S i O 2 system is particularly applicable to silica and semi-silica compositions exposed to soda vapor. The high silica portion of the phase equilibrium diagram [7, 8] is shown in Fig. 3. The system contains three binary compounds: the orthosilicate (2Na20" SiO2), the metasilicate (Na20 �9 SiO2), and the disilicate (Na20 �9 2SIO2).

48 J E S S E J. B R O W N , JR .

1700

1600 Crisfobolife + Liquid . . . . . .

1500

1400

Liquid

leO0

I100

NazO \ 2NozO'SiOz Liquid \ +Liquid . . . . . .

I . -~ , " I N O 2 U ' ~ I U z ----[--1.1~/,-,/.,,, + Liquid

+

Liquid

I000 ZNazO',SiOe 0.25i

900 - Naz0.Si0 ~ ' n e T * - Quartz + Liquid

800- 9_" ,,~ o Na20"25i02 §

700- c~ ~IN t__.__.T0z. . O ~U. _ ~6,780 z

~ m

30 40 50 60 70 80 90 I00 "~'-NozO SiOz

FIG. 3. Phase diagram of the N a 2 0 - S i O 2 system [7]. Reprinted by permission of the American Ceramic Society.

The orthosilicate is the most basic sodium silicate; however, the compound is not stable at its melting point. In mixtures containing less than 40.7 wt % SiO e, the orthosilicate decomposes at 1120 + 5~ to a solid, presumably NaeO, and a liquid before its melting point is reached. The NaeO released tends to absorb CO e, so that, if any CO e is present some of this gas is retained in the melt. Thus, the retention of CO e may be expected to lower the dissociation temperature of 2NaeO �9 SiO e.

Silica crystallizes from sodium silicates in three forms, cristobalite, tridymite, and quartz. The inversion temperatures are 1470 and 870~ Cristobalite, the high-temperature modification, melts at 1713~ The cristobalite liquidus decreases from the melting point of SiO e to the inversion point (located at 88.7 wt % SiO e) between cristobalite and tridymite. The tridymite liquidus then descends from this point and meets the liquidus curve of quartz at 75.5 wt % SiO e (870 + 10~ The tridymite liquidus extends metastable below 870 to 793~ ending at the disilicate-


quartz eutectic. Only the high-temperature modification of quartz, stable above 573~ is encountered at the liquidus.

The sodium disilicate-quartz eutectic occurs at 74.2 _+ 0.3 wt % SiO 2 and 789~ [8]. A metastable eutectic between Na20 �9 2SiO 2 and tridymite exists at 74.6 wt % SiO 2 and 782~

The field of sodium disilicate extends from the quartz-disilicate eutectic to the metasilicate-disilicate eutectic originally located at approximately 62.1 wt % SiO 2 and 837 _+ 1~ The compound N a 2 0 . 2 S i O 2 melts congruently at 874 _+ 3~ and has several polymorphic forms. There are probably three stable forms in addition to several metastable ones. The high-temperature form melts congruently at 874 + 3~ There are rapid reversible transitions at 707 and 678~

N a 2 0 . 2 S i O 2 takes into solid solution excess Na20 and SiO2, with the difference that unmixing temperatures lie entirely below the liquidus. Compositions containing excess Na20 show unmixing at 706~ and those with excess SiO 2 at 768~

The melting point of this Na20 �9 SiO 2 is 1089 + 0.5~ [9], and the sodium orthosilicate-sodium metasilicate eutectic is at 41.1 wt % SiO 2 and 1022~

3. THE N a 2 0 - a l 2 0 3 - S i O 2 SYSTEM

The primary phase fields that appear on the liquidus surface within the NazO-AI203-SiO 2 system are cristobalite, tridymite, quartz, sodium disil- icates, sodium metasilicate, mullite, corundum, nepheline, carnegieite, albite, and sodium orthosilicate solid solution.

In appropriate NazO-AI203-SiO 2 melts, tridymite is very difficult to obtain, whereas quartz crystallizes easily. Both sodium disilicate and sodium metasilicate form readily from their respective melts. Although sodium disilicate has several polymorphs, only the high-temperature form occurs in NazO-AI203-S iO 2 melts.

Since mullite melts incongruently at 1810 _+ 10~ its composition lies outside its primary phase field. For compositions in the field of mullite with a liquidus temperature below about 1400~ Schairer and Bowen [4] encountered considerable difficulty in crystallizing this compound. For compositions with liquidus temperatures below around 1200~ many months were required to form small mullite needles.

From most of the compositions in the field of corundum, both corundum and beta-alumina crystallize at about 100 ~ below liquidus temperatures. For liquidus temperatures below 1400 ~ crystallization of both of these phases proved difficult.


The compound albite, N a 2 0 . A 1 2 0 3 . 6 S i O 2 , melts congruently at 1118 _+ 3~ For those compositions in the albite field with a liquidus of about 1040~ or less, albite crystallized in a few days at a temperature 50 to 75 ~ below liquidus temperatures. To the left of the carnegieite-silica line, a rapid decrease in the viscosity of melts in the albite field occurs. As a result, it is possible to crystallize albite even at temperatures around 750~ In the albite field toward the system albite-silica, crystallization of albite is very difficult.

Pure soda nepheline is the low-temperature polymorph of NaA1SiO 4. The high-temperature modification, carnegieite, melts congruently at 1562 + 2~ and is the stable form above 1254 + 5~ [10] (the inversion temperature between carnegieite and soda nepheline). Carnegieite is isometric at high temperatures but undergoes a rapid reversible metastable transformation at 692.1~ on heating and at 687.0~ on cooling to a twinned low-temperature form [ 11].

In the nepheline-albite system, both nepheline and carnegieite take up limited amounts of albite in solid solution, about 34% and about 14.5%, respectively. As a result the inversion temperature of nepheline to carnegieite is raised to about 1280~ Carnegieite takes up to 24% NazSiO 3, sodium metasilicate, in solid solution, while there is no appreciable solid solution of NazSiO 3 in nepheline. Consequently, the inversion temperature is lowered to 1163~ If no solid solution occurs in either carnegieite or nepheline, the inversion temperature remains constant. Schairer and Bowen [4] suggest a complete series of solid solution between pure carnegieite and Na20 �9 A120 3 crystals, with the probability of a continuous rise in the liquidus curve from 1526~ at pure carnegieite to an estimated temperature of 1850 _+ 30~ at Na20 - A120 3. Results by Matignon [12], however, are not in accord with these data. He prepared Na20 - AI20 3 and measured its melting point at 1650~ Brownmiller and Brogue [13] prepared a homogeneous so|id phase of the composition Na20 �9 A120 3 and found no melting or dissociation at that temperature. Further study of the carnegieite-sodium aluminate system is indicated to resolve these discrepancies and therefore fix the liquidus surface.

As previously mentioned, silica brick can withstand temperatures close to its melting point. This may be attributed to silica's ability to form immiscible liquids in the presence of lime or iron oxide. No immiscibility, however, is present with silica in the NazO-S iO z -A I20 3 system. With small additions of alumina or alkali oxides to silicate systems in which immiscibility occurs, the miscibility gap is narrowed and ultimately eliminated. A range of from 0.3% to 0.5% alkali or 2.5% to 3.0% AI20 3 is sufficient to cause this effect, specifically alkalies form low melting eutectics with silica. As a result, the presence of these oxides even in small


1800

1600

1400

_~ 1200

E #. ~ooo

8OO

to two liquids at 1789 ~ and 20% mi02 to peritectic at

- Na20(~790) ~ KaO(-780 )

600 I I 1 I I I 10 20 30 40 50 60

Si02 Wt. Percent Second Oxide

F~G. 4. Liquidus curves of binary basic oxide-silica systems.

amounts decreases the resistance of silica brick to alkalies. The strong fluxing action of soda on alumina-silica mixtures high in silica is indicated by the resulting liquidus being as low as 800~ Graphically, the nature of this oxide effect is represented in Fig. 4. The liquidus curves in binary silicate systems have been drawn with the existence of immiscibility shown by the occurrence on the curve of a monotectic horizontal. As can be seen, no immiscibility (horizontal) is present for K20 , Na20 , Li20 , and A1203 [14]. In addition, the refractoriness under load is lowered due to the increased formation of a liquid phase at high temperatures in unused brick containing 2.0% CaO added as a mineralizer [3].

The effect of the addition of soda on the refractoriness of A1203-SiO 2 compositions can be seen from the Na20-A1203-SiO 2 phase diagram. Concerning the present research project, particular attention is called to the incompatibility of the mullite and nepheline phases at equilibrium. In this ternary system, there is no high melting compound analogous to leucite ( K 2 0 . A1203.4SIO2). An examination of the K20-A1203-SiO 2 diagram shows that compositions in the triangle leucite-KA1SiO4-A1203 begin to melt only at 1553 _+ 5~ the temperature of the ternary eutectic between leucite, orthorhombic KA1SiO 4 and corundum [15]. No similar area of compositions with a high temperature at the beginning of melting


exists in the Na20-A1203-SiO 2 system. All compositions in the nepheline-albite-corundum region begin to melt at 1063 _+ 5~ the temperature of the respective ternary eutectic (Na20, 13.8 wt %; A1203, 23.8%; SiO2, 62.4%). Addition of soda to pure mullite refractories yields liquid at low temperatures until carnegieite is formed. Due to the extreme viscosity of a liquid of albite composition and of liquids in the corundum and mullite fields with liquidus temperatures below about 1400~ equilibrium between liquid and crystals is very slow. Even when only a few crystals are present these crystals are very small; consequently, equilibrium is not attained for several days. With large mullite crystals and smaller amounts of liquid present considerable difficulty in establishing equilibrium is to be expected.

4. ALKALI Loss

Alkali loss has been observed in compositions of low silica content in several alkali-silica systems with other oxides. Goldsmith [16] found a loss of Na20 by volatilization from melts in the NaA1SiO4-CaO �9 A1203 system, particularly in those melts rich in the latter component. He noted that the loss of Na20 influenced the amount and even the presence of nonsodic or soda-poor phases such as CaO �9 2A1203 and beta alumina. Schairer and Bowen [4] determined that in multi-component silicate systems, alkali losses are not controlled by alkali content alone. Rather the presence of other oxides has been observed to affect the alkali losses. For example, silica and alumina tend to stabilize compositions and prevent or minimize alkali losses, whereas CaO or MgO tends to increase the losses.

5. SYSTEM K 2 0 - A 1 2 0 3 - S i O 2

Concurrent ly with the phase equil ibrium studies on the NazO-A|203-S iO 2 system, the melting relations in the KzO-A |203-S iO 2 system were investigated extensively by Schairer and Bowen [15]. By means of standard quenching experiments, the liquidus surfaces and the fields of primary phases, cristobalite, tridymite, quartz, mullite, corundum, potash feldspar, leucite, hexagonal and orthorhombic KA1SiO 4, potassium tetrasilicate, and potassium disilicate were located.

As a result of the hydroscopic nature of powdered potassium silicate glasses and KzO-AI203-SiO 2 glasses, particularly those with a silica content of less than 75%, difficulties were encountered when subsequent quenching runs were made to verify previous runs. Even when stored in desiccators, the alkali-silica glasses absorb moisture. When these glasses were reheated to drive off the water, the quenching crystals grew too large


for prompt attainment of equilibrium between crystals and liquid, and thus the quenching runs indicated that the composition of the mixtures had shifted slightly. As Kracek, Bowen, and Morey [17] found, the moisture in these glasses is not merely absorbed on the surfaces of the grains, as in the case with most ordinary glasses, but penetrates into the interior of the grains as a consequence of the high solubility of water in these glasses.

Similarly, Schairer and Bowen [4] noted that Na20-S iO 2 glasses are hygroscopic, particularly those with less than 60% silica, and that NazO-AI203-SiO 2 glasses with less than about 40% silica are hygroscopic. The soda glasses, however, are much less hygroscopic than the corresponding potassia glasses. In addition, the compositional shifts associated with potash loss from hydration of the potassia glasses do not occur in soda glasses. This may be attributed to the fact that Na2 O is much less volatile than K 2 O.

As is the case for the NazO-AI203-SiO 2 ternary system, the KzO-AI203-SiO 2 system is limited by the three respective binary systems: KzO-A1203, K20-SiO2, and Al203-SiO 2.

6. THE K 20-A1203 SYSTEM

Apparently a detailed investigation of the K20-A1203 system has not been conducted or reported since the phase relations between these two compounds have not been published in a phase diagram. Studies of this system have most likely been limited due to experimental difficulties associated with the high temperatures required for melting. In addition, the hygroscopic nature of compositions in the KzO-A120 3 system precludes the collection of reliable data.

Brownmiller [18], however, studied eight compositions in this system. No melting occurred in any of the preparations when heated to 1600~ Microscopic examinations of all but two of the samples heated at 1600~ showed the presence of two phases: K z O . A 1 2 0 3 and beta-alumina. When the ratio of K20 to A1203 was only 1 : i the compound K20 �9 A120 3 crystallized. At 1650~ the highest temperature possible in this study, Brownmiller reported no melting or dissociation of K20 �9 A1203; however, this compound was notably hygroscopic. For the composition 96.5 wt % alumina, 3.5 wt % K z O, corundum and beta-alumina crystallized when heated at 1550~

Strokov et al. [19] studied the reactions between alkali carbonates and alumina. When the molar ratio of K2CO3 " A 1 2 0 3 was 1 : 1 only KzO "

AlzO 3 formed on fusion between 900 and 1110~ If the amount of KaCO 3 present was such that the ratio was 2 :1 or 3:1, the excess K20 evaporated and only K20 �9 A120 3 crystallized.


7. THE K2SiO3-SiO 2 System

The phase relations in the K2SiO3-SiO 2 system are illustrated by the phase diagram of Fig. 5. This diagram is relevant to the study of corrosion of high-silica-content refractories subjected to potassia vapor. The system contains three compounds, K 2 0 . SiO2, K 2 0 " 2SiO 2, and K20 �9 4SiO 2, whose melting points are 976, 1045, and 770~ respectively.

Preparations in this system, particularly those richer in K20 than the disilicate, are extremely hygroscopic. Kracek et al. [17] observed that moisture in the glasses tends to produce excessively large crystals on devitrification, and thus introduces inhomogeneity into the preparations. For this reason researchers avoided the use of the hydrothermal methods of crystallization formerly used to determine the phase relations in this system. Consequently, the diagram shown in Fig. 5 represents a major revision from earlier diagrams of this system.

1700

1600

15 O0

1400

I100

I000

900

8 0 0

. -~-KlO

I I I I I , I �9 I I I

LIQUID

GRI5 TOBALITE A § LIQ

/ " LIQUID -

':', / �9 0 �9 I . . . e i

. . ~ I

. . . . . . . . . - ' ! I r I

, , ! , ",, , ,

40 5 0 60 70 80 90 I00 SiO z

FIG. 5. Phase diagram of the K2SiO3-SiO 2 system [17]. Reprinted by permission of the American Ceramic Society.


As the diagram indicates, the K20 �9 SiO 2 and K20 �9 2SiO 2 compounds melt congruently. The tetrasilicate, K 2 0 " 4SiO 2, is also shown as melting congruently. However, the authors found that this compound has a tendency toward incongruent melting. A mixture whose composition is theoretically K 2 0 �9 4SiO 2 crystallized to be wholly K 2 0 . 4 S I O 4 , with only a faint trace of an additional phase present. When heated in a quenching furnace at 765~ the K20 �9 4SiO 2 crystals completely disappeared but a small quantity of quartz remained. Heating for a much longer time did not appreciably modify the results. Because of the sluggish dissolution of quartz, Kracek et al. did not attribute too much significance to the presence of quartz in the melt. This result indicates a tendency of the K 2 0 . 4 S i O 2 compound toward incongruent melting. Nevertheless, further information on this point showed that the compound has a congruent melting point. As a result Kracek et al. reported that K 2 0 " 4 S i O 2 probably melts congruently at 770~

Preparations of 46% to 56% SiO 2, when completely divitrified, contain crystalline K 2 0 " 2SiO 2. In this region of composition, Kracek et al. [17] reported an inversion of 594~ for K 2 0 . 2 S i O 2. In an earlier study, the same researchers observed that the high-temperature form of this compound takes into solid solution small amounts of both K2 O and SiO 2. As a result the differential heating curves show effects at about 810~ and 980 to 990~ on the two sides of the K 2 0 . 2 S i O 2 composition. Kracek et al. attributed these arrests to the decomposition of the solid solutions (crystallized from the melt) into the pure high-temperature form of K 20 �9 2SiO 2 and liquid at 814 and 993~ However, in their subsequent paper no mention was made of this phenomenon; the breaks in the K 20 �9 2SiO 2 liquidus curves were deleted from the diagram without explanation.

Two small reversible inversions, with some tendency toward delayed occurrence, at 740 _+ 5~ and 760 + 10~ on heating and at 732 _+ 2~ and 717 _+ 2~ on cooling were also noted in this region of compositions. Finally eutectic melting of K zSiO 3 beginning near 780~ and ending at 805 to 815~ depending on the heating rate occurred in this compositional range.

8. FIELDS OF THE PRIMARY CRYSTALLINE PHASES

IN THE K20-A1203-S iO 2 SYSTEM

The primary phase fields that appear on the liquidus surface within the KzO-AI203-S iO 2 system are shown in Fig. 6 [20]. The field of cristobalite extends from 1470 + 10~ the temperature of the boundary curve between cristobalite and tridymite to the melting point of cristobalite at 1713 + 5~

56

Crislobolilt

Ouorlz,

SiOz 1725 ~

JESSE J. B R O W N , JR.

Crystalline Phases Notation Oxide Formula

Cristobolite 1 Tridymite SiO 2 Quartz Corundum AIz05 Mullite 3AI203" 2SIO2 Polo sh K20 "AI203 "GSiO2

Feldspar Leucite KzO'AI203"4SiO2

Koliophilite K20" AI203" 25 i02

k~.sio~

AW i V V V V V /~ KO AiM

3AIzO3-2SiOz I~ ,.,.185o �9

KzO..AIzO~ AV~ ,v 2020 ~

FIG. 6. Phase diagram of the K 2 0 - A 1 2 0 3 - S i O 2 system [2]. Temperatures up to approximately 1550~ are on The Geophysical Laboratory Scale" those above 1550~ are on The 1948 International Scale. Reprinted by permission of the American Ceramic Society.

In appropriate melts, tridymite crystallizes easily in its stabi|ity range except in the very viscous melts whose compositions lie on or very near the leucite-silica line. Quartz, however, is very difficult to obtain in melts at appropriate temperatures. On the other hand, in N a z O - A I 2 0 3 - S i O 2

melts, tridymite is very difficult to obtain, whereas quartz crystallizes without particular trouble [4]. Buerger [21] indicated that tridymite and cristobalite have open crystal structures capable of enclosing foreign ions, which may prevent repacking at lower temperatures and thus the inversion to quartz.

In the K z O - A I 2 0 3 - S i O 2 system, the field of quartz is small. Schairer and Bowen [4] studied only three compositions in this field. Tridymite crystals grew in these compositions and the metastable tridymite liquidus


was obtained. Kracek et al. [17] found only a 2~ difference in temperature between the stable quartz-potassium tetrasilicate eutectic and the metastable eutectic with tridymite. Accordingly, the boundary curve between quartz and tridymite was placed at the 867~ isotherm.

As previously discussed, both potassium tetrasilicate and potassium disilicate melt congruently at 770 _+ 2~ and 1045 + 2~ respectively. Both compounds crystallize readily from appropriate melts when no potash feldspar crystals are present. In the region of the boundary curve between the tetrasilicate field and potash feldspar field, it is possible to follow the metastable extension of the potassium tetrasilicate liquidus surface beneath the stable potash feldspar field. Near the boundary curve between the disilicate field and the potash feldspar field, a metastable extension of the disilicate liquidus surface beneath the stable potash feldspar field is also possible providing once again that no potash feldspar crystals are present.

With liquidus temperatures below about 1400~ for those compositions in the K20-A1203-SiO 2 system in the mullite field, considerable difficulty in crystallization is encountered. For compositions with liquidus temperatures around 1200~ many months are required at temperatures about 75~ below estimated liquidus temperatures for crystallization. In the field of corundum, both corundum and beta-alumina crystallize in preparations held about 100~ below liquidus temperatures.

The ternary compound, potash feldspar, has an incongruent melting point at 1150 + 20~ In all compositions that lie in this field, considerable difficulty is experienced in obtaining potash feldspar crystals. Even after several months of crystallization at temperatures about 50 to 75~ below liquidus temperature, Schairer and Bowen [15] observed only a few percent of very small feldspar laths in the glasses. For those extremely viscous melts that also lie on or very near the leucite-silica line, Schairer and Bowen never crystallized feldspar.

The K20-A1203-SiO 2 diagram shows a large field of leucite, K 2 0 . A1203 �9 4SIO2, with a congruent melting point at 1686 _+ 5~ Leucite crystallizes readily from melts in this field except for (1) those compositions near the ternary reaction point (K20, 32.1%; A120 3, 5.3%; SiO 2, 62.6% at 810 _+ 5~ and (2) those compositions with liquidus temperatures below about 1400~ and very near that portion of the leucite-silica line that lies in the leucite field.

A field of hexagonal KA1SiO 4 and also a field of orthorhombic KA1SiO 4 are shown in Fig. 6. The orthorhombic form of K20 �9 A120 3 �9 2SiO 2 has a congruent melting point known to be above 1700~ and probably near or above 1750~ but the precise melting point is unknown. No euhedral crystals of hexagonal KA1SiO 4 were found in the preparations studied by Schairer and Bowen.


Knowledge of the polymorphs of KA1SiO 4 and their stability relations is often contradictory. The compound was first described as the mineral kaliophilite, a hexagonal mineral from Mount Somma. Bowen [22] conducted the first systematic synthetic study of pure KA1SiO4 and of its relations to NaA1SiO 4. In that study, a melt of pure synthetic KA1SiO 4 was prepared with extreme difficulty from KHCO 3, A120 3, and SiO 2. The glass was easily crystallized at 1300~ to hexagonal crystals. At higher temperatures, crystals of twinned orthorhombic form were encountered. Crystals of the hexagonal form were converted to the orthorhombic form but it was not possible to convert the orthorhombic crystals back to the hexagonal form. Therefore, the possibility of an enantiotrophic inversion at 1540~ or perhaps a little lower, was indicated.

Bannister and Hey [23] identified the hexagonal mineral kalsilite from Uganda, as a polymorphous modification of KA1SiO4, which is not isomor- phous with nepheline, NaA1SiO 4. The x-ray structure of Uganda kalsilite, however, differs from hexagonal kaliophilite from Mount Somma. Rigsby and Richardson [9], from a study of artificial kalsilite and orthorhombic KA1SiO 4 in blast-furnace linings and from synthetic studies of potassium aluminum silicates, indicate the existence of three forms of KA1SiO4: (1) a low-temperature hexagonal modification, probably similar to natural kaliophilite and formed below 900~ (2) a high-temperature orthorhombic modification, sometimes twinned and sometimes free from twinning, formed at temperatures above 1000~ and (3) the hexagonal mineral kalsilite, which can be formed only in the presence of some Na20 at temperatures between 650 and 1200~ and is converted to orthorhombic KA1SiO 4 below 1300 or above 1500~ depending on the soda content.

Tuttle and Smith [24] made a study of the KA1SiO4-NaA1SiO4 system by both dry and hydrothermal methods. They reported that hexagonal kalsilite is the stable form of pure KA1SiO 4 up to 840~ at atmospheric pressure, and above this temperature an orthorhombic form is stable. In addition, they observed a different "high-orthorhombic" form in some of the melts with moderate amounts of NaA1SiO 4 present. Kalsilite takes up limited amounts of NaA1SiO 4 in solid solution. A miscibility gap exists between kalsilite with limited NaA1SiO 4 and low-nepheline with limited but larger amounts of KA1SiO 4. Schairer and Bowen [15] in their liquidus studies of NaA1SiO4-KAISiO4-SiO 2 found a ternary invariant point with leucite, an orthorhombic (K, Na)A1SiO 4 solid solution and a hexagonal Na, K nepheline in equilibrium with a liquid at about 1460~

In the studies of the K 2 0 - A 1 2 0 3 �9 SIO2-A1203 system and that portion of the ternary system that includes the fields labeled hexagonal, KA1SiO 4 and orthorhombic KA1SiO 4 were investigated in the years between 1931 and 1935. The crystals obtained from the melts were small and rounded,


sometimes intergrown with fine-grained material formed during the quenching. Only at temperatures above 1500~ were euhedral crystals observed. These were always the orthorhombic form of KA1SiO 4. This area of the ternary requires further study in light of the confusing results on the polymorphs of KA1SiO 4 discussed earlier.

The evidence presented by Bowen [22] of an enantiotrophic inversion between hexagonal and orthorhombic KA1SiO 4 at about 1540~ was accepted for the purpose of drawing the KzO-AI203-SiO 2 phase diagram. Schairer and Bowen [15] assumed no solid solution in this system and have placed the binary and ternary inversion points at approximate compositions of K20 , 35.6%; A120 3, 19.0%; SiO2, 45.5% and K20 , 28.5%; A1203, 22.0%; SiO 2, 49.5%; respectively, at about 1540~

The compound K 2 0 . A 1 2 0 3 . SiO 2 described by Weyberg [26] was prepared from melts with the aid of a flux (kaolin with K zCrO4). Bowen [22] in 1917 obtained this same compound while preparing the orthorhombic form of KA1SiO 4 with the aid of fluxes. The compound was present in the sample in small amounts as octahedra.

In studies on the KzO-AI203-SiO 2 system by Schairer and Bowen [15], an attempt was made to prepare the compound K20 �9 A120 3 �9 SiO 2 from the appropriate amounts of potassia, silica, and alumina. At temperatures above 1700~ the charge did not sinter and thus was inhomogeneous due to the presence of some undissolved alumina. This result indicates that K20 �9 A120 3 �9 SiO 2 likely has a very high melting point.

The effect of the addition of potassia on the refractoriness of AI203-SiO 2 compositions can be seen from the phase diagram in Fig. 6. All compositions in the corundum-mullite-leucite area are completely solid at temperatures below 1315 +_ 10~ In the mullite-potash feldspar- leucite area; all compositions are completely solid at temperatures below 1140 _+ 20~ For those in the mullite-potash feldspar-silica area, only at temperatures below the ternary eutectic at 985 + 20~ are the compositions totally solid.

B. Reactions with Sodium Carbonate Vapors

Studies regarding the reactions between sodium carbonate and potassium carbonate vapors and refractories have been limited in scope. (Many studies conducted by industry may not have been reported in the literature.) The majority have involved the examination of the degenerated features of refractories after an average service campaign. Most of the articles published have concerned the reactions between alkalies and alumino-silicate refractories utilized in glass melting and blast furnaces.


Investigations of this nature, however, consider the combined effect rather than the individual effects of each component present in the furnace atmosphere.

To determine the mechanisms of alkali attack and resulting reaction products, laboratory-prepared powders of alkali and alumino-silicate refractories have been fired at various temperatures. Results obtained from these reaction tests have supported the findings from the field samples. Alkali slag tests are performed by the various manufacturers to determine how their products withstand corrosion by alkalies.

Other studies have involved reactions of alkali salts with the minerals or raw materials used in making refractories. To a lesser extent, experiments have been conducted on unused brick, but these were primarily concerned with mullite or zircon refractories.

Although many different types of investigations in the area of alkali corrosion of refractories have been performed, there still appears to be doubt regarding the best method for testing. In addition, there are few comparative studies in the literature showing the differences in alkali corrosion of the different alumino-silica refractories. Studies concerning silicon carbide bricks exposed to alkalies are practically nonexistent.

1. REACTION WITH LABORATORY PREPARED SAMPLES

Farris and Allen [27] used two techniques to study the effect of alkali on various refractories. One, a reactive test, concerned the alkali attack on refractories in the range of 42% to 90% alumina content. Specifically, a superduty and bricks of 60%, 70%, 70% high fired, and 90% alumina were ground and mixed with alkali and fired to temperatures from 1575 to 2625~ The resulting reactions products were identified by x-ray diffraction techniques (XRD).

The second approach, an alkali cup test, was performed on straights ranging in composition from 42% (superduty) alumina to 70% alumina. This procedure, however, proved less desirable due to the need to accelerate the reaction procedure. An abundance of the reactant material was placed in a pocket in the sample and fired one or more times to achieve deterioration.

In both tests, soda as well as potassia were used as reactant materials. Although Farris and Allen in their discussion Concentrated predominantly on the results obtained with Na20, the amount of Na20 utilized in the reaction test was 13% added as NazCO 3 powder. The reaction products were identified as nepheline ( 1 : 1 : 2 ) and a ( 3 : 2 : 4 ) sodium aluminum


silicate. Farris and Allen pointed out that these silicates form below the lower test temperature of 1575~ and some exist through 2625~

Specifically, nepheline exists through the 2325~ test temperature; however, it was not found in any of the samples at 2625~ Those refractories higher in silica content developed more nepheline than high-alumina products. In fact, nepheline did not exist in either the 70% high fired or 90% alumina materials. If either the amount of soda increases or the silica content of the mixture is low, a sodium aluminum silicate forms. Most probably this compound has the sodium-to-alumina-to-silica ratio of 3 : 2 : 4; however, it could be the 2 : 1 :2 compound. Those samples lower in silica content develop higher amounts of this phase. The amount increases with increase in temperature in all the refractories tested except in the superduty sample. Rather, this product shows a decrease in the quantity of that phase as a result of the formation of glass with an increase in temperature. The fusion point of this glass was reached at the highest test temperature of 2625~

Above 2000~ the authors observed the presence of beta-alumina. In the higher alumina products, the amount was still on the increase at the highest test temperature. Unexplainably the quantity of beta-alumina decreased above the 2300~ in the 70% alumina sample. Furthermore, a decrease in beta-alumina occurred above this temperature in the superduty product. This was due to the increased formation of glass. The magnitude of these decreases was not discussed by the authors. However, their plot of relative XRD peak intensity vs. temperature clearly indicates that the decrease is greater for the 70% alumina product than the superduty sample.

The authors summarized the results of the soda reaction test as follows: For the high-silica samples the reaction products were identified as nepheline and sodium aluminum silicate (3 : 2 : 4). The amount of nepheline decreased with increasing alumina content, whereas the amount of (3:2 : 4) compound increased. The lack of silica in the high alumina samples permits the free alumina not tied up as a sodium aluminum silicate to react with the alkali to form beta-alumina. Higher amounts of soda produce sodium a|uminates.

The authors described the reaction mechanism as occurring by impregnation of the sample, bond reaction, and bond depletion. Immediately the alkali consumes the cristobalite and attacks the glass of the matrix bond. Next the fine crystalline mullite bridges associated with the bonding matrix react with the soda. The mullite x-ray peak intensity is quickly reduced. The extent of attack on the coarser crystalline mullite is intensified at higher temperatures or as the reaction proceeds. When equilibrium


between the alkali and the available cristobalite, glass, and mullite has been achieved, the free alumina reacts with alkali forming beta-alumina. This phase develops at temperatures above 2000~ and exists through about 3000~ An increase in the amount of soda will result in the formation of sodium aluminates.

The formation of the above soda compounds is accompanied by a corresponding change in the mineralogy of the refractory. Initially a very rapid depletion of cristobalite occurs with increasing temperature. By 1575~ the cristobalite has totally disappeared. The sequence of attack continues as the mullite is consumed, the amount decreasing with increasing temperature. For the high-fired samples, the mullite exists longer at higher temperatures. This is due to the larger quantity of coarse crystalline mullite in the original sample.

The change in mineralogy is dependent on the alumina content of the sample. For the lower alumina products, the alpha alumina content remained essentially the same. However, a slight increase in the alumina content of certain samples, the superduty and the 60% alumina products, was observed at the lower test temperature.

The 90% alumina refractory showed a sharp drop in the amount of alpha-alumina present as the temperature was increased. The depletion of alpha-alumina does not occur unless excess soda is available and the temperature is above 2000~ or unless the amount of SiO 2 present is not sufficient to tie up the soda as a sodium aluminum silicate. Finally, the free alumina existent in the lower silica-containing samples is depleted by the formation of beta-alumina and alkali aluminate.

The work conducted by Yamaguchi [28] deals in part with the effects of sodium carbonate vapor on the minerals corundum and mullite. In the case of corundum, the major reaction product was apparently sodium aluminate, whereas with mullite, carnegieite solid solution was observed.

In the second phase of his study, Yamaguchi [28] investigated the corrosion of certain burned refractories by sodium carbonate vapor. He suspended the test piece with a platinum wire from the bottom of an alumina crucible placed upside down. The entire assembly was heated at 1200~ for various times. Included in this part of the study was a fireclay refractory composed of mullite and silica minerals. Mullite reacted with Na20 to form nepheline and alumina. The nepheline increased in amount as the remaining soda vapor reacted with the newly formed alumina and the preexistent silica. For refractories composed of mullite and corundum, carnegieite solid solution was the major reaction product. The formation Yamaguchi described resulted when NazCO 3 vapor reacted with the alumina liberated from mul|ite and preexistent as corundum, forming


sodium aluminate. The sodium aluminate thus formed dissolved into nepheline resulting from the reaction between mullite and soda.

Studies performed by Rigby and Hutton [29] yield essentially the same results as reported by Yamaguchi. Their experimental procedure consisted of fabricating six samples ranging in composition from 100% silica-0% alumina to 0% silica-100% alumina with increments of 20% silica and alumina. Each fired silica-alumina specimen was ground, and 100 parts of each specimen were mixed with 42 parts of sodium carbonate. The mixtures were pressed into cylinders and subjected to successive 3-h heat treatments at temperatures beginning at 800~ and continuing to 1600~

The two samples comprised of 100% and 80% silica, respectively, formed sodium silicates that melted near 800~ The authors pointed out that with the sample containing 60% silica-40% alumina, the soda first combined with cristobalite to form sodium silicate and then reacted with the mullite to form nepheline and liberate corundum.

On the other hand, the specimens with a silica/alumina ratio of 40:60 consisted largely of mullite and formed a sodium aluminate rather than a sodium silicate. First, the soda attacked the excess alumina to form a sodium aluminate. As the reaction with soda proceeded, the mullite decomposed, forming nepheline and more sodium aluminate.

The specimen corresponding to an 80% alumina brick, ratio 20:80 silica to alumina, had the mineralogical composition of mullite and corundum. The resulting phases present in this sample when mixed with the specified quantity of NazCO 3 were sodium aluminate and corundum.

On a laboratory-prepared specimen comparable to corundum brick, the presence of beta-alumina in addition to sodium aluminate was noted. However, beta-alumina was formed above 1300~ Unlike Yamaguchi [28], Rigby and Hutton [29] did not differentiate as to which was the major or minor reaction point.

Rigby and Hutton [29] concluded that for refractories with a silica/alumina ratio greater than 1.5, shrinkage may occur at the hot face of the brick. The higher the silica content, the greater the shrinkage and the lower the liquidus temperature. On the other hand, expansion occurs for refractories with a ratio less than 1.0 which may be attributed to the formation of either sodium aluminate or beta-alumina.

Similar to the procedure utilized by Rigby and Hutton, Lambertson [30] ground a 60% alumina-diaspore, fireclay brick and mixed this powder with sodium sulfate. The mixture was heated in a platinum crucible to 1300~ X-ray and microscopic examination showed the decomposition of mullite to form glass and corundum. As the reaction proceeded, nephelite replaced the glass.


Under equilibrium conditions as soda is added to a 60% alumina-40% silica composition, albite should begin to form. With an addition of approximately 6.4% soda, all the silica and mullite of the refractory is changed to albite and corundum. Grieg and Barth [31] in their studies of the system nephelite-albite, however, encountered extreme difficulty in crystallizing albite. With additional soda, Lambertson stated that nephelite begins to form in equilibrium with corundum, and with more than 16.5% soda the corundum reacts to form N a 2 0 - AlzO 3.

2. ExeosuRE TO SODA VAPOR IN GLASS MELTING FURNACES

a. Reactions on Silica Brick. Insley [32] approached the subject by examining the refractories used above the glass melt. In this zone of the furnace, the refractories are exposed to attack by sodium compounds present in the furnace atmosphere as either gases or dust. The exposed surfaces of the silica brick used in the crown and side walls of the tank had a vitreous luster, which when viewed under the microscope showed a considerable content of glassy material. However, little corrosion of the brick itself was observed. Insley surmised that this phenomenon was a result of either the liquid being so viscous or so small in amount that it did not drain from the brick.

On the surfaces of the silica brick utilized in the side walls, Insley observed an incrustation with a white, coral-like appearance. This "frost" was easily broken away from the surface; consequently, very little evidence of corrosion of the silica brick was noticed. A petrographic examination of the "frost" revealed very small tridymite crystals in the part of the brick next to the frost. The contact material between the brick and the incrustation had a vitreous appearance composed of glass and large tridymite crystals. Further within the incrustation, the crystals were larger, but the quantity of glass was smaller.

Insley [32] believed that the soda present in the furnace atmosphere combined with the silica to form a liquid sodium silicate. In turn, this liquid acted as a crystallizing medium for the tridymite. The continued growth of the new large tridymite crystals forced the previously formed crystals away from the face of the brick.

This was verified in a subsequent study by NARCO [25] in which a more detailed explanation of the mechanism of frost formation was given. It is well known that silica bricks exposed to soda vapor develop a liquid phase at around 1450~ The amount of liquid in the brick will increase as the amount of flux is increased. The liquid thus formed is gradually absorbed into the interior of the brick until the voids are filled. On the


surface the liquid phase promotes the development and growth of tridymite. The tridymite crystals left on the surface continue to grow, the ability of the silica brick to absorb more liquid is reduced as the voids become filled. The liquid remaining on the surface absorbs more soda, causing a gradual loss of the residual tridymite crystals. These crystals become deeply corroded. The embayments thus formed intersect and a skeletal structure (frost) remains.

A study of soda ash foundry slag attack on the refractory lining of a teapot-type, reservoir ladle by Dear [33] shows that siliceous fireclay ladle bricks are rapidly eroded and corroded during service. Optical microscopy of the used brick reveals the presence of four layers: (1) a thin, inhomogeneous glass layer at the exposed surface of the brick, (2) a transition layer at the slag-brick interface, showing copious mullite development in a glass phase, (3) a greatly darkened vesicular layer immediately behind the transition layer, followed by (4) the main body of the brick, darkened but otherwise not greatly altered [33].

In the teapot-type reservoir ladle soda ash is used as a purifying agent. All of the reactions between soda ash and molten metal are not known although literal quantities of CO gas are evolved at the slag line. The presence of CO may partially explain the darkening of the brick because CO would act as a reducing agent toward the iron compounds in the brick.

Dear [33] described the mechanism of attack as the solution of first the glass, cristobalite, and other forms of silica in brick, followed by the solution of the finely divided mullite and clay body. The presence of well-crystallized mullite needles at the slag-interface indicated that the liquid at the contact surface was enriched with alumina.

b. Reactions with Clay Refractories. Insley [32] observed that the action of dust and vapors in the furnace is more detrimental to clay than to silica refractories. Wilson [34] noted soda combining with the alumina and silica in the clay refractories to form a liquid at operating temperatures. The liquid was found to have the approximate composition Na20 �9 A1203 �9 5SiO 2. Any excess of alumina over the amount required to form the liquid will crystallize as corundum. On the other hand, if there is an excess of both silica and alumina, the crystalline products will be either mullite and corundum or mullite and silica (tridymite or cristobalite). In the majority of cases examined, only corundum crystallized.

Insley [32] analyzed the clay refractories above the glass line in the melting end of the tank and in the refining end. Each region showed different results characteristic of the temperature and conditions at that particular location.

The interior of the brick in the melting end was composed of mullite and glass. The size of the crystals was dependent on the distance from the


exposed surface. Near the exposed surface, the amount of mullite decreased and the amount of glass increased. Closer to the surface, mullite disappeared completely and corundum appeared. The surfaces of the brick were covered with a brown and white deposit, which curled in waves. The brown color was probably due to the presence of iron oxide either as inclusions or as solid solution.

In the refining end of the tank, the effect of soda on the clay refractories resulted in the formation of other substances in addition to the sodium alumina-silicate glass and corundum. In one tank examined, stalactites were found hanging from the clay refractories. The principal crystalline phase of these stalactites was nephelite ( N a z O . A 1 2 0 3 . 2 S i O 2) with carnegieite (high-temperature form of nephelite), melilite (2CaO �9 SiO 2 �9 A1203), and aegirite (Na20 �9 FeO �9 SiO 2) present in lesser amounts. The corundum crystals formed were very small and had undergone solution. In every instance, the corundum crystals were surrounded by nephelite.

Insley [32] speculated that the first substances formed were corundum and glass. As the soda content increased, reaction with the glass occurred to form nephelite or carnegieite depending on furnace temperature. He found it difficult to explain the fact that nephelite formed in the cool end of the tank (refining end), while a glass lower in soda formed in the melting end where presumably the atmosphere contained more soda. He offered as a possible explanation that nephelite could only form at the lower temperature. Thompson and Kraner [35] also commented on the effect of alkalis on clay refractories used in glass melting. The sodium vapor and dust react to form either nephelite or to drain from the refractory a liquid phase richer in silica. The result is that nephelite and coarsely crystalline mullite appear. Additional fluxing action by soda leaves corundum [3]. Specifically, Thompson and Kraner found that the extreme outer surface of a clay brick that had been exposed to furnace fumes and temperature on a bridge wall is composed of corundum. Even the glass has been drained out of this portion. The next or innermost portions of the brick consist of mullite. The corundum is the result of the decomposition of mullite, which could not have occurred at the prevailing temperatures without the presence of the alkaline flux. The extreme outer portion of the brick is porous due to the rapid solution and removal of the siliceous material. Nephelite, the authors stated, generally forms as a transition layer between the refractory and the glass. Thompson and Kraner could not determine whether nephelite exists as a layer or only as a pasty liquid during furnace operation; however, the mineral was found between the two in all alumina-silica refractories.

Turner and Turner [36] conducted studies on the corrosion of fireclay from four different sources. Although their work deals with clays rather


than firebrick, it does provide some insight into the temperatures at which the brick would be affected. They pointed out that at 800~ sodium carbonate attacks fully burnt fireclay, even though this temperature is below the melting point of the carbonate. This essentially agreed with the work of Cobb [37] in which he observed that the reaction between sodium carbonate and silica occurred at temperatures as low as 690~ with alumina, reaction was noted at 720~ This reaction was accompanied by an evolution of carbon dioxide.

Turner and Turner [36] found the attack on the clay fired at 1400~ to be less corrosive than on the same clay fired only at 1300~ A rapid increase in the extent of corrosion occurs with an increase in test temperature. Between 800 and 900~ the rate of attack increases severalfold. The resistance to corrosion was found to be dependent on the density of the clay body as well as on chemical composition.

Petrie and Brown [38] investigated superduty firebrick from the top courses of the glass tank checkers. Parts of these brick were completely eroded, leaving a glazed surface on the remaining portion of the brick. The glaze was thin, indicating the low viscosity of the glass phase during service. It was separated from the unaltered portion of the brick by a thin white porcelain-like layer. Their observations concerning high-duty firebrick agree with those made by Insley [32].

A later study by Insley [39] on the surface deposits formed on high alumina refractories used in glass regenerators revealed that the interior of the brick was composed of mu|lite corundum and a small amount of glass. The white vitreous layer between the interior of the brick and the altered surface contained well-developed mullite, glass, and very few corundum crystals. The outer zone contained predominantly nephelite, with small corundum crystals present as inclusions. At operating temperatures below 2700~ carnegieite formed a protective coating on the surface of the brick.

Results obtained by Thompson and Rexford [40] concur with the findings of Insley [39] with respect to the development of nephelite and carnegieite. Specifically, Thompson and Rexford set standard 9-in. brick of 49% alumina (high-duty), 50% alumina, and 60% alumina on the floor of a regenerator port in a glass tank. The temperature was approximately 2400~ and the samples were removed after about six months. The gases leaving the tank at a relatively high velocity were laden with dust and vapor and were impinged directly on the test straights. Visually two effects were observed: (1) a more fluid slag apparently formed on the fireclay brick than on the 60% alumina specimen, and (2) the amount of erosion was less severe for the two types of high alumina than for the fireclay brick.


Currier [41] observed that when these batch dusts and vapors come in contact with the surface wet by the internal liquid developed in a brick at high operating temperatures, the dusts adhere to this face. The alkalies and other bath constituents then react with the alumina and silica in the brick to form additional low melting compositions. If the temperature is sufficiently high, the liquid will drip onto the brickwork below. This effectively keeps the attacked brick clean, thereby offering new surfaces for the carryover. In this manner, the refractory is gradually consumed. In the lower courses where the temperature is not high, the dust collects, forming a hard crust on the brick and thus constricts or clogs the flue openings.

For superduty brick the lower alkali content and higher alumina content than high-duty brick results in the formation of less liquid at a given temperature. Batch dusts also react less readily with refractories of this class. The extra-hard-burned brick, in general, serves better due to their lower porosity, higher strength, and greater volume stability. In addition, high firing tends to enhance mullite growth or crystallinity, which would improve chemical corrosion resistance. Therefore, the use of a superduty brick maintains the original flue lines, unlike high-duty straights, which often sag or form a hard crust.

c. Reactions with High-Alumina Refractories. High-alumina refractories should resist the fluxing action of soda better than clay refractories. If the silica content is low, then the amount of sodium alumino-silicate liquid formed would be small compared to the amount of excess alumina. The liquid would be an efficient medium in which corundum could crystallize but would not be so great in amount as to flow out of the refractory. The heavy crust of corundum formed on the surface thus acts as a barrier against further attacks of the soda.

On the other hand, Thompson and Rexford [40] reported that for high-alumina brick at temperatures around 2200~ the liquid flow dimin- ishes to the point that sufficient time is available for the liquid on the surface to become rich in alkalies. Nephelite thereby forms in the surface layer causing slabbing of the brick surface. If reactions are completed, only about 4% to 5% soda need be added to high-alumina compositions to consume mullite and produce corundum, nephelite, and albite. With an addition of 12% soda, the phases present are corundum and nephelite.

_

Therefore, with higher amounts of alumina, the tendency to form a slag with the batch dusts or soda vapor is reduced, however, the possibility of slagging or peeling becomes a consideration. For example, Thompson and Rexford [40] found that for the hottest zone of checkerwork, slabs more than 0.5 in. thick separated from the faces of the brick. The slabbed portion of the brick was found to contain 10% soda and to be composed


predominately of nephelite. Numerous corundum crystals as well as scat- tered patches of mullite through the ground mass also existed in the slabbed layer. Based on chemical analysis, calculations of the probable mineral composition indicated that the slabbed portion contained 57.7% nephelite, 20.9% mullite, 11.9% corundum, and 9.8% miscellaneous constituents. Microscopic examination, however, showed that more corundum and nephelite and less mullite were present.

These reactions would be expected to cause a permanent volume expansion inasmuch as nephelite has a lower true specific gravity than the weighted mixture of the original corundum, silica, and mullite. The bulk volume increase of about 14% was associated with the change in mineralogy. This fact, coupled with the increase in thermal expansion of the outer layer, which contained alkalies, results in slabbing or shelling of the surface.

D. Reactions with Potassium Carbonate Vapors

1. EXPOSURE TO POTASSIA VAPOR IN A BLAST FURNACE

The service requirements utilized in a blast furnace vary considerably at different locations within the furnace. Refractories are subjected to variations in temperature and pressure as well as to different gases and corrosive agents. For example, the temperature is relatively low in the upper part of the stack but increases in the tuyere and hearth zones. A marked thermal gradient occurs from the inner surface of the lining through to the shell of the furnace. The blast-furnace atmosphere consists of reducing gases, volatile metals, and alkalies, which may collect in the cooler part of the lining and cause disintegration of the refractory [42].

In the blast furnace potassia has been found to be a major destructive component in wear processes. The descending charge not only carries K20 and Na20 in the ore and coke (ash), but also any K20 , Zn, or coke ash (coke burned at tuyere level) that was deposited on the burden from the ascending gas stream. At the tuyere zone most of the soda and some of the potassia remaining in the slag are flushed from the furnace. However, most of the potassium reacts with coke or CO to form volatile metallic potassium compounds [43]. Under certain conditions potassium can react with CO and CO 2 to produce K20 or even K2CO3, causing the subsequent deposit of carbon and the freezing of K20, which combines with the oxides of the refractory [44].

Alkali minerals have been found on the inside surface and on the porous parts of the bosh and stack lining. The alkali content of the inside


surface of the stack ran as high as 31. No disintegration of the brick was observed in this zone. Rather this zone may be identified by the vitreous appearance of the brick [42, 44, 45].

Apparently the surface glaze on the superduty and high-duty brick probably resulted from a rapid reaction between the brick and the K zCO 3. The reaction products become liquid at operating temperatures yet retain a high viscosity. Therefore, this viscous liquid seals the pores on the face and prevents the molten K zCO 3 from penetrating the underlying portions of the brick. On the other hand, the KzCO 3 enters the pores of the brick, reaction with the alumina-silica components occurs, giving off CO 2 in the process. The reaction material was often vesicular and accumulated almost immediately on melting of KzCO 3 [45].

In the center of the lining Van Vlack [42] detected the absorption of 8% alkali. The high porosity of this region was attributed to carbon disintegration of the refractory. For the inside of the bosh area a coating of carbon and alkalies was observed with the sample containing as much as 42% alkali. The refractories involved are in the 40% to 45% alumina class.

Cracks noted by Lister [46] in the unreacted portion of the brick were the result of mechanical stresses in the lining due to thermal and structural causes. However, Lister reported that all the cracks in the hot end of the furnace, particularly those in the alkali attack zone, were not relict mechanical cracks. Many of these must have formed from mechanical stress on brickwork weakened by alkali attack. Lister hypothesized that the first factor affecting lining life is the formation of cracks due to heating up stresses; next comes alkali penetration and attack, which is believed to be the most important event governing lining life. Alkali attack appears to weaken the lining to such a point that the brickwork becomes more prone to scouring by the lime-rich slag. Alkali attack can occur in two or possibly three ways:

~

.

Straightforward reaction with the alumino-silicate constituents of the refractory, to produce expansive-type phases such as kalsilite (K 20 �9 A120 3 �9 2SIO2). Considerable volume increases accompany these reactions, thus causing weakening and disintegration of the brick.

The solution of the brick material to produce a glass with a different thermal expansion than the refractory. In many of the samples Lister examined, this glass forms the matrix bond of the refractory in this zone. Consequently, if the lining is subjected to a temperature change, the glass may crack causing weakening or disintegration of the brick.

2 PREDICTION OF ALKALI OXIDE CORROSION 7/

3. In many of the samples in the alkali-rich zone the alkali literally fills all pores in the structure. Thus, one possible mechanism of alkali disruption may involve alternative freezing and thawing of this alkali. During cooling, alkali freezes and then cracks, which forms voids. If more molten alkali penetrates into these voids as the temperature of the hot face is increased, then all voids are filled before the frozen alkali can melt. When the original alkali does melt there is no place for the alkali to migrate or expand thus structural disruption may result.

The alkali-bearing minerals found in the blast furnace include kaliophilite-nepheline, leucite, plagioclase, and alkali carbonates [40-44]. Kaliophilite was the dominant reaction product on the inside surface of the lining and in the joints between brick. The formation of this compound is accompanied by about a 45% volume increase. No kaliophilite, however, formed in the interior of the brick except in the carbon-disintegration

zones. Van Vlack [42] found leucite immediately adjacent to the kaliophilite

zone in the brick. In general, leucite formed in more siliceous areas between grog grains of the brick. Volume expansion was also encountered with the formation of leucite although the amount is less than for kaliophilite. Most likely the pore volume of the brick was sufficient to accom- modate the leucite produced. Van Vlack concluded that the nature of the boundary between the grog and leucite zone and the kaliophilite zone indicated that the first alkalies to enter the brick formed leucite. The absorption of more alkali apparently produces kaliophilite from the existing leucite and mullite in the grog grains. Leucite was sometimes observed within the grog grains with mullite and corundum. The instability of mullite with keliophilite was further demonstrated in a recent work by Schairer and Bowen [47]. In the K20-A12O3-SiO 2 system, Schairer and Bowen showed that leucite and corundum are mutually stable.

Van Vlack [42] observed that the bosh region contained alkalies in excess of the requirements of kaliophilite and nepheline. This excess did not result in the formation of alumino-silicate or silicate minerals of higher alkali content but was present as alkali carbonates. The carbonates formed toward the end of the furnace campaign from the free alkali deposited in the bosh lining. For the particular furnace examined, Van Vlack determined that alkali-bearing minerals were present only where porosity and proximity of the surface permitted expansion. Therefore, the presence of alkalies did not prove seriously detrimental in the bosh and


stack lining. Where disintegrated brickwork contained alkalies, the destruction was also associated with the existence of carbon.

To examine the mechanism of peeling of fireclay brick in the low-temperature region of a blast furnace, McCune et al. [45] investigated the mineralogical changes in used brick from a blast furnace. In addition, their work included laboratory tests using mixtures of crushed refractories (high-duty, superduty, and 70% alumina) and potassium carbonate. Test temperatures for the superduty mixtures were at 1500 and 2200~ while the mineralization of crushed high-duty and 70% alumina samples was studied at 900, 1100, 1400, and 1700~

In general, the authors found that the constituents of used blast furnace linings in the upper or lower temperature regions were the same as the reaction products that formed in mixtures of crushed firebrick and alkali. McCune et al. [45] determined that the initial alkali contact, the potassia content being 0% to 5%, resulted in a decrease in cristobalite. Next leucite was observed to coexist with some cristobalite, with the potassia content increasing to between 5% to 10%. Therefore, the peeling of a fireclay brick in the upper regions of the stack is undoubtedly related to the formation of leucite and kaliophilite, which occurs with 5% to 10% K20. The peeling of a refractory in this area of the blast furnace seems to take place as alkali vapor or liquid penetrates a short distance into the brick. With the absence of a silicate liquid, reaction with the solid bond constituents occurs. Below 1500~ alkali affects the cristobalite content of the groundmass, or bond. A disruption of the bond would involve considerable stress or loss of strength for the refractory. The occurrence of leucite and kaliophilite would cause further stress on the bond at the interface between the unaltered and altered portion of the brick. How- ever, if a silicate liquid is present then, as observed by Van Vlack, reaction between K20 and silica can take place at a sufficiently low temperature to seal the surface pores and thereby limit the amount of alkali penetration [42, 45].

In the upper stack regions of the blast furnace, McCune et al. [45] reported that peeling was found close to the maximum depth of penetration. On the other hand, in the lower stack region as well as the lower hearth region, fractures near the limit of alkali penetration are rarely encountered. This, McCune et al. attributed to the formation of a silicate liquid, which prevented the development of breaking stresses. In the higher temperature regions of the stack, peeling occurred where 20% to 30% K20 was present. Evidently, the higher temperatures resulted in the depletion of the silicate liquid in forming kaliophilite. Above 1600~ and for a K20 content of less than 15%, the adverse effects of kaliophilite


formation would be reduced if enough liquid is developed to eliminate peeling stresses. McCune et al. concluded that a high-silica/low-alumina dense brick would tend to react with K 2 O to surface seal and thus prevent alkali destruction of the brick by peeling. Their studies of 70% alumina- K2CO3 mixtures showed a rapid formation of kaliophilite, which indicates that high-alumina refractories would be subject to failure by peeling in blast furnace applications.

Farris and Allen [27] briefly summarized the role of potassia in blast furnace wear processes as follows:

1. In high-fired superduty brick, initially the cristobalite reacts at about 1100~ and is consumed at 1400 to 1700~ depending on the K20 concentration. With the presence of an appropriate amount of K20, mullite is consumed at temperatures as low as 1700~ Corrosion products formed from these reactions are generally leucite (K20" AI20-4SiO 2) and kaliophilite (K20" A120" 2SIO2).

2. The reaction phases are similar for the 70% alumina brick exposed to K20; however, corundum and potassium aluminate are also produced by potassia reaction.

3. Basically the mechanism of attack is alkali impregnation and reaction accompanied by expansion.

2. REACTIONS WITH MULLITE BRICK

Laboratory tests were conducted by Yamaguchi and Okawara [48] in an attempt to identify the processes of corrosion of mullite refractories by K z CO 3 vapors. A small amount of leucite and kalsilite formed at the initial stage of alkali attack. Gradually these phases were replaced by the final products of potassium aluminate and a glassy phase. A later study by Yamaguchi [28] as well as a subsequent investigation by Thomas [49] confirmed the presence of potassium aluminate in a corroded mullite test sample. However, both papers reported the formation of a potassium alumino-silicate though neither specified whether the compound was leucite or kaliophilite. Thomas observed that more porous brands of mullite showed a high degree of flexure, which resulted in severe cracking and spalling in a potash environment.

The work by Boyer [50] was concerned with mullite refractories utilized in blast-furnace applications. The major reaction appeared to be KzCO 3 fluxing of the mullite-glass samples to form minor kaliophilite. Boyer described the reaction as a soft black hygroscopic crust confined to the


exposed surface of the brick. A shallow zone of minor phase alteration consisting of a dark discoloration of the refractory glass accompanied the surface attack. Boyer attributed the metal inclusions within the discolored glass to the reduction of the refractory titanate. The mullite phase showed no alteration. It should be pointed out that the test temperature of 1725~ would not be adequate to produce potassia corrosion or reaction with mullite.

III. Predictions of Alkali Corrosion of Alumino-Silicate Refractories Using Phase Diagrams

The common refractory compositions in the AI203-SiO 2 system were indicated in Fig. 1. These compositions also occur in the boundary A1203-SiO 2 system of the Na20-A1203-SiO 2 ternary diagram shown in Fig. 2. A simplified version of this ternary diagram is shown in Fig. 7.

Cristobol

Si02 1723 A ~59~

Tridymite 789,~

Albite

Nepheline

Cornegieite

Mullite o~

Corundum

IzO~-2SiO z 1850

1840

lumi

NotO AIz03 No20-AIzO 3 1580 N~ 2040

1650 1980

FIG. 7. Simplified version of the Na20-AI203-SiO 2 ternary phase diagram.


Na~,O

Liquid _,~ 1800oC . ~ M u l l i t e + L

/I -" ~ M u l l i t e + Corundum+L I I I I~ ~--- - -Corundum + L

Carnegieite + L - - ~ I ~ T 4 0 0 ~ L _ _ ~ - / I / ~ Tridy mite + Mullite+L

Carnegieite+ Nepheline + [_ ~ _ . . _ _ - - Albite + Mull i te +L I I -I IO00~

~ J f - J ~ = Albite + Nepheline +Cor. J [ - ~ ~ Albite + Mullite +Corundum I I ~ ~ Mullite + Albite +Quartz

= I i, = 6 0 0 o c 20 I0

Firecloy

FIG. 8. The NazO-fireclay vertical section of the Na20-A1203-SiO2 system.

When an alkali such as Na20 reacts with an alumino-silicate refractory, the reaction products that form will be described by the Na20-refractory vertical section. The vertical section can be constructed by tracing the crystallization paths for a series of ternary compositions that lie on the Na20-refractory line. For example, the low Na20 portion of the Na20-fireclay vertical section is shown in Fig. 8. As Na20 reacts with fireclay, the alkali content will gradually increase and, for any given temperature, the phases forming in the refractory (generally near the surface)will be described by an isotherm in the vertical section. If the corrosion takes place over a long enough time, the corroded portion of the refractory will experience a large alkali concentration gradient. This will create a series of reaction products that is predicted by the vertical section. If corrosion is from the surface of a refractory, then the surface will have a high alkali concentration and the concentration will be progres- sively lower deeper into the refractory.

The Na20-fireclay vertical section contains a tremendous amount of information that describes the sequence of reactions that occur at all temperatures up to and including the melting point of the refractory. (Remember that the phase diagram describes what will happen but gives no information about kinetics. The phase diagram will also give us quantitative information, but this must be gathered from the ternary diagram, not the vertical section.) As can be seen from the vertical section, very small amounts of Na20 cause liquid to form in fireclay at temperatures as low as 1063~ A fireclay normally contains mullite and various polymorphs of silica; however, as it reacts with Na20, other crystalline phases form at relatively low Na20 concentrations. Depending on temperature and the Na20 concentration, corundum, albite, carnegieite, and nepheline


can form. Also note that at about 18% Na20 all liquid recrystallizes up to temperatures as high as 1600~

Now that the phase diagram has told us what will form during Na20 reaction with a fireclay refractory, we must evaluate whether or not these phases create a serious problem. This information can be obtained from various sources, especially any text on high-temperature silicate chemistry. The important information follows:

1. If the SiO 2 content of a silicate liquid is high, the liquid will be very viscous, even at high temperatures. The resulting glass will have a relatively low thermal expansion coefficient. In other words, a high- silica liquid may effectively glaze the refractory and reduce the rate of alkali reaction. The high-silica glass will not be prone to cracking from thermal shock.

2. As the Na20 content (and to a lesser degree the A120 3 content) increases, silicate liquids become increasingly more fluid. This means they will be affected by gravity and not provide a pseudo-protective glazing effect. Also, high Na20 silicate liquids form high thermal expansion glasses that will be prone to thermal shock and spalling. If the Na20 content is moderately high, the glass Will also be hygroscopic.

3. The crystalline phases~corundum, albite, nepheline, and carnegieite --al l possess a high thermal expansion coefficient. (This is typical of most alkali compounds.) As these phases form in increasingly higher amounts, the thermal shock resistance of the reacted portion of the fireclay refractory will decrease significantly.

In summary, we can conclude that very small amounts of Na20 will cause liquid to form in a fireclay refractory above 1063~ If the liquid is confined to the surface of the refractory, a glazing effect will prevent serious damage to the refractory. If Na20 concentrations are significant, even a fraction of 1 wt %, a fireclay refractory will be prone to erosion at high temperatures and serious thermal shock at low temperatures. Fire- clay refractories would seem to have poor-to-moderate resistance to Na 2 O.

Similar analyses can be made for the other alumino-silicate refractory compositions, which are summarized below.

Silica compositions: It is well known that Na20 reacts rapidly with high-silica refractories and this is reflected in Fig. 3. Ten percent Na20 will form 50% liquid at temperatures as low as 1100~ High-silica refractories are generally not used in alkali environments except at low temperatures.


Na20

1800

Cristobalite + L Cristobalite + Mullite + L

Tridymite Tridymite + Mullite + L

Tridymite + Albite + L

Tridymite + Albite + Mullite

Quartz + Albite + NS 2 I , 600

20 10 5 AI20 3 95 SiO 2

FIG. 9. The Na20-semi-silica vertical section of the Na20-AI203-S iO 2 system.

Semi-silica compositions: The Na20-95% SIO2/5% A 1 2 0 3 vertical section (Fig. 9) shows that as Na20 reacts with a semi-silicate refractory, all SiO 2 phases should disappear, and tridymite, the only stable form of silica, will form very slowly. Albite will form only at temperatures below 1050~ and liquid is present down to 740~ Semi-silica refractories are not resistant to alkali attack, and may be even less alkali-resistant than high-silica refractories.

Mullite refractories: Figure 10 shows the NazO-mullite vertical section. The phases formed are corundum, ~-A1203, and nepheline. In both cases, the mullite phase decreases in amount as Na20 reaction increases. The

N a A I O 2 + L ~ ~ " ~ . - ~ - - - ~

Nephel ine+Corundum+L~ I | r

NaA I 0z+/3 -A la03+Carnegiei te~~ \ /

20 I0 Na20

=/•Mullite § Corundum §

Mullite + Corundum +L 1600 ~ . 8 -A Iz03+Corundum +L 1400

=2,~ Corundu m ~AI bite §

-A I203-~Corundum + Nepheline

~ Corundum +Nepheline + coo Albite

---Corundum+Mull i te+ Albite

Mullite

FIG. 10. The Na20-mull i te vertical section of the Na20-A1203-SiO 2 system.


13 . . . . Corundum + L -~,2U3 + L ~ j 13-AI203 . Corundu.m +

NA * L ~ ~ , ~ ~ ~ ~ 1 Na + ~-AI20 a + L ~ J

Na20

Na + 13-AI203 + Carnegieite -

.••'•180110 - Corundum + L it - ' ' ' I Mullite + L

, 00

13-AI203 + Carnegieite + Corundum Carnegieite + Corundum + Nepheline -

Corundum + Mullite + Albite ----

20 10

Carnegieite + Corundum + L

95 AI203 5 SiO2

F~G. 11. The Na20-95% AI203 vertical section of the Na20-A1203-SiO 2 system.

overall diagram looks similar to the Na20-fireclay section, except that /3-A1203 and Na20 �9 A1203 are now formed at high alkali concentrations. Moreover, as the refractory composition becomes higher in A1203, the amount of liquid formed decreases sharply. On cooling, the liquid recrystallizes rather than forming glass. Again it can be noted that after about 15% alkali, refractoriness increases.

2000

1800

1600

1400

1200

I 000 60 70 80 90

N~ Mol % AI203

FIG. 12. The high-alumina portion of the Na20-AI20 3 phase diagram [51]. Reprinted by permission of the American Ceramic Society.


High-alumina refractories: Figure 11 shows the Na20-95% A120 3 vertical section. Mullite, corundum, and glass are reduced or eliminated in favor of /3-A1203, Na20 �9 A1203, and nepheline as the alkali reaction proceeds. Refractoriness is increased after reaction with only 10% alkali, significantly less alkali reaction than that required for increased refractoriness in mullite.

Corundum refractories: As expected from the NazO-AI20 3 phase diagram in Fig. 12 [51], corundum refractories should be very resistant to soda attack. No liquid should form below 1580~ however /3-A120 3 should form at all temperatures. The formation of /J-AI20 3 produces a volume increase of 18.5% compared to corundum. Obviously corundum refractories are resistant to chemical attack by soda vapors but not to the chemical spalling produced by soda. This leads to the well-known conclusion that /3-A120 3 refractories, rather than corundum refractories, perform best in alkali environments. Also it is clear that the refractoriness of /J-AI20 3 is much less than that of high-alumina refractories.

IVo Comparison of Predicted Na 20 Corrosion of Alumino-Silicate Refractories with Experimental Results

Hayden [52] conducted a controlled series of alkali corrosion tests of commercial refractories using the "crucible cover" technique illustrated in Fig. 13. This test involved placing a refractory sample over a crucible containing a premeasured amount of alkali carbonate and heating the assembly to 1350~ for 12 h. A furnace containing a load of test samples is

SAMPLE

FIRECLAY CRUCIBLE

J I

' ' [ " \ ' ' i., ALUMINA 1... :k 17. c,uc,stt

\ ~!!~ ,~:~..:.~/\ ~ii!i~, ,~!i;l / TABUtm -\ ~!?': :~:~:iI / \ ~ii.>:...:: .,?.:.~-/--~u.~NA

FIG. 13. The alkali "crucible cover" testing assembly.


shown in Fig. 14. After the test, the samples were photographed and the mineralogy of the reaction areas was determined.

The photographic results of Hayden's study are shown for semi-silica, fireclay, mullite, high-alumina, and corundum refractories in Figs. 15 through 19, respectively. As can be seen by comparing these figures, the phase diagram prediction was accurate. High-silica refractories exhibited severe corrosion and extensive liquid formation. As the alumina content increased, the corrosion was less severe. Mullite and AlzO3-rich refractories show no glass formation; the liquid that formed recrystallized on cooling. Extensive cracking of the reacted region of the refractory was noted. Finally, the corundum refractory showed extensive volume expansion and cracking, as was to be expected from the formation of /3-A120 3.

The mineralogy of the reacted areas of the refractories showed excellent agreement with the phase equilibria predictions. The high-silica refractories contained quartz, cristobalite, large amounts of liquid, and trace amounts of mullite. The fireclay refractory contained liquid, mullite, and carnegieite. Nepheline was not observed nor would it be expected to be present at temperatures above 1050~ The mullite sample contained mullite, corundum, nepheline and 13-A120 3. The high-alumina refractory contained mullite, corundum, /3-A120 3, nepheline, N a 2 0 . A1203, and evidence of recrystallized liquid. Finally, the corundum refractory was found to contain large amounts of 13-A120 3 in addition to corundum.

V. Conclusions

(Because of space limitations it is impossible to treat the K2 ~ corrosion of alumino-silicate refractories here. The KzO-fireclay vertical section of the K z O - A I 2 0 3 - S i O 2 ternary system is illustrated in Fig. 20. For those

FIG. 14. Furnace setup to react multiple refractory samples with alkali at elevated temperatures.

FIG. 15. Photograph of semi-silica refractory before and after reaction with soda vapor at 1350~ h.

FIG. 16. Photograph of fireclay refractory before and after reaction with soda vapor at 1350~ h.

FIG. 17. Photograph of mullite refractory before and after reaction with soda vapor at 1350~ h.

FIG. 18. Photograph of 90% alumina refractory before and after reaction with soda vapor at 1350~ h.

FxG. 19. Photograph of corundum refractory before and after reaction with soda vapor at 1350~ h.


KAS z +L .. KASz+Corundum +L KAS 4 +Corundum+L KAS4 + KASz + L

Corundum +KAS z +KAS 4 Mullite+KAS4+Corundum - -

Mullite +KASs+KAS4 - -

I

K20 30

fCorundum+L

L I Q U I D / ~ C o r u n d u m + KAS4 + L Mullite + L

" A ---J 16oo_... Mu l l i f e+c+L l / ~ M u i l i t e + K A S , + L

_'~1 I/I,.I-~s + T + L - ~ M u i l i t e + K A S s+L

I I I -I 800

Ii ~ M u l l i t e + O + K A S s

20 10 42Ale03 58 Si02

FIG. 20. The K20-fireclay vertical section of the Na20-AI203-SiO 2 system.

interested, it is instructive to compare this vertical section with its NazO-fireclay counterpart. To a large degree, the corrosive reactions are similar.)

Vertical sections of ternary oxide phase diagrams are relatively easy to construct by one who has training in oxide phase equilibria. In those instances where high-temperature reactions are relatively fast, as in the case of alkali oxides reacting with common refractories, the vertical section of the ternary diagram containing the corrodant and the refractory composition of interest is an extremely valuable guide in determining the high-temperature reactions that will occur. This information, together with a moderate knowledge of oxide liquids (glass) and selected properties of the crystalline reaction products, enables the researcher to evaluate rather accurately the corrosive effect of the oxide of interest on the refractory. Most importantly, this information reduces the need for extensive experimentation to evaluate the high-temperature corrosion of many ceramic materials.

Acknowledgments

The author wishes to acknowledge Ms. Rhonda F. Hayden who conducted a large portion of the literature survey presented here as part of her M.S. thesis program at Virginia Tech.

References

1. S. Aramaki and R. Roy, The alumina-silica phase diagram. J. Am. Ceram. Soc. 42(12), 644 (1959); 45(5), 239 (1962).

2. E. M. Levin, C. R. Robbins, and H. F. McMurdie, eds., "Phase Diagrams for Ceramists," pp. 123, 157, 181. Am. Ceram. Soc., Columbus, OH, 1964.


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A. M. Alper, ed., "Phase Diagrams: Materials Science and Technology," Vol. 2. Aca- demic Press, New York, 1970. J. F. Schairer and N. L. Bowen, The system Na20-AIzO3-SiO 2. Am. J. Sci. 254(3), 129-195 (1956). L. DePablon-Galan and W. R. Foster, J. Am. Ceram. Soc. 42, 491-498 (1959). E. F. Osborn and A. Muan, "Phase Equilibrium Diagrams of Oxide Systems," Plate 4. Am. Ceram. Soc., Columbus, OH, 1960. F. C. Kracek, The system sodium oxide-silica. J. Phys. Chem. 34, 1583-1598 (1930). F. C. Kracek, Phase equilibrium relations in the system Na2SiO3-Li2SiO3-SiO 2. Am. Chem. Soc. J. 61, 2863-2877 (1939). G. R. Rigsby and H. M. Richardson, Occurrence of artificial kalsilite and allied potassium aluminum silicates in blast-furnace linings. Am. J. Sci. 253(12), 734 (1955). N. L. Bowen, The binary system Na2Al2Si20. (Nephelite, Carnegieite)-CaAl2Si20 s (Anorthite). Am. J. Sci. 33(4), 551-573 (1912). N. L. Bowen and J. W. Grieg, The crystalline modification of NaA1SiO4. Am. J. Sci. 10(5), 204-212 (1925). M. H. Matignon, The system Na20-A120 3. Am. J. Sci. 254(3), 135 (1956). L. T. Brownmiller and R. H. Brogue, The system CaO-Na20-AI20 ~. Am. J. Sci. 23(5), 505-524 (1932). J. W. Grieg, Immiscibility in silicate melts. Am. J. Sci. 13(1), 133 (1927). J. F. Schairer and N. L. Bowen, The system K20-AI203-SiO 2. Am. J. Sci. 253(12), 681-746 (1955). J. R. Goldsmith, Some aspects of the system NaAISiO4-CaO �9 AI20 3. Am. Mineral. 34, 471-493 (1949). F. C. Kracek, N. L. Bowen, and G. W. Morey, Equilibrium relations and factors influencing their determination in the system KzSiO3-SiO z. J. Phys. Chem. 41(9), 1183-1193 (1937). L. T. Brownmiller, A study of the system lime-potash-alumina. Am. J. Sci. 29(5), 260-277 (1935). F. N. Strokov et al., Reaction of alkali carbonates with alumina and with silica, on fusion. Sb. Rab. G. Inst. Prikl. Khim. 32, 4-15 (1940); Chem. Abstr. 37, 28919. E. F. Osborn and A. Muan, "Phase Equilibrium Diagrams of Oxide Systems," Plate 5. Am. Ceram. Soc., Columbus, OH, 1960. M. J. Buerger, The silica framework crystals and stability fields. Z. Kristallogr., Kristall- geom., Kristallphys., Kristallchem. A90, 186-192 (1935). N. L. Bowen, The sodium-potassium nephelites. Am. J. Sci. 43(4), 115-132 (1917). F. A. Bannister and M. H. Hey, Kalsilite a polymorph of KAISiO4, from Uganda. Mineral. Mag. 26, 218-224 (1942). O. F. Tuttle and J. V. Smith, An annual report of the Director of the Geophysical Laboratory (1952-1953). Year Book--Carnegie Inst. Washington 52, 39-96 (1953). North American Refractories Company, "Refractories for the Glass Industry: Properties, Theoretical Considerations, and Recommendations." NARCO, Cleveland, OH. Z. Weyberg, On the alumino-silicate, K 2 0 . A I 2 0 3 . S i O 6. Neues Jahrb., Centralbl., pp. 326-330 (1908). R. E. Farris and J. E. Allen, Aluminous refractories-alkali reactions. Iron Steel Eng., February, pp. 67-74 (1973). A. Yamaguchi, Reactions between alkaline vapors and refractories for glass tank furnace. Int. Congr. Glass [Pap.], lOth, 1974, Vol. 2 (1974).

2 P R E D I C T I O N O F A L K A L I O X I D E C O R R O S I O N 83

29. G. R. Rigby and Hutton, Action of alkali and alkali-vanodium oxide slags on alumina-silica refractories. J. Am. Ceram. Soc. 45(2), 68-73 (1962).

30. W. A. Lambertson, Reactions of sodium sulfate with alumina-silica refractories. J. Am. Ceram. Soc. 35(7), 161-165 (1952).

31. J. W. Grieg and T. F. W. Barth, System N a 2 0 . AI20 3 �9 2SiOzNa20. A1203.6SIO 2. J. Am. Ceram. Soc. 30(11), 51 (1947).

32. H. Insley, Notes on the behavior of refractories in glass melting furnaces, J. Am. Ceram. Soc. 1(8), 583-593 (1924).

33. P. S. Dear, Study of soda-ash foundry-slag attack on the refractory lining of a teapot-type, reservoir ladle. Bull. Am. Ceram. Soc. 17(3), 4-8 (1938).

34. J. Wilson, J. Soc. Glass Technol. 2, 197-213 (1918). 35. F. S. Thompson and H. M. Kraner, Refractories for the manufacturer of glass. Ind. Eng.

Chem. 25, 856 (1933). 36. D. Turner and W. E. S. Turner, Some observations on the corrosion of fireclay materials

by alkali salts. J. Soc. Glass. Technol. 207-227 (1923). 37. J. W. Cobb, J. Soc. Chem. Ind. 29, 311 (1910). 38. E. C. Petrie and D. P. Brown, Observations on the shelling of checker brick. J. Am.

Ceram. Soc. 31(1), 15 (1948). 39. H. Insley, Some observations of surface deposits formed on glass-furnace regenerators.

J. Am. Ceram. Soc. 9(10), 635-638 (1926). 40. C. L. Thompson and E. P. Rexford, A study of an alumina-silica checker brick from the

regenerator of a glass tank. J. Am. Ceram. Soc. 21, 55 (1938). 41. A. E. Currier, Refractories for use in glass furnace regenerators. Bull. Am. Ceram. Soc.

29(3), 90-95 (1950). 42. L. H. Van Vlack, Chemical and mineralogical changes in stack and hearth refractories of

a blast furnace. J. Am. Ceram. Soc. 31(8), 220-235 (1948). 43. R. B. Snow, Effect of chemical attack and operational parameters on the wear of blast

furnace refractories. ERDA Contract Ret'. Meet. Chem. Energy Storage Hydrogen Energy Syst., 1976, No. E (49-18)-3731 (1976).

44. L. Halm, Note on refractory wear problems in blast furnaces due to the presence of alkaline compounds. Br. Iron Steel Inst. 8003; C.D.S. Circ. 4, 955-964 (1969).

45. S. E. McCune, T. P. Greaney, W. C. Allen, and R. B. Snow, Reaction between K20 and A1203-SiO 2 refractories as related to blast-furnace linings. J. Am. Ceram. Soc. 40(6), 187-195 (1957).

46. R. R. Lister, Investigation into failure mechanisms in alumino-silicate blast furnace boshes. J. Iron Steel Inst., London, December (1968).

47. J. F. Schairer and N. L. Bowen, Melting relations in systems NazO-AI203-SiO 2 and KzO-AI203-S iO 2. Am. J. Sci. 245(4), 193-203 (1947).

48. A. Yamaguchi and S. Okawara, Corrosion by vapors of potassium salts. J. Ceram. Soc. Jpn. 77, 357-366 (1969).

49. E. A. Thomas, A study of soda and potash vapor attack on superstructure refractories. J. Can. Ceram. Soc. 44, 37-41 (1975).

50. W. H. Boyer, "Subject: Alkali Test," Lab. Serv. Rep. Serial No. P75-261. 51. R. C. DeVries and W. L. Roth, Critical evaluation of the literature data on /3-alumina

and related phases. I. Phase equilibria and characterization of beta alumina phases. J. Am. Ceram. Soc. 52(7), 367 (1969).

52. R. F. Hayden, Study of the resistance of refractories to alkali vapors. M.S. Thesis, Virginia Polytechnic Institute and State University, Blacksburg (1980).

Appfication of Phase Diagrams to the Production

of Advanced Composites

WILLIAM B. JOHNSON AND ALAN S. NAGELBERG

Lanxide Corporation Newark, Deleware 19714

I. I n t r o d u c t i o n . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 85

II. R e v i e w of the Directed Metal Oxidation Process . . . . . . . . . . . . . . . . . . . 88

A. A l u m i n a M a t r i x C o m p o s i t e s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 88

B. Zirconium Diboride-Reinforced Zirconium Carbide Composites . . . . . . . . 93

I I I . Thermochemical Considerations of Matrix Formation . . . . . . . . . . . . . . . . 95

A. M a t r i x G r o w t h in the Absence of Fillers: A l u m i n u m O x i d e S y s t e m . . . . . . . 95

B. I n t e r a c t i o n s wi th F i l l e r s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 104

C. M a t r i x F o r m a t i o n in the Silicon Nitride System . . . . . . . . . . . . . . . . . . 107

IV. P h a s e E q u i l i b r i a in Carbide/Boride Systems . . . . . . . . . . . . . . . . . . . . . 112

A. P h a s e E q u i l i b r i a in t he Z r B 2 / Z r C / Z r S y s t e m . . . . . . . . . . . . . . . . . . . 112

B. T h e T i - B - C a n d H f - B - C S y s t e m s . . . . . . . . . . . . . . . . . . . . . . . . . 118

V. C o n c l u s i o n s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 121

References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 122

I. Introduction

Ceramic materials have properties that make them ideal candidates for many applications where hardness, wear resistance, stiffness, and elevated temperature stability are required. Due to the refractory character of ceramics, they often are the only choice for a material that can potentially satisfy the most demanding requirements. In addition to offering high melting or decomposition temperatures, many ceramics possess other attractive features such as low-density, high-temperature strength and resistance to creep, thermochemical stability leading to lack of reactivity with other materials and atmospheres, and high wear resistance.



86 WILLIAM B. JOHNSON AND ALAN S. NAGELBERG

Probably the single most important disadvantage of ceramic materials is their relatively low fracture toughness, i.e., their low resistance to the propagation of even a very small crack. This lack of toughness translates for design engineers into a low damage tolerance, or an unacceptably high potential for catastrophic brittle failure. In addition, traditional ceramic processing methods place limitations on the shapes and sizes of components that can be made conveniently.

One method of improving the fracture toughness of a ceramic is by incorporation of a second phase, forming a ceramic matrix composite (CMC) [1]. Various approaches have been demonstrated, including the introduction of ceramic whiskers [2], e.g., SiC, into a ceramic matrix, incorporation of a second ceramic phase that undergoes a phase transition on application of a stress field (such as that near a crack tip) [3, 4]; the incorporation of strong fibers that bond weakly with the ceramic matrix [5] and, finally, the introduction of a ductile metal phase to impart toughness by crack tip blunting [1]. Often these materials cannot be produced using traditional ceramic processing routes because of potential reactions/interactions between the various phases in the system. For example, the incorporation of dimensionally stable fibers or whiskers can prevent the sintering of the matrix [6], or the high temperatures required for sintering can cause deleterious reactions between the fiber and the matrix.

In some applications the lack of toughness of ceramics or CMCs prohibits their use. In cases where enhanced stiffness, wear resistance, or elevated temperature capabilities greater than those provided by metals are necessary, metal matrix composites (MMCs) offer a reasonable com- promise between ceramics or CMCs and metals. Typically, MMCs have discrete ceramic particulate or fiber reinforcement contained within a metal matrix. In comparison to CMCs, MMCs tend to be more workable and more easily formed, less brittle, and more flaw tolerant. These gains come primarily at the expense of a loss of high-temperature mechanical properties and chemical stability offered by CMCs. These materials thus offer an intermediate set of properties between metals and ceramics, though somewhat closer to metals than ceramics or CMCs. Nonetheless, like ceramic matrix composites, they involve physical mixtures of different materials that are exposed to elevated temperature processes, and therefore evoke similar thermodyamic considerations for reinforcement stability.

Within the last decade several novel processing methods have been developed to overcome the limitations of traditional methods. For CMCs, these include chemical vapor infiltration [7-10] repeated infiltration/ pyrolysis of sols or polymers [11], crystallization of glass matrices [12-14] and the direct formation of a ceramic matrix by the directed oxidation of a

3 PRODUCTION OF ADVANCED COMPOSITES 87

molten metal [15-21]. Metals and ceramics have been combined to form MMCs by blending metal and ceramic powders before consolidation [22], infiltration of metal into a porous ceramic body [23], pressure casting [24], compocasting [25], exothermic dispersion (XD TM) processing [26], and pressureless metal infiltration [17, 27, 28]. The directed metal oxidation technology to form CMCs is particularly intriguing because more than one toughening mechanism can be used at the same time; for example, metal phase toughening and toughening through the incorporation of a second phase. Further, it can do so in a net-shape or near-net-shape process.

This chapter, then, deals primarily with the directed metal oxidation process, although selected examples of stability in metal matrix composites are also discussed briefly. The focus is, of course, on the applications of phase equilibria, and more generally, thermodynamic principles that are applicable to the formation of composites in the presence of molten metals. Because these general principles are the same regardless of whether the end product is an MMC or a CMC, little generality is lost by focusing the discussion on CMCs formed by directed metal oxidation.

The directed metal oxidation process is quite intriguing from the viewpoint of thermochemistry. It involves a multiphase system, which normally includes a gas phase, a liquid metal phase, and one or more solid ceramic phases. The molton metal phase is typically a multi-component alloy, and the ceramic solid often contains two or more distinct phases. This complexity, then, requires a knowledge of both multi-component metallic and ceramic phase diagrams, as well as an understanding of the interactions between the phases. Further, because the processing normally involves the presence of a gas phase, detailed knowledge of potential reactions between the molten and solid phases with the gas must also be considered. As might be imagined, traditional binary and /o r ternary phase diagrams often provide only a starting point for thermochemical understanding. Additional equilibrium diagrams must often be constructed to determine applicable processing conditions, potential reactions, appropriate alloys, etc. Often, the complexity precludes quantitative analysis. In these cases, qualitative trends can still be determined and, in fact, can be quite useful.

Because the details of processing in each class of CMCs (e.g., oxide, carbide, or nitride matrix) are slightly different, the appropriate thermochemical approach for each class may also be different. For example, in the formation of alumina matrix materials by directed metal oxidation, the alumina product grows from a molten aluminum alloy by reaction with an oxygen-containing gas phase. On the other hand, in the formation of platlet-reinforced zirconium carbide, the gas phase is not involved in the reaction at all, being inert to the reactants and products. Thus, a general approach to deal with the myriad of possible products formed by the


directed metal oxidation process is not easily described. The approach taken here, then, will be to outline a series of examples involving directed metal oxidation where the thermodynamics has been elucidated. The understanding derived from the description is not of academic interest, but rather can be (and often has been) used to effect important processing changes, which in turn produce better composites.

Before describing in detail these specific examples, the next section provides a brief overview of the directed metal oxidation process. The information provided will allow a better understanding of the phase diagram and thermochemical applications that follow.

II. Review of the Directed Metal Oxidation Process

In this section the preparation of ceramic composites by the directed metal oxidation process is described. First, in Section II.A, the aluminum oxide system is used as an example to explain the nature of the process, and a further example, a ZrB 2 reinforced ZrC composite, is discussed in Section II.B.

A. Alumina Matrix Composites

The directed metal oxidation process is illustrated schematically in Fig. 1. Using the alumina system as a specific example, a preform of the desired reinforcement is formed and placed in contact with an aluminum alloy. The entire assembly is coated with a growth barrier, which stops the oxidation process when the growth front contacts it. With the appropriate choice of alloy and temperature [15] the aluminum will be oxidized to form a dense interconnected A120 3 ceramic matrix containing a small amount of residual metal. Further, this growth will proceed through the preform without displacing it, forming a composite with a ceramic matrix, which contains the reinforcement phase of the preform. For example, with an Al-10% Si-3% Mg alloy and a SiC preform placed at 1100~ in air, the composite contains an interconnected A120 3 matrix with ~-10 vol % residual aluminum [16] (Fig. 2).

The growth rates are quite high, growing through fillers as fast as 2.5 cm/day with the length and width of samples limited only by the size of the preform and the pool of molten aluminum. The rapid growth rate permits thick composites to be prepared, with thicknesses of up to 20 cm reported [ 15].

3 P R O D U C T I O N O F A D V A N C E D C O M P O S I T E S 89

FIG. 1. A schematic illustration of CMC growth to net-shape using a directed metal oxidation process where the preform is formed by cold-pressing.

FIG. 2. An optical micrograph of an AI203/AI matrix grown through a SiC reinforcement phase.

90 W I L L I A M B. J O H N S O N A N D ALAN S. N A G E L B E R G

The mechanism of the growth process has also been studied [29, 30]. It is proposed that a metastable oxide layer (either ZnO or MgO depending on the alloy) on the surface of the growing reaction product allows rapid oxygen transport to the metal phase and inhibits the formation of a protective alumina layer on the molten metal. Once through this layer, the oxygen dissolves in a thin layer of molten metal and subsequently precipitates as A1203 on the growing A1203 structure. The microstructure [31] and effects of processing parameters on the growth rate [32] have also been discussed.

Further details of each of the steps in the formation of a ceramic composite article are given in the following sections.

1. PREFORM FABRICATION

The first step in the manufacture of a CMC component is the production of a net-shape or near-net-shape preform from either ceramic particles or ceramic fibers. Particulate preforms can be produced by any of the conventional ceramic processing routes, such as uniaxial or isostatic pressing, injection molding, sediment casting, extrusion, or slip casting. A

FIG. 3. A schematic illustration showing the various steps employed to form a tubular component by the directed metal oxidation process. The preform is formed by slip casting.


schematic of the formation of a preform using slip casting is shown in Fig. 3. For fiber reinforcements, preforms can be produced by techniques such as fabric layup, filament winding, weaving, or braiding.

For preforms of fiber reinforcements, a thin coating is applied to the fibers using chemical vapor infiltration (CVI). This coating step is essential both to protect the fiber from chemical attack by the strongly reducing aluminum alloy and to provide for a weak fiber/matrix interface in the composite. Because the coating is thin, the CVI step requires only a few hours, unlike CVI matrix formation processes, where long times are necessary to achieve sufficient densification.

The final preparation step is the application of a growth barrier, which allows the fabrication of components to net-shape or near-net-shape. This is accomplished by coating all preform surfaces, except the one in contact with the alloy, with a gas-permeable barrier, e.g., CaSO 4. This growth barrier inhibits the growth process locally when the growth front has traversed the preform and contacts the barrier. The barrier functions by forming a reaction product with the metal that strongly reduces further growth.

2. CERAMIC MATRIX GROWTH

The preform is now ready for matrix growth. It is placed in contact with alloy, either with solid alloy before heating, or with molten alloy after preheating, as shown in Fig. 3. The layup is further heated to a temperature typically between 900 and 1200~ Following a short initiation period, a highly interconnected ceramic matrix phase of alpha alumina begins to grow into the preform starting from the alloy/preform interface. Present within the reaction product matrix is an interconnected network of microscopic metal channels that are at most a few micrometers in diameter (Fig. 2). Driven by surface energy forces, the alloy wicks from the reservoir through the microchannels to the growth front where it reacts to continue the matrix growth process.

Matrix growth will continue as long as alloy chemistry, temperature, and oxygen availability are sufficient to sustain growth. This allows components of thick cross section to be produced. Growth rates depend on the process conditions and the desired microstructural characteristics of the matrix. Depending on the size and shape of the component, several hours to several days may be required to complete the growth. A more detailed qualitative description of the effect of processing conditions on growth rates and microstructures is given in a subsequent section.

For many applications, the interconnected residual metal in the composite provides a desirable increase in the low-temperature toughness of the


material. For other purposes, the metal may be undesirable. In these cases, additional processing steps are utilized to remove the metal, change its composition, or seal it off from the environment. For example, where access of the molten metal to the surface can lead to detrimental continued oxidation or reaction with adjoining structures, a high-temperature surface sealing process can be used to form an adherent ceramic skin on the composite. This surface layer enhances the composite's environmental stability over a wide range of conditions at elevated temperature.

3. CERAMIC MATRIX MICROSTRUCTURE

The a-alumina structure, in the absence of a reinforcement, consists of long continuous columnar regions containing predominantly low-angle AlzO3/A1203 grain boundaries [31]. Within the columnar regions, the alumina exhibits a preferred orientation, the c-axis being parallel to the growth direction. The AlzO3/A120 3 boundaries are devoid of any grain- boundary phase. The metallic component of the composite is present as both interconnected tortuous microchannels and as isolated small pockets.

The incorporation of a particulate reinforcement into the ceramic matrix produces a refinement of the matrix, i.e., both the ceramic ligament and metal channel sizes are decreased [31]. Although the reinforcement particles effectively break up any macroscopically apparent columnar growth of the matrix, some preferred orientation of the A120 3 is still observed by x-ray diffraction analysis. In the specific case examined, approximately one-half of the surface area of the reinforcement particles was found to be directly bonded to the interconnected A1203 of the matrix; the balance was in contact with the metallic constituent.

The microstructures of ceramic matrices grown from two different classes of alloys have been reported. The external growth surface of ceramic matrices grown from an A1-Si-Mg alloy in the absence of a reinforcement was covered by a thin (~ 1- to 4-1xm) layer of MgO that sometimes contained up to 5% MgAI204 [33]. The external MgO layer typically was separated from the interconnected A120 3 matr ixby a thin aluminum alloy (1- to 3-1xm) layer. Only rarely was an A1203 grain found in direct contact with the external oxide layer. Within the bulk of the composite, the metallic channels typically were 3 to 8 txm in width.

When a Zn-containing alloy was used to form the composite matrix, the external surface of the growth was covered by a thin layer of ZnO [29]. No ZnAI204 was observed. For these composites, a thin metal layer separating the external ZnO layer from the A120 3 matrix was also observed, but it was significantly thinner than the layer found in using A1-Si-Mg alloys.


These composites typically contained a refined microstructure with 1- to 3-/~m-wide channels.

B. Zirconium Diboride-Reinforced Zirconium Carbide Composites

More recently, ceramic composite materials have been described that incorporate zirconium diboride platelet reinforcements in a zirconium carbide matrix [34, 35]. These materials are prepared by the directed reaction of molten zirconium with boron carbide (B4C) to form a ceramic material composed of zirconium diboride platelets in a zirconium carbide matrix with a controlled amount of residual zirconium metal.

In this case, a liquid metal, molten zirconium, can be reacted (oxidized) by a bed of solid oxidant (i.e., B4C) to form products that are different than the bed (Fig 4). Just as described earlier, there is a directed oxidation reaction of a liquid metal and the product may contain some residual metal. In this case, the boron carbide bed is consumed according to the reaction:

(2x + y + 1)Zr + xB4C ---) 2xZrB 2 + ZrC x + yZr (1)

where y is the number of excess moles of zirconium added to the reactants and x is the stoichiometry of the carbide phase.

One advantage of this processing approach is that the amount of retained metal phase in the composite can be controlled from essentially zero up to greater than 30 vol % simply by controlling the stoichiometry of the reaction [34]. For example, for reaction (1)with y equal to zero, there will theoretically be no free Zr in the body, whereas when y equals 0.5, there will be ~ 18 vol % Zr in the final composite.

The ability to control the amount of residual metal is one key advantage to the process, because the properties of the material are dependent on

FIG. 4. A schematic illustration of the processing used to produce zirconium diboride reinforced zirconium carbide materials by directed metal oxidation.

FIG. 5. The microstructure of zirconium diboride-reinforced zirconium carbide composites produced by directed metal oxidation. (a) A composite prepared with ~ 22 vol % metal. (b) A composite produced with less than 2 vol % metal.


the metal content [35]. For example, the fracture toughness increases linearly from ~ 11 MPa-m 1/2 at 0 to 2 vol % Zr, to 22 MPa-m 1/2 for 30 vol % metal. The high toughness even at low metal content arises because of the in situ formation of ZrB 2, which forms as platelets (Fig. 5), giving rise to crack deflection and crack bridging during deformation. Additional toughening is provided by the metal phase, with the extent of metal phase toughening increasing as the metal content increases.

The kinetics of formation of this zirconium diboride platelet reinforced zirconium carbide have been discussed, as have possible formation mechanisms [36] and detailed microstructural and orientation relationships between the phases [37]. These materials, in addition to being very refractory, are quite hard. Potential applications typically involve wear resistance, either at low to moderate temperatures or for short times at very high temperatures, such as in biomedical and rocket nozzle or rocket motor application [ 38].

III. Thermochemical Considerations of Matrix Formation

Typically, the directed metal oxidation process involves the simultaneous reaction of molten metal, e.g., A1 with O 2, and infiltration of the reaction product and metal into a porous preform of the desired reinforcement. The directed metal oxidation process can also form composites in the absence of a reinforcement phase, termed matrix-only growth. Al- though the former process is more interesting commercially because of the ability to tailor the composite properties and because the product does not show significant preferred orientation, the latter case is simpler conceptu- ally and theoretically. Thus, the thermodynamic discussion will begin with growth in the absence of reinforcements and then cover the additional complications that arise from their presence.

A. Matrix Growth in the Absence of Fillers: Aluminum Oxide System

As described, the addition of Mg to an aluminum alloy permits rapid oxidation of the alloy at elevated temperatures in an oxygen-containing atmosphere. The product is a composite composed of a matrix of interconnected A |20 3 with a small fraction of residual A1 alloy dispersed in the matrix as interconnected metal channels. From a thermodynamic or phase diagram viewpoint, three distinct oxides can form when an A1-Mg alloy is exposed to such an atmosphere: A1203, MgAI204, or MgO. The specific


1/2 0 2 (g)

-~ ////t /// I

AI203+ (AI, Mg) / / / MgO; (AI,Mg)

l / I l l S / / / II AI

/ ~-7.97 at.%Mg (7.24 wt.%) (Liquid) I l.V.," ./ / ,

L_ 0.446 at.%Mg (0.40 wt.%)

MgAI204 + (AI, Mg)+ MgOT~Z-~ / /

ALQ 3 + (AI, Mg)+ MgAI2q--~ z z .,,-'Z__--34.23

Mg (Liquid)

FIG. 6. Ternary A1-Mg-O phase diagram at 1373 K showing the alloy composition ranges over which A1203, MgAI204, and MgO are stable. From Salas et al. [18].

oxide that does form depends on the alloy composition and temperature at the reaction front. For example, the phase diagram (Fig. 6) predicts that at 1373 K, molten A1-Mg alloys with Mg concentrations above ~ 7 wt % should form MgO in an oxygen atmosphere, whereas magnesium aluminate should form for Mg concentrations between 0.4 and 7 wt %. A120 3 is the stable oxide only for Mg concentrations lower than 0.4 wt %. The predominance of the MgO-(A1-Mg) two-phase region results from three factors: (1) the higher stability of MgO compared to A1203 on a per mole oxygen basis, (2) the additional stability of the ternary oxide MgA|204 over MgO and A120 3, and (3) the negative deviation from ideality for A1-Mg liquid alloy solutions.

Although thermodynamics predicts that only alloys with less than ~ 0.4 wt % Mg should form A1203 matrices at 1373 K, experimental results show that alloys with Mg contents well above 10 wt % form A120 3 matrix composites. This occurs because of the formation of a layer composed predominantly of MgAI204 at the original alloy/atmosphere interface, which depletes the aluminum alloy of Mg, at least locally. The reduced Mg activity in the alloy permits A1203 to form at the MgAlzOa/atmosphere interface as fresh alloy wicks through the MgA|204 layer. As growth continues, both the MgAlzO 4 and A1203 layers increase


in thickness as long as the alloy contains Mg levels greater than those in equilibrium with A1203. The MgAl204 layer thickens because the growth alloy wicking through the reaction product reacts with A1203 to form MgAI204 according to

3 Mg + 4 A1203 ---) 3 MgA1204 + A1 (2)

reducing the Mg content to that at the alloy/MgA1204/A1203 three-phase equilibrium (0.4 wt % Mg for an A1-Mg binary alloy at 1373 K). At the outer A1203/atmosphere surface, the A1203 layer grows by reaction with the atmosphere, as described later. The A1203 layer grows much faster than the MgA1204 layer because of the preponderance of A1 in the growth alloy. The final thickness of the MgA1204 layer depends directly on the amount of Mg in the growth alloy and is normally quite thin, on the order of a few to tens of microns thick.

During growth, the external surface of the A1203/A1 composite is covered by an MgO layer. It is the persistence of this layer at the outer growth interface that permits the continuous rapid formation of the A1203 ceramic matrix [30]. (Clearly, if A1203 formed as a continuous layer at the surface, its low permeability for A1 and/or O would not support the observed rapid reaction rates.) The presence of a metastable external MgO layer in direct contact with a molten Mg-depleted A1 alloy supports the A1203 matrix formation process by generating a significant oxygen solubility gradient, A[O]/Ax, across the thin metal layer of thickness, Ax, that separates it from the growing A1203 ceramic matrix. Driven by this solubility gradient, oxygen rapidly diffuses across the metal layer, reacts with A1, and precipitates A1203 on the growing "columns" of ceramic [30] (Fig. 7). The rate of oxygen diffusion, and thus the ceramic formation rate, will be influenced by the size of this gradient. The MgO layer allows rapid oxygen transport by Mg vacancy formation during a surface reaction with oxygen, followed by oxygen dissolution and vacancy destruction at the MgO/alloy interface. Grain-boundary diffusion of Mg compensated by electronic conduction maintains the layer and charge neutrality. Vlach et al. [39], building on these ideas, illustrated the concept using a phase diagram that plots oxygen solubility in A1 versus alloy composition for the alloy equilibrium with MgO and A1203 (Fig. 8).

Frequently, Si is also added to the aluminum alloys used for composite growth. The silicon appears to play a role in the growth nucleation process both during the initial stages of the growth process and during renucle- ation of composite growth in the event that the surface MgO/a l loy / ceramic structure breaks down [40].


~ j ~0 + Mg "~ MgO + Yl~g+ 2h* --"- surface reaction 2 2 Mg

grain boundary diffusion of Mg Mg " compensated by electronic conduction

. .

MgO + Vl~g.--* Mg I~g + O + 2e' - - - - oxygen dissolution

Al -a l l oy

alumina

II AI

0 diffusion of dissolved oxygen

2 AI + 3 0 - - - AI20~. A120 z precipitation

aluminum alloy convection by wicking through reaction product

FIG. 7. Schematic of the composite growth surface showing the sequence of layers involved in the reaction process. From Nagelberg [29].

T {oi

AI

\ Effect of Composite hickness

[ ! N, MgOSolubility

Liquid Alloy

Mg

AI + Mg

FIG. 8. Schematic plot of oxygen solubility vs. alloy composition illustrating the source of the oxygen gradient that drives the composite growth rate. From Vlach et al. [41].


In addition to this nucleation effect, though, there is a thermochemical effect of silicon. Thermochemical calculations by Vlach et. al predict that the addition of Si to the alloy can have a pronounced effect on the calculated oxygen solubility in the alloy in equilibrium with MgO, MgAI204, or A120 3 [33, 41]. Thus, the oxygen chemical potential gradient across the thin metal layer between the growing alumina and the MgO (or MgAI204) surface skin is modified. With a 10 wt % Si addition to the alloy and a fixed Mg concentration, the oxygen solubility in equilibrium with A120 3 is changed only slightly by the addition of silicon as shown by the two nearly horizontal lines near the bottom of Fig. 9. Conversely, the

Effect of Alloy [Si] on Solubilities and Rate

I r'' [O] ~ \ \-*--- in (AI, Mg, 10 at. % Si) / \- '~-k~-in (AI, Mg)

Liquid Alloy

\ \ \

ot-AI203 Solubility

in (AI, Mg, 10 at. % Si)

AI, Mg)

AI Mg AI + Mg

FIG. 9. Schematic plot of oxygen solubility vs. alloy composition showing the effect of Si on the solubilities of MgO and AI20 3 and hence the oxygen gradient across the near surface alloy layer. Vlach et al. [41].


calculated equilibrium oxygen concentration at the MgO/alloy interface (or MgAlzO4/A1 , not shown) is significantly increased by the addition of silicon. Assuming the metal layer thickness does not change, the oxygen gradient across the metal layer will thus increase relative to an alloy without silicon.

To more clearly understand this effect, consider reaction (3), which has an equilibrium constant given by Eq. (4):

MgO = Mg + 0 (3)

K 3 = aMga O = XMg'YMga O (4)

where the underlined species indicate that they are dissolved in the melt at less than unit activity. As indicated by Eq. (4), the product of the Mg and O activities will be a constant at any temperature as defined by the free energy of formation for aluminum saturated MgO. For a given Mg concentration, the addition of an element that affects the activity coefficient of Mg, '~Mg, will then lead to a change in the oxygen solubility. For example, adding Si to the alloy will lower Y Mg because Mg and Si have a positive interaction coefficient in the melt, as indicated by the high melting point of MgzSi in the Mg-Si system. Correspondingly, then, at a fixed Mg concentration one would expect that the equilibrium oxygen concentration at the MgO/alloy (or MgAlzO4/alloy) interface will increase as the Si concentration is increased. Conversely, the interaction in the molten A1 alloy between A1 and Si is relatively small, as indicated by the nearly ideal liquid metal thermodynamic behavior, and thus Si only has a minor effect on the equilibrium oxygen solubility at the AlzO3/alloy interface. There- fore, the addition of Si should enhance the oxygen concentration gradient across the metal layer that separates the MgO from the A1203 matrix, as shown in Fig. 9, and thereby increase the growth rate.

There is an additional thermochemical effect of Si additions to the alloy not considered by Vlach et al. The addition of Si to an A1-Mg growth alloy and the accompanying modification of the Mg activity coefficient will affect the alloy composition for the three phase equilibrium between MgAI204, A1203, and Al-alloy. At the MgAlzO4/AI203 interface in the composite the alloy is equilibrated with MgAI204 and A1203 according to reaction (2). This equilibrium interrelates the activity of Mg and A1 according to

K 2 aAl2/aMg 3 (5 )

As the alloy continues to wick past this interface to the external surface,


the alloy composition is fixed by this reaction. At the MgO/alloy growth interface, the Mg activity is described by reaction (3) and the oxygen activity given by Eq. (4). Across the external metal layer at the growing A 1 2 0 3 columns, the oxygen activity is given by reaction (6) and Eq. (7):

AI2O 3 = 2 A1 + 30 (6)

2 3 K 6 -- aA1 a o (7)

The gradient in oxygen activity is the difference between the oxygen activity given by reactions (2) and (6):

Aa o = ao[reaction (2)]-ao[reaction (6)] (8)

divided by the metal layer thickness. By substituting for the a Mg in Eq. 4 from Eq. 5, then solving for a o in the resulting expression and in Eq. 7, the difference in oxygen activity can be expressed only as a function of the

aAl"

= .-. - 2 / 3 _ . - 2 / 3 Aa o K21/3K3uAI UAl (9)

The effect of adding silicon on the oxygen gradient is then expressed only by its reduction of the aluminum activity, primarily due to dilution since its interaction with silicon is small. A reduced aluminum activity results in a higher oxygen solubility at both the alloy/A1203 and alloy/MgO interfaces. Once again, assuming the metal layer thickness does not change, the addition of Si is found to enhance the oxygen gradient across the metal layer separating MgO from the growing A1203 matrix, though by a somewhat smaller amount than the previous analysis.

Experimental results find just the opposite of this prediction, i.e., as the Si content of the starting alloy is increased the growth rate decreases (42). There are several possible explanations for this discrepancy. First, the assumption that the metal layer thickness between the external MgO layer and the A1203 does not change with Si alloy additions might not be valid. The surface energy relationships that are a factor in determining this thickness can be affected by alloy composition. Second, the preceding analysis has not considered the mass transfer step through the external MgO layer. Here also, the presence of silicon in the alloy may affect the ionic and electronic transport properties of the MgO, thereby changing the kinetics.

More important than these factors, though, the preceding discussion assumes that the alloy composition near the growth front is determined


solely by the al loy/MgA1204/Al20 3 equilibrium. In reality, the growth process is dynamic. During growth, the A1 near the growth front is constantly being removed from the alloy to form A1203. The Si, Mg, and A1 concentrations near the growth front are thus continually changing. For example, the Si concentration near the growth surface increases with time because Si remains unoxidized due to the large difference in stability between SiO 2 and A120 3. Thus, Si tends to become enriched near the surface as A1 is preferentially oxidized. The Si, of course, tends to diffuse away from the growth surface through the metal channels to reduce this concentration gradient. The diffusion process, though, may be slower than the growth process, allowing a Si concentration gradient to be maintained. This buildup of Si at the growth front affects the kinetics of growth because as Si builds up at the surface it changes the alloy chemistry. These concentration changes at the growth interface may affect other thermodynamic system parameters, such as the activity coefficient of oxygen. The oxidation process may even become starved for A1 because of the buildup of Si, leading to mixed reaction/mass-transfer kinetics. The kinetics are clearly more complex than the simple thermodynamic models described earlier. More detailed quantitative modeling would require coupling time- dependent phenomena, e.g., diffusion and metal mass transport, with the thermodynamic description.

Thus, the simple thermodynamic models relating oxygen gradient to growth rate are, in the case of Si additions to the alloy, counter to the experimental results. Yet, the description still may give some guidance in assessing the effect of various changes in the processing conditions. For example, the same thermodynamic arguments described earlier can cor- rectly predict the effects of the addition of other alloying elements, as is described more fully later. Generally, the approach is to examine the effect the additions have on the oxygen gradient to assess changes in the kinetics. If all the interaction coefficients between the alloying elements were known, i.e., if the activity coefficients of each element could be calculated for any alloy addition, potential effects could be more accurately predicted given a valid kinetics model. Generally, this is not possible, both because the data are not available, and because of the dynamic nature of the growth process. Yet these qualitative predictions are useful in predicting first-order effects on the growth rate.

The extent of these interactions can be predicted using the same tools used to predict multi-component phase diagrams [43]. In the absence of detailed thermodynamic relationships for the system being considered, thermodynamic data for binary (or ternary) systems that are subelements can be used to supply useful trends. Examination of phase diagrams and


the stability of individual compounds or intermetallic phases can also be used to obtain useful information about elemental interactions.

For example, Ni additions to an A1 alloy used to grow an A1203 matrix composite would be expected to more strongly interact with A1 than Mg. The strong interaction between Ni and A1 is indicated by the high heat of formation for several Ni-A1 binary intermetallic phases and their high melting points (in several cases higher than the constituent Ni and A1). Conversely, Ni does not form highly stable intermetallics with Mg and therefore would not be expected to influence the Mg activity coefficient to the same extent that it affects the A1 activity coefficient. In addition, since Mg is only present at the composite growth front at a relatively low concentration as compared to A1, the effect of Ni on the Mg activity is further reduced. Ni additions to the growth alloy should then lower the oxygen concentration gradient, and thereby the growth rate. Such an effect has been observed experimentally [42].

The addition of alloying elements can also indirectly affect the oxygen solubility gradient across the thin metallic surface layer by modifying the activity of an element that interacts with A1. For example, Fe forms very stable intermetallics with silicon, while the Fe-A1 and Fe-Mg intermetallics are less stable. Additions of Fe to an A1-Si-Mg alloy, therefore, will decrease the Si activity, and thereby increase the Mg activity. This will lead to a lower oxygen potential gradient and a lower growth rate.

The addition of alloying elements to the growth alloy, e.g., transition metals, can have other ramifications on the growth process and the resulting composite microstructure and properties. Transition metal additions lead to a modification of both the ceramic matrix and the composition of the alloy retained within the matrix [44]. The resultant matrix has a higher ceramic content and finer metal channels, presumably because of the slower growth rate and /o r changes in the wetting characteristics of the AlzO3/alloy interface. Further, the transition metal additions can lead to the formation of alloy constituents and /o r intermetallic phases with higher melting points than the base A1 alloy. Depending on the nature of these phases, they can impart improved corrosion/erosion resistance to the metallic constituent of the composite.

The addition of alloying additions more noble than aluminum, i.e., those that form less stable oxide reaction products, results in an enrichment of these metals in the alloy near the growth front because A1 is constantly being removed by oxidation to form A120 3. If the addition is highly soluble, the local thermochemistry changes according to interactions with the other alloy constituents, as described earlier. On the other hand, if the addition has a limited solubility, the alloy composition at the growth front


can become saturated, and a solid (or solids) may precipitate out of solution. Should this occur, A1 access to the reaction front and the diffusion of more noble elements away from the reaction front may be inhibited, thus slowing or even stopping the growth process.

B. Interactions with Fillers

The incorporation of reinforcements into a ceramic or metal matrix composite introduces additional stability constraints. The reinforcement should be compatible not only with the matrix itself, but also with the processing atmosphere and temperature. For example, the use of fine particles of TiB 2 as a reinforcement in an A1203 matrix formed by a directed metal oxidation process is not compatible with the oxygen- containing growth atmosphere because the TiB 2 rapidly oxidizes. Con- versely, the incorporation of the same TiB 2 particulate into an A1N matrix is feasible because the TiB 2 is stable in nitrogen growth atmospheres, i.e., the reaction

TiB 2 + 3 /2N 2 ~ TiN + 2BN AG ~ > 0 (10)

has a positive free energy of reaction at the growth temperature. Further- more, it is expected that TiB 2 would be stable in the presence of a molten A1 alloy at typical processing temperatures as determined by the reaction:

TiB 2 + 4 A1 ~ TiA13 + AIB 2 AG ~ > 0 (11)

Eq. (11) is not written as a simple exchange reaction, but rather includes the stable titanium aluminide and aluminum boride phases expected to be in equilibrium with molten A1. Consideration of the formation of intermetallics must not be overlooked during analysis of reinforcement stability because these, not the free element, are the most likely phases to form (presuming, of course, they exist at the growth temperature).

Alternately, the use of a preform constituent that reacts with the infiltrating ceramic matrix during processing can be advantageous to the properties of the final composite product. For example, as noted, the presence of transition metals in the alloy used to grow an A120 3 matrix composite results in a refinement in the matrix microstructure. A similar effect can be achieved by adding a reducible compound containing the desired transition metal alloying addition; for example, adding NiO or NiAlzO4 to a preform in order to dope the growth alloy with Ni [45]. The rapid reaction with the infiltrating aluminum alloy is driven by the stability


difference between NiO (or NiA1204) and A1203, as illustrated by the reactions

3 NiO + 2Al(1) ~ A1203 + 3Ni AG ~ << 0 (12)

or

3 NiA1204 4- 2Al(1) ~ 4 A1203 4- 3Ni AG ~ << 0 (13)

The large difference in stability between NiO and A1203 makes nearly all A1-Ni alloy compositions in equilibrium with A1203, as shown by the wide A1203 stability region in Fig. 10, even though numerous very stable nickel aluminides can form. Once again, a balance between the alloy Iiquidus surface composition, matrix formation rate, and the diffusion of transition metal away from the reaction front must be maintained. Typically the presence of a reducible compound in the preform, especially if it is the sole constituent, results in a further refinement of the composite microstructure.

A third, intermediate matrix/reinforcement stability case exists. In this instance, the reinforcement can be stabilized by appropriate selection of

, , , . .

z

. . J

I AI Ni 3

- 2 "

AI - 4 -

-6 -

- 8 -

- 1 0 -40

AI Ni

3Ni

AI(I)

- " 13 Ni AI203 NiAI204

I I I I -32 -24 -16 -8

Log ( Po2 )

0

Fro. 10. The AI-Ni-O stability diagram as a function of oxygen partial pressure and Ni activity in molten aluminum at 1273 K.


the composition of the alloy. The production of A1 MMCs containing a SiC reinforcement is a prime example of this case. When pure aluminum is used as the matrix for a SiC reinforcement, the following reaction occurs during the required elevated temperature processing,

4 A1 + 3SiC ~ 3 Si + A14C 3 (14)

Consideration of the A1-Si-C ternary diagram [46] would suggest that the following reaction also be considered:

4 A1 + 4 SiC ~ 3 Si + A14SiC 4 (15)

Both reaction products have been reported, but the formation of AI4C 3 seems to be kinetically favored at temperatures below ~ 1620 K [47, 48]. Unlike the former case where NiO and A120 3 had quite different stabilities, the stabilities of SiC and A14C 3 are quite close to each other. Therefore, it becomes possible to stabilize the SiC in the molten aluminum by adding Si to the alloy. For a binary A1-Si alloy, exact calculations can be made from the Si activity data in the melt. Using the Si activity data of Murray and McAlister [49], and the free energy of formation data from JANAF [50] the equilibrium Si concentration increases from 8.4 wt % at 880 K to 12.8 wt % at 1100 K [51]. Thus, at metal matrix processing temperatures of around 1000 K, the addition of ~ 10 wt % Si is required to stabilize SiC and prevent the formation of A14C 3. Of course, at higher temperatures, higher concentrations of Si are required. Further, for ternary or more complex alloys, where there may be interactions between the additional alloying elements and either Si a n d / o r A1, the amount of Si required may be higher or lower depending on the details of the interactions. Generally, quantitative evaluations are not possible because interaction coefficients are not available, but by using arguments similar to those presented here, it is possible to make qualitative predictions. For example, Mg additions to an A1-Si alloy would be expected to increase the amount of Si in the alloy required to prevent AlaC 3 formation because Mg has a positive interaction coefficient with Si, as evidenced by the stability of MgzSi. This interaction will be much stronger that the interaction with A1, so that the Si activity will be lowered more than the A1 activity, thus tending to drive reaction (14) toward the right. To suppress this tendency, additional Si must be added to the alloy to overcome the positive interaction with Mg.


C. Matrix Formation in the Silicon Nitride System

As has been reported recently [52] the directed metal oxidation process has been extended to the formation of Si3N 4 matrix composites. Using Si-Mn or Si-Fe growth alloys, temperatures between 1500 and 1700~ and a nitrogenous atmosphere, a matrix of Si3N 4 containing residual silicon has been produced. This matrix has been reinforced by growing through preforms of Si3N4, SiC, and carbon fibers coated with SiC. The latter two of these three cases provides a particularly interesting example of the type of thermochemical and phase diagram analysis that can lead to a more complete understanding of processing conditions. These two cases are the subject of this section, first with the growth of Si3N 4 into SiC reinforcements, and then into SiC-coated carbon fibers.

To illustrate the thermodynamic complexities that arise because of the presence of a molten metal in the directed metal oxidation process, a detailed analysis is presented both with and without the Si metal present. The former analysis is applicable to more traditional ceramic processing such as sintering or hot-pressing of SiC/Si3N 4 composites, whereas the latter is applicable to the directed metal oxidation process, or any other composite process where molten Si may be present.

The phase diagram for processing of a SiC/Si3N 4 composite when no molten Si is present is quite straightforward. It involves the reactions between Si3N 4 and SiC in a nitrogen gas. This interaction, which has been previously described when considering the hot-pressing of SiC whisker reinforced Si3N 4 (53), can be described by reactions (16) and (17):

3 SiC + 2N 2 = Si3N 4 4- 3 C (16)

Si3N 4 -- 3 Si + 2 N 2 (17)

where all species are assumed to be at unit activity. The nitrogen partial pressure can then be expressed as a function of temperature only, according to Eqs. (18) and (19):

1 0 In p N 2 -- + g A G [reaction ( 1 6 ) ] / R T (18)

1 In PN2-- ~AG 0 [reaction ( 1 7 ) ] / R T (19)

where 2~G ~ [reaction (i)] is the standard free energy of reaction i. Equation (18) then mathematically expresses the two-phase equilibrium between SiC and Si3N4, and Eq. (19) expresses the decomposition of Si3N4 into its constituent elements.


1

7

I 5

i O

t500 teoo 2i00 2400 2700 ~~erature

F~o. 11. SiC-Si3N 4 stability diagram for unit activity of all species.

Figure 11 from this analysis indicates that there is a wide two-phase region where SiC and Si3N 4 coexist. For example, for a partial pressure of nitrogen equal to 1 atm (lOgl0pN 2 = 0), SiC and Si3N 4 are stable between ~ 1750 and 2125 K. This temperature range is the processing window for SiC/Si3N 4 composites at 1 atm nitrogen pressure, in which no decomposition will occur in the absence of molten Si.

This analysis must be modified for the directed metal oxidation process or other processing where a molten Si phase is present. When this phase is included in the analysis, reactions (16) and (17) must be modified, becom- ing:

3 SiC + 2 N 2 -- Si3N 4 + 3 fin Si(l) (20)

Si3N 4 = 3 Si(/) + 2 N 2 (21)

where the underlined species are not pure (activity less than unity). Further, the decomposition of SiC must be considered, according to

SiC = Si(1) + Cinsi(/) (22)

The stability diagram then is described by Eqs. (23) through (25). The molton Si is assumed to be approximately pure, because the solubilities of


both carbon [54] and nitrogen [55] are very low near the melting point of Si:

3 In PN2 = + l A G ~ [reaction (20 ) ] /RT + ~ln a c (23)

' G~ (21 ) ] /RT In p N 2 - - ~ A (24)

In a c = - A G O [reaction (22 ) ] /RT (25)

From a phase equilibria viewpoint, the system has three degrees of freedom (e.g., a c, PN2 and T) rather than the two found previously. The net effect of this change is to relax the condition of unit activity of carbon because carbon can dissolve in the molten silicon phase. Using standard thermodynamic data (56), a stability diagram at fixed temperature can be drawn (Fig. 12), using Eqs. (23), (24), and (25). In this diagram, the thermodynamic stability of each phase at a given temperature occurs within a given range of nitrogen partial pressure and carbon activity. The stability region of each phase, SiC, Si3N4, and Si(/), will thus be shown as an area on a log a c versus log PN2 plot. Pure carbon in its standard state (graphite) occurs only when log a c is equal to zero (i.e., a c = 1.0.).

Figure 12 indicates that at any fixed temperature the three condensed phases, Si, SiC, and Si3N4, can coexist at equilibrium only at a fixed carbon activity and nitrogen partial pressure, i.e., at the triple point for that temperature. For example, at 1600~ this point occurs when log a c

-0.5

-1 O r -1.5 O

- J -2

-2.5

-3 -6

..................................................... / . , ' y . . . . . . . . . . . . . . . - - - ' - - - - - - - - - - - , , , ~ " .... s~

Si(I) 1800 C - - - 1700 C - - - 1600 C " - -

S I I I I I - -4 -3 -2 -" 0 1 2

Log (p N 2 , Pressure in atm)

FIG. 12. A Si-SiC-Si3N 4 stability diagram plotted for three different fixed temperatures: 1600, 1700 and 1800~ From Johnson [52 ].


FIG. 13. Micrograph of a SiC reinforced Si3N 4 matrix composite prepared by directed metal oxidation. From Johnson [52].

equals ~ -1 .5 and log P N2 equals ~ -1.6. During processing, the growing Si3N4/Si(t ) interface is in contact with 1 atm nitrogen so that the molten Si phase will nitride to form Si3N 4. Beneath this interface in the bulk of the composite, the nitrogen pressure and activity of carbon are not fixed by the external environment and will adjust to reach the equilibrium values given at the triple point for that temperature. Thus, the resulting composite will not degrade due to internal reactions either during growth, or subsequently under high-temperature operating conditions. These arguments have been verified experimentally in growths through SiC particulate preforms producing a Si3N4/Si matrix that shows little or no attack of the SiC particulate (Fig. 13).

With the analysis embodied in Fig. 12, the formation of carbon fiber reinforced Si3N 4 matrix materials by directed metal oxidation of molten silicon can also be described. To avoid fiber degradation during processing, the fiber should, ideally, be thermodynamically stable with respect to


the atmosphere, nitrogen; the molten metal, a silicon alloy; and the product, Si3N 4. Fig. 12 clearly shows, as expected, that carbon fibers are not stable with molten silicon and will react to form SiC. If the carbon fiber is coated uniformly with a dense layer of SiC, though, the thermodynamic analysis is exactly as described earlier. As far as the growing interface is concerned, the SiC coating on the fiber is the only component that will be exposed to the atmosphere, the Si3N 4 product, or the molten silicon. As long as the coating remains intact, the growth will occur much as it does for growth through SiC particulate fillers. Of course, should the SiC coating be defective or incomplete, then the molten Si will react with the underlying fiber to form SiC. Experimentally, these predictions are again verified. When using an uncoated or poorly coated carbon fiber, an extensive amount of SiC is observed in the matrix, but when a well-coated fiber is used (Fig. 14), the reaction proceeds essential as it does with particulate SiC.

FIG. 14. Micrograph of a C-fiber reinforced Si3N 4 matrix composite prepared by directed metal oxidation. From Johnson [52].


IV. Phase Equilibria in Carbide/Boride Systems

As described in Section II. B, a variation of the directed metal oxidation process has been used to form a ZrB 2 reinforced ZrC composite by the reaction of molten zirconium with a particulate bed of boron carbide. This process offers a particularly interesting example of the use of ternary phase diagrams in ceramic processing. The high processing temperature of 1850 to 2000~ and the presence of molten zirconium during reaction mean that the thermodynamic state predicted in the phase diagram is rapidly approached, thereby allowing the phase diagram to be an excellent predictor of the processing conditions. This section begins with a description of the zirconium system, and then describes the differences observed in the titanium and hafnium analogs.

A. Phase Equilibria in the ZrB 2 / ZrC / Zr System

As described earlier, the processing of Z r B 2 / Z r C / Z r composites by the directed metal oxidation process is accomplished by reacting molten zirconium with a porous bed of boron carbide in an inert atmosphere. Because the system contains only 3 components, Zr, B, and C, the Z r - B - C ternary phase diagram provides an excellent starting point for the phase analysis of this system. Fortunately, this ternary system has been studied previously [57]. This system contains no ternary compounds and only four binary compounds, ZrB2, ZrCx, B4C and ZrBI2. The ZrC x and B4C phases exist with a range of stoichiometry, e.g., x in ZrC x can vary between 0.6 and 1.0, while the two boride compounds have a limited range of stoichiometry.

The reaction given by Eq. (1) is very exothermic. To determine the temperature range of interest for thermodynamic analysis, the adiabatic reaction temperatures (ARTs) have been calculated for the Zr, Ti, and Hf systems (the latter two are described in the next section) using the following assumptions:

1. The appropriate reaction is

(3 + x ) M + B4C = 2MB 2 + MC1. 0 + xM (26)

where M is Zr, Ti, or Hr.

2. The reaction did not occur until after the metal melted, i.e., all reactants were preheated to the melting temperature of the metal. This is consistent with the differential thermal analysis (DTA) studies in the Zr-BaC system.


,

.

,

Once the metal melted, the system became isolated from the furnace by the molten metal phase, i.e., it became a constant volume system.

A small amount of Ar was trapped in the boron carbide preform. The number of moles was estimated from the ideal gas law using the melting temperature of the metal and the void volume in a typical boron carbide preform.

There was 10 mole % excess metal present. The calculated ARTs are presented in Table 1 for the Zr system, as well as the Ti and Hf analogs, which are discussed in the next section. The theoretical total pressures generated in all systems are greater than 1 atm. The composition of the gas phase varies from system to system, though it is composed primarily of metal and boron vapor in all cases suggesting that the ART is controlled by the highly endothermic vaporization of the metal and b o r o n ( A H v a p = 450 to 650 kJ/mol).

In practice, significant pressure buildup above 1 atm is highly unlikely because the graphite containment crucible is somewhat porous, and because the molten metal would not sustain it. Instead, one might expect gas bubble porosity to be present, which would lead to some increase in volume of the system. However, in the normal reactive infiltration mode, where metal is melted and reacted into a preform, reaching the ART and generating gas bubbles is highly unlikely for two interconnected reasons: high radiative heat losses and moderate reaction rates. The radiative heat losses, which scale as T 4, are enormous at these ultrahigh temperatures.

TABLE I

ADIABATIC REACTION TEMPERATURES (ARTs)FOR THE REACTION

OF GROUP IVB METALS AND BORON CARBIDE a

System

TBp Pure Pmetal P B at Pc at PAr at Ptotal ART Metal at ART ART ART ART at ART (~ (~ (atm) (atm) (atm) (atm) (atm)

Ti-B3C 3428 3357 1.583 0.021 0.002 0.305 1.19

Zr-B3C 4209 4429 .376 0.601 0.01 0.368 1.36

Hf-B4C 4388 b 4690 0.512 0.512 5 • 10 -7 0.383 1.22

aCalculations are for a constant volume of 10 cm 3, 1 • 10 -5 mol of Ar, 10 mole % excess metal and all reacting species preheated to the melting temperature of the metal. Products in each case are the liquid boride, liquid carbide, and liquid metal phases bEstimate based on estimated heat capacities and heats of melting for HfB2(I) and HfC(I) because these values are unavailable.


To support a temperature differential of ~ 2000~ between the furnace temperature and the ART, the reaction rates would have to be very rapid indeed. Yet, the measured reaction rates are fairly moderate [36], clearly not fast enough to prevent significant heat losses during the reaction. The ARTs are almost certainly not reached during processing, but they do provide a useful starting point for further analysis by suggesting that the sample temperature during processing is probably significantly higher than the actual furnace temperature.

As somewhat of an aside, the conclusion that the ART will not be reached need not always be true for these reactions if different processing approaches are used. For example, for reactions between powder mixtures of reactants, the self-propagating high-temperature synthesis (SHS) route, reaction rates are orders of magnitude higher, so that temperatures much closer to the ARTs are expected and bubble generation is much more plausible. Assuming the reactions were performed so the model assumptions described above hold, the Ti system would be most likely to produce bubbles because only in the Ti system is the metal boiling point lower than the ART. Moreover, the probability of reaching the ART is higher for the Ti system because of the relatively lower radiative heat losses expected at its lower ART temperature. This may, in fact, be one of the contributing factors to the high levels of porosity often observed in such SHS reactions [58].

These considerations indicate that the thermodynamic analysis of the Zr system must not be restricted only to the processing temperature of 1900~ The kinetics of this reaction have been studied using thermo- gravimetic analysis (TGA), interrupted growth experiments, ART calculations, and reaction velocity measurements. The proposed reaction mechanism is as follows [36]:

Molten zirconium reacts with B4C to form ZrB 2 and ZrC. Concur- rently, boron and carbon can dissolve into the liquid zirconium and graphite may precipitate at the interface with the B4C.

The high heats of these reactions raise the local temperature well above the nominal furnace temperature of 1900~

A B-rich liquid forms at the reaction interface with the B4C as a result of this temperature rise. The first B-rich liquid in the system outside of molten boron occurs at 2165~ from eutectic reaction of BaC , C, and ZrC1. 0. On further local heating (T > 2220~ a continuous B-rich liquid can form between pure boron and the eutectic liquid. The exact composition of the B-rich liquid is unknown, but it appears to lie somewhere within the triangle formed by B4C, C, and ZrB 2 .


4. The two-phase boride/carbide layer described in item 1 extends downward by reaction as the boron-rich liquid below it is drawn into the porous boron carbide by capillary action. Boron and carbon rapidly diffuse through this thin B-rich liquid layer, causing the two-phase layer to thicken as the directed reaction proceeds.

The results of that study estimated a maximum temperature between 2300 and 2400~ near the reaction interface. The ternary Z r - B - C phase diagrams [57] (Figs 15 and 16) provide an invaluable guide to formulating and understanding this proposed mechanism, as described more fully next.

As shown in the isothermal ternary cross section (Fig. 15), at the furnace temperature of 1900~ there is a significant solubility of boron and carbon in molten zirconium, and there is no other liquid phase present in the system. At the higher temperatures expected at the reaction interface, though, there is also a boron-rich liquid present. This phase field becomes larger as the temperature is increased. It forms first when B melts at ~ 2100~ and grows as the temperature increases, forming through a class II ternary reaction at 2160~ with ZrBI2, from decomposition of B4C, and from ZrB 2. Concurrently, a second B-rich liquid forms from a ternary eutectic at 2165~ between B4C , ZrB2, and C. As the temperature is increased further, these two liquid fields become one at 2220~ at the pseudobinary eutectic between ZrB 2 and B4C. At 2390~ the ZrB2-C tie-line (quasibinary) goes through eutectic melting to this boron-rich liquid (Fig 16). The corresponding Z r C • ZrB 2 pseudobinary eutectic melting does not occur until a much higher temperature, 2830~ [10].

Using the Z r - B - C phase diagrams, some important features about the process of forming the Z r B z / Z r C / Z r composites by directed metal oxidation are apparent. With a maximum interface temperature of 2300 to

FIG. 15. [57].

Zr

B B ZrBlz~B4C Z//~B4 c

ZrB 2

ZrCx v Zr ZrCx C

T = 1900 ~ T - 2400 ~

The Zr-B-C isothermal ternary cross sections at 1900 and 2400~ After Rudy


o T,C

m

2800 - L1, zr.

m

4oo - . . . . .

/ 2 0 0 0 ~ / L 1 + Z r / L I + Z r B 2 + ZrCx

/

I ~ , ~ ZrB 2 + ZrC x

, rCx

z r ' l l , Zr+ ZrB2+ ZrC x

c c

, Z r C l l

m

L2

L2+ZrB

~- ZrB2+~"

�9 ZrB 2 + B 4 C

C + ZrB 2

I II

0 20 40 60 80 100 Mole % B C

0.8 0.2

~ L 2 + B 4 C

�9 L2+B4C+ ZrB2

\ B4C

FIG. 16. The Zr-B4C isopleth, where the dotted line is the proposed reaction pathway. From Johnson et al. [36].

2400~ an approximate reaction pathway as shown by the dashed line in the Zr-B0.sC0. 2 isopleth of Fig. 16 can be suggested. The observed phases in an interrupted growth experiment [36] were consistent with the proposed reaction pathway, except for a Z r B z / Z r C x / C three-phase layer that is predicted to occur between the ZrB2/C layer and the ZrBz/ZrC x layer. Although this layer was not observed, it should be difficult to detect because the carbon would occur only as a very finely dispersed minor constituent in the bottom of the ZrBz/ZrC x layer. Although the reaction pathway shown in Fig. 16 is an approximation because it assumes that the phase stoichiometry does not deviate off the Zr-BaC tie-line, this simpli- fication was found to be sufficient to explain the observed reaction behavior.

The proposed mechanism and the pathway shown in Fig. 16 should lead to the following sequence of layers, from top to bottom during processing at the reaction temperature: unreacted molten Zr, the ZrBz/ZrC0.6/Zr product layer, a relatively thin solid two-phase ZrBz/ZrC x layer, a B-rich liquid layer, and finally any unreacted B4C. It appears from the observed microstructure that the B-rich liquid at the reaction interface consumes, at least partially, the B4C reactant, because the C-tendrils and ZrB2, which come from this liquid on freezing, surround mostly unreacted B4C particles. During the reaction, then, there is an activity gradient across the sample, with the zirconium activity decreasing and carbon and boron activity increasing in going from top to bottom. Finally, when all the B4C is consumed, the activity gradient cannot be maintained, the B-rich liquid


disappears, and the sample homogenizes to give the final Zr/ZrC0.6/ZrB2 product.

Although there is an activity gradient across the reacting system, the activity of the components in any given layer is not easily determined because superimposed across the above sequence of reaction layers is a large temperature gradient, a result of the exothermic reactions. Because the exothermic reactions involve the formation of ZrB 2 and ZrC x, the maximum temperature is expected near the ZrBz/ZrC/B-r ich liquid interface where these reactions occur. One consequence of the temperature gradient is that, unlike a ternary diffusion couple, a three-phase region can exist over an extended region without violating the phase rule.

The microstructure of the ZrBz/ZrC0.6/Zr layer coarsens with time at temperature (Ostwald ripening), giving rise to the observed variation in feature size through the reaction product layer. Boron and carbon also diffuse into the unreacted liquid zirconium at the top, which when cooled gives the solidification microstructure observed in an interrupted growth experiment. When left for longer times at temperature, grain growth allows the entire sample to homogenize, giving the observed substantially uniform feature sizes seen from top to bottom in the final product.

The proposed mechanism and its attendant use of the appropriate phase diagrams has been useful in its processing implications for the formation of Z r B z / Z r C / Z r and related composites formed using this processing approach. As the phase diagram suggests, and the experimental results confirm, a ZrBz-ZrC two-phase layer forms below the Z r B z / Z r C / Z r product layer. To react with the B4C , the Zr reactant must be transported through both of these layers, and in so doing, both layers thicken with time. Thus, the reaction rate will decrease with increasing time at temperature because of the increased distance through which the Zr must be transported, and as a result the reaction rate is parabolic. The observed parabolic rate constant is large enough that very thick parts, as thick as nearly 4 in., can be achieved in only 2 h at 1900~ Further, by increasing the time at temperature, thicker components should be possible as well. Thus, the parabolic nature of the reaction kinetics does not limit the utility of the process.

The final step in the proposed mechanism suggests that Ostwald ripening is responsible for homogenizing the finals product. This also has processing implications. For very thick components, even though the reaction may be complete (i.e., all of the B4C is consumed) in a short time, additional time at temperature may be desirable to achieve a more uniform product through the thickness of the component. Or, to put it another way, at a fixed processing time of, say, 2 h, thin components will tend to be more uniform than thick components, because for the thin

118 W I L L I A M B. J O H N S O N A N D ALAN S. N A G E L B E R G

component the reaction is complete early in the dwell so the bulk of the time at temperature is essentially an anneal that allows homogenization.

Finally, the proposed mechanism for the formation of the ZrB2/ZrC/Zr composites has been useful in understanding and formation of similar systems, e.g., the Ti and Hf analogs. These results are described more fully in the next section.

B. The Ti-B-C and Hf-B-C Systems

The Ti -B-C phase diagrams are very similar to the Z r - B - C ternary diagram except that in the Ti system there are three stable borides--TiB, TiB2, and a low-temperature phase, TisB4--whereas in the Zr system there is only one stable boride, Z r B 2. Based on the mechanism presented previously, the presence of the additional boride phases has important implications regarding the formation of boride/carbide composites in the Ti system. A comparison of the two ternary systems at just above the melting temperatures of the metal phase and at a higher temperature (Figs. 15 and 17) reveals several features. At the reaction temperature the liquid phase in the Ti system is in equilibrium with TiB and TiC, so that these phases will form in preference to T iB 2 and TiC. The formation of the TiBz-containing composite is inhibited by the presence of the TiB-TiC-Ti three-phase triangle, which "gets in the way" by reacting to form TiB rather than TiB 2. In the Zr case, on the other hand, the liquid

B B A..c A..c "'K k

Ti C Ti C TiC TIC

x ](

0 0 T - 1 7 0 0 - C T - 2 4 0 0 C

FIG. 17. The Ti-B-C isothermal ternary cross sections at 1700 and 2400~ After Rudy [571.


can be in equilibrium with ZrB 2 and ZrC so that it is possible to readily form the Z rBz-ZrC-Zr composite.

An additional factor contributes to the kinetics of the reaction between Ti and B4C. The TiB2-C pseudobinary goes through eutectic melting at a higher temperature than in the ZrB2-C system. Thus, at the higher temperatures that occur at the reaction front, the postulated boron-rich liquid is not in equilibrium with the carbide as it is in the Zr case, but rather is in equilibrium with the titanium diboride phase. This implies that the reaction pathway must go through a region containing TiB 2 or C in the Ti system, whereas that is unnecessary in the Zr system. Using these pathways, the following layers might be expected to be observed in interrupted growth experiments in the Ti system progressing from the infiltrating metal to the unreacted B4C: Ti, T iB/TiC/Ti , TiB/TiC, TiBz/TiC , TiB2, TiBz/BaC/C, and B4C. The monoboride phase is observed even though it is not in the reaction pathway because away from the reaction interface the temperature is low enough so that TiB is stable. Although detailed studies of this system have not been completed, prelimi- nary interrupted growth studies indicate that all these layers are observed. Thus, during reaction, the Ti must penetrate through three solid layers, TiB/TiC, TiBz/TiC, and TiB 2, whereas in the Zr system, the Zr need only go through one solid/solid layer (ZrBz/ZrC) to reach a BaC-containing layer.

One potential route around this problem is to infiltrate above the melting temperature of TiB, i.e., above 2190~ In this case, the molten Ti will be in equilibrium with TiB 2 and TiC, rather than TiB and TiC. The reaction conditions will then become similar to those in the Zr system. On cooling below 2190~ though, the residual Ti in the composite should, at least thermodynamically, form TiB, according to a class II ternary reaction:

Ti(l) + TiB 2 = TiB + TiC (27)

This reaction, though, will be kinetically slow because B must diffuse through a TiB product phase that will form on the TiB 2. By performing the reaction in an inductively heated furnace in which these very high temperatures can be reached, a TiBz/TiC/Ti composite has been formed by the reaction of Ti with B4C [59].

The preceding analysis predicts that it will be difficult to form a TiBz/TiC/Ti composite using furnace infiltration temperatures near the Ti melting point. On the other hand, a different composite, a T iB /T iC/T i material, has been successfully formed by reacting at ~ 2000~ for times of 6 to 24 h. Under these conditions, molten Ti is in equilibrium with the

120 W I L L I A M B. J O H N S O N A N D A L A N S. N A G E L B E R G

TiB and TiC solid phases. Although the solid layers described here still form, they tend to be thinner. Coupled with the additional processing time, the Ti flux is then high enough to form the TiB/TiC/Ti product.

The Hf -B-C system presents a situation that falls somewhat between the Ti and Zr systems [60]. Although the HfB phase is stable in the Hf-B binary system, it melts at 2100~ below the melting point of the Hf parent metal (2227~ During a directed metal oxidation reaction of molten Hf with B4C at just above the melting point of Hf, e.g., 2400~ the H f - B - C isothermal ternary cross section (Fig. 18) indicates that the molten metal is

B B

HfB

Hf C Hf HfC C

T = 2000 oC =

H f L ~ ' ~ C

T = 2 7 0 0 ~

F~6. 18. The Hf-B-C isothermal ternary cross sections at three different temperatures. Points A and B represent two possible stoichiometries for processing hafnium boride/hafnium carbide composites. After Rudy [57].


in equilibrium with the diboride and carbide, just like the Zr system. Furthermore, at the higher temperatures expected at the reaction front, the boron-rich liquid phase is in equilibrium with HfB 2 and HfC, again similar to the Zr system. Thus, infiltration behavior should be very similar to that observed in the Zr system. On cooling below 2100~ though, the residual Hf in the composite should, at least thermodynamically, form HfB according to a class II ternary reaction:

Hf(l) + HfB 2 = HfB + HfC (28)

As described just for the Ti system, however, this reaction will be kinetically slow because B must diffuse through a HfB product phase that will form on the HfB 2. With relatively rapid cooling, the H f B z / H f C / H f composite can be formed. Depending on the original stoichiometry, the formation of the HfB phase can be enhanced or limited. By reacting under conditions where there is a large excess of Hf metal, e.g., point A on Fig. 18, there will be more tendency to form HfB because the HfB can precipitate out from the excess molten Hf during cooling from the reaction temperature. On the other hand, if the amount of Hf is more nearly stiochiometric, (i.e., that amount of Hf required to form HfB 2 and HfC from B4C) , e.g., point B on Fig. 18, the HfB will form according to Eq. (28), which as described earlier will be inhibited kinetically, thereby limiting the HfB content in the final composite. These effects have, in fact, recently been experimentally demonstrated in the Hf-B4C system [61].

V. Conclusions

In this chapter we have attempted to describe various thermochemical techniques for dealing with processing of complex composite systems. Although the discussion focused primarily on the directed metal oxidation process, the general approaches used are applicable to the processing of any composite material. These systems are typically multiphase systems, which can include a gas phase, a multicomponent liquid metal phase, and one or more solid ceramic phases. Although this complexity generally precludes a detailed quantitative analysis because of a lack of thermodynamic data, various qualitative predictions have been used to illustrate general trends, and to indicate expected effects on processing. Binary and /o r ternary phase diagrams have been used as a starting point for thermochemical analysis. In some cases, for example, the processing of carbide/boride composites, the processing temperatures are high enough


and the ternary phase diagrams complete enough, so that little additional phase diagram data are necessary.

Generally, though, additional equilibrium diagrams have been constructed to provide additional insight into the composite formation process. The predictions about actual processing conditions obtained from these diagrams must, of course, be tempered with the knowledge that conditions during the actual process are not in equilibrium. Yet, local equilibrium may hold, and the equilibrium diagrams can be useful in understanding mechanisms if kinetic effects are considered, as they were for the directed metal oxidation of molten aluminum alloys to form alumina. Furthermore, the phase diagrams provide an invaluable tool for predicting potential reaction products between constituents within the composite. Again, kinetics may not be fast enough to yield the thermodynamically stable product, but the diagram at least does indicate what to anticipate. Should undesirable products be found, means to prevent their formation may also become apparent from the thermochemical analysis. Such was the case in the prevention of AlaC 3 formation in A1 MMCs by the addition of Si to the aluminum. The phase diagrams are thus an invaluable, but not infallible, tool, which, if used carefully, can be predic- tive and can help solve problems that arise during composite processing.

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19. Y. Kgawa, A. Okura, and S. Watanabe, In-situ formation of A1203/Si composites by directed oxidation of liquid AI alloy. Achievements in composites in Japan and United States. Proc. Jpn.-U.S. Conf. Compos. Mater. 5th, 1990, pp. 79-86 (1990).

20. P. Xiao and B. Derby, Study of the formation of AI203/A1 composites by oxidation. Eur. Ceram. Soc. Cony., 2nd, Augsburg, 1991.

21. S. Kharti, M. Koczak, T. Chang and Y. Kagawa, Analysis of A1203/AI-Si composites fabricated by directed oxidation of aluminum alloys. Proc. Int. Symp. Sci. Eng. Ceram. 1st, pp. 19-24.

22. P. Hing and G. W. Groves, The strength and fracture toughness of polycrystalline magnesium oxide containing metallic particles and fibers. J. Mater. Sci. 7(4), 427-434 (1972).

23. B. Flinn, M. Ruhle, and A. G. Evans, Toughening in composites of A120 3 reinforced with AI. Acta Metall. 37(11), 3001-3002 (1989).

24. J. T. Blucher, Discussion of a liquid metal pressure infiltration process to produce metal matrix composites. J. Mater. Proc. Technol. 30, 381-390 (1992).

25. R. Mehrabian, R. G. Riek, and M. C. Flemings, Preparation and casting of metal-particulate nonmetal composites. Metall. Trans. 5(8), 1899-1905 (1974).

26. A. R. C. Westwood, Materials for advanced studies and devices. Metall. Trans. A 19A, 749-758 (1988).

27. M. K. Aghajanian, J. T. Burke, D. R. White, and A. S. Nagelberg, A new infiltration process for the fabrication of metal matrix composites. SAMPE Q. 20(4), 43-46 (1989).

28. M. K. Aghajanian, M. A. Rocazell, J. T. Burke, and S. D. Keck, The fabrication of metal matrix composites by a pressureless metal infiltration technique. J. Mater. Sci. 26(2), 447-454 (1991).

29. A. S. Nagelberg, Growth kinetics of A1203/metal composites from a complex aluminum alloy. Solid State Ionics 32-33, 783-788 (1989).

30. A. S. Nagelberg, S. A. Antolin, and A. W. Urquhart, Formation of A1203/metal composites by the directed oxidation of molten aluminum-magnesium-silicon alloys. Part II. Growth kinetics. J. Amer. Ceram. Soc. 75(2), 455-462 (1992).


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3,

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E. Breval and A. S. Nagelberg, Microstructure of an Al203/metal composite containing an AI203 filler material. Mater. Res. Soc. Symp. Proc. 132, 93-98 (1989). A. S. Nagelberg, The effect of processing on the growth rate and microstructure of A1203/metal composites. Mater. Res. Soc. Symp. Proc. 155, 275-282 (1989). S. Antolin, A. S. Nagelberg, and D. Creber, Formation of A1203/metal composites by the directed oxidation of molten aluminum-magnesium-silicon alloys. Part I, Microstruc- tural development. J. Am. Ceram. Soc. 75(2), 447-454 (1992). W. B. Johnson, T. D. Claar, and G. H. Schiroky, Preparation and processing of platelet reinforced ceramics by the directed reaction by zirconium with boron carbide. Ceram. Eng. Sci. Proc. 10(7-8), 588-598 (1989). T. D. Claar, W. B. Johnson, C. A. Andersson, and G. H. Schiroky, Microstructure and properties of platelet reinforced ceramics formed by the directed reaction of zirconium with boron carbide. Ceram. Eng. Sci. Proc. 10(7-8) 599-609 (1989). W. B. Johnson, E. Breval, and A. S. Nagelberg, The kinetics of formation of a platelet reinforced ceramic composite prepared by the directed reaction of zirconium with boron carbide. J. Am. Ceram. Soc. 74 (9) 2093-2101 (1991). E. Breval and W. B. Johnson, Microstructure of platelet reinforced ceramic prepared by the directed metal reaction of zirconium with boron carbide. J. Amer. Ceram. Soc. 75(8), 2139-2145 (1992). V. A. Ravi, T. D. Claar, J. A. Hornor, and W. B. Johnson, Platelet reinforced ceramics for severe thermal shock applications. In Composites for High Temperature Applica- tions" (V. A. Ravi, ed.), pp. 175-183. Metallurgical Society, Warrendale, PA, 1992. K. C. Vlach, O. Salas, H. Ni, V. Jayaram, C. G. Levi, and R. Mehrabian, Nucleation and growth of alumina/metal composites by oxidation of aluminum alloys. J. Mater. Res. 6(9), 1982-1995 (1991). A. S. Nagelberg, Observations on the role of Mg and Si in the directed oxidation of AI-Mg-Si alloys. J. Mater. Res. 7(2), 265-268 (1992). K. C. Vlach, C. G. Levi, and R. Mehrabian, On the kinetics of the oxidative growth of alumina-metal composites. In "Processing and Fabrication of Advanced Materials for High Temperature Applications" (V. A. Ravi and T. S. Srivatsan, eds.) pp. 157-172. Metallurgical Society, Warrensdale, PA, 1992. A. S. Nagelberg, A. S. Fareed, and D. J. Landini, Production of ceramic matrix composites for elevated temperature applications using the DIMOX T M directed metal oxidation process. In "Processing and Fabrication of Advanced Materials for High Temperature Applications" (V. A. Ravi and T. S. Srivatsan, eds.), pp. 127-142. Metallur- gical Society, Warrensdale, PA, 1992. C. H. P. Lupis, "Chemical Thermodynamics of Materials," pp. 285-292. Elsevier, New York, 1983. E. Breval, M. K. Aghajanian, J. P. Biel, and S. Antolin, The structure of aluminum nitride/aluminum and aluminum oxide/aluminum composites produced by the directed oxidation of AI. J. Am. Ceram. Soc. 76 (7), 1805-8 (1993). S. Antolin, M. Aghajanian, M. Rocazella, R. Kantner, and A. Nagelberg, private communication. L. L. Oden, and R. A. McCune, Phase equilibria in AI-Si-C system. Metall. Trans., A 18A, 2005-2014 (1987). J.C. Viala, P. Fortier, and J. Bouix, Stable and Metastable phase equilibria in the chemical interaction between aluminum and silicon carbide. J. Mater. Sci. 22, 1842-1850 (1990). D.-J. Lee, M. D. Vaudin, C. A. Handwerker, and U. R. Kattner, Phase stability and interface reactions in the AI-SiC system. Mater. Res. Soc. Symp. Proc. 120, 357-365 (1988).


49. J. L. Murray and A. J. McAlister, The AI-Si (aluminum-silicon) system. Bull. Alloy Phase Diagram, 5(1), 74-84, 89-90 (1984).

50. M. W. Chase, Jr., C. A. Davies, J. R. Downey, Jr., D. J. Frurip, R. A. McDonald, and A. N. Syverud, J. Phys. Chem. Ref. Data 14, Suppl. 1 (1986).

51. D. J. Lloyd and E. Dewing, Stability of SiC in molten aluminum, in Proc. Int. Symp. Adv. Struc. Mater. 1988, pp. 71-77 (1989).

52. W. B. Johnson, Reinforced Si3N 4 matrix composites formed by the directed metal oxidation process. Ceram. Eng. Sci. Proc. 13(7-8), 573-580 (1992).

53. K. G. Nickel, M. J. Hoffmann, P. Greil, and G. Petzow, Thermodynamic calculations for the formation of SiC-whisker reinforced Si3N 4 ceramics. Adv. Ceram. Mater. 3(6), 557-562 (1988).

54. L. L. Oden and R. A. McCune, Phase equilibria in the A1-Si-C system. Metall. Trans., A 18A(12), 2005-2014 (1987).

55. W. Kaiser and C. D. Thurmond, Nitrogen in silicon. J. Appl. Phys. 10(3), 427-431 (1959). 56. M. W. Chase, Jr., C. A. Davies, J. R. Downey, Jr., D. J. Frurip, R. A. MacDonald, and

A. N. Syverud, J. Phys. Chem. Ref. Data 14 (1985). 57. E. Rudy, Ternary phase equilibria in transition metal-boron-carbon-silicon systems. Part

V. Compendium of phase diagram data. Tech. Rep. AFML-TR--Air Force Mater. Lab. (U.S.) AFML-TR-65-2 (1969).

58. Z. Munir, Synthesis of high temperature materials by self-propagating combustion methods. Ceram. Bull. 67(2), 342-349 (1988).

59. D. White, personal communication. 60. V. A. Ravi, personal communication. 61. V. A. Ravi, Formation of HfB2/HfC/Hf composites. Scr. Metall. (1993).

Use of Phase Diagrams in the Study of Silicon Nitride Ceramics

TSENG-YING TIEN

Department of Materials Science and Engineering University of Michigan

Ann Arbor, Michigan 48109

I. I n t roduc t i on . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 127

II. R e p r e s e n t a t i o n of Silicon N i t r i d e - M e t a l Oxide Systems . . . . . . . . . . . . . . . 128

III. S i 3 N 4 - M e t a l Oxide Systems . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 131

A. Silicon Ni t r ide Systems with O n e Meta l Oxide . . . . . . . . . . . . . . . . . . . 131

B. Silicon Ni t r ide Systems with Two Meta l Oxides . . . . . . . . . . . " . . . . . . . 137

IV. Alloy Des ign . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 140

A. A c i c u l a r / 3 - S i 3 N 4 Gra ins with G r a i n - B o u n d a r y Phases . . . . . . . . . . . . . . 141

' B. T w o - P h a s e Compos i t e s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 151

V. Conclus ion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 154

R e f e r e n c e s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 155

I. Introduction

The family of silicon nitride ceramics is one of the most promising materials for structural applications. The superior thermal and mechanical properties of these materials is due to the highly covalent nature of their chemical bonds. Because of the covalent nature of the bonds, solid-state diffusion is very slow, thus preventing densification of the silicon nitride during sintering. However, if a liquid phase is introduced into the system during sintering, high densities are achieved. This liquid phase is produced by the addition of sintering additives to the silicon nitride, which are usually metal oxides that form a low-melting-point eutectic liquid with the oxide surface layers of the silicon nitride powder. During and after densification the starting silicon nitride powders transform from ce-Si3N 4 to /3-Si3N 4 by the solution-precipitation process. The newly formed /3-Si3N 4 phase has an elongated hexagonal rod morphology, which forms



128 TSENG-YING TIEN

an interlocking microstructure and is responsible for the superior mechanical properties of these materials.

Although densification is facilitated by the use of oxide additives, on cooling, the high-silica-containing liquid becomes glass and envelops the acicular/3-Si3N 4 grains, forming an intergranular phase. This intergranular glassy phase becomes soft above the glass transition temperature, resulting in reduced creep resistance. However, to counteract this problem, the glass layer can be crystallized or reacted with the major phase. This then increases the high-temperature mechanical properties of silicon nitride ceramics.

A number of sintering additives are used to densify silicon nitride, each affecting the properties of the intergranular liquid phase differently, thus affecting the properties of the crystallized phase differently. Different compositions of the liquid affect the growth kinetics and the morphology development of the silicon nitride grains and the characteristics of the grain-boundary phase. For developing silicon nitride ceramics with optimum mechanical properties, the knowledge of the phase relationships of silicon nitride-metal oxide systems is essential.

II. Representation of Silicon Nitride-Metal Oxide Systems

An excellent representation of Si3N4-metal oxide phase diagrams can be found in Gauckler's paper [1]. Gauckler's representation can be illustrated using the Si3N4-AI203 system. Gauckler considered the silicon nitride-metal oxide systems to be a reciprocal salt systems. These systems are actually one section in the quaternary system, Si-A1-N-O. The basic concept is based on the fact that the Si3N 4 + 2A1203 = 3SiO 2 4- 4A1N reaction is reversible, and the valency states of these elements are nonvariant. In this case, the phase diagrams of this system can be expressed as a quadrilateral using the four compounds as components at the corners (Fig. 1). When a diagram is plotted in equivalent percent, i.e., the total charges at each corner are equal, the diagram is a square plane as shown in Fig. lb. When the diagram is plotted in atomic percent, the diagram is a trapezium (a quadrilateral, no two sides of which are parallel) as shown in Fig. la.

According to the chemical reaction, Si3N 4 + 2A1203 = 3SiO 2 + 4A1N, the compositions in these systems contain three independent variables and one dependent variable, i.e., when three of the components are fixed, the fourth component is determined according to the chemical reaction. Fig. 2 is the phase diagram of the Si3N4-SiOz-AlzO3-A1N system pre-

4 S T U D Y O F S I L I C O N N I T R I D E C E R A M I C S 129

AI AI

a / / ~ b

AI203 03 / /~ " AIN / / f " "~A IN

_/,,. -<" :7 , ,X,, Si Si3N4 N Si Si3N 4

Atomic % . Equivalent %

FIG. 1. The quaternary system S i -AI -N-O [1]. (a) The Si3N4-SiO2-AI203-AIN system is plotted in atomic percent. The phase

diagram of this system is a trapezium. It is a quadrilateral; no two sides of which are parallel.

(b) The Si3N4-SiO2-AI203-A1N system is plotted in equivalent percent. The phase diagram of this system is a square.

SiC Schm Iz 3AI C~SiO 2 AI203

1; " ~ //X/"1 "'~

q; ,' "/ ll/I//llllr , i / \

4~ !1 I ~ liN Il lb.o -~

~ g Si2N20

0 20 40 60 80 100 Si 3 N 4 Equi volent % A! A! N

FIG. 2. The Si3N4-SiO2-AI203-A1N system. The diagram was plotted in equilibrium percent by Gauckler [2]. This was the first time ever that this was presented as a reciprocal salts system.

130 TSENG-YING TIEN

Be

N 0

A1N " ~ ~ V A 1203

A1

FIG. 3. The S i - A 1 - B e - N - O system. The subsystem S i 3 N 4 - S i O 2 - A l 2 0 3 - AIN-BeO-Be3N 2 is shown as a triangle prism. The diagram was plotted in equivalent percent [ 1].

sented by Gauckler [2]. Note the area inside the square contains binary joins and ternary compatibility triangles. These compatibility triangles can be treated individually and plotted either in weight, atomic, or mole percent.

The silicon nitride system with two oxide additives can be expressed as a triangular prism [1]. Fig. 3 shows the Si3N 4 system with two oxides additives, A1203 and BeO. This is a part of the S i -A1-Be-N-O system. When the valency states of these elements are nonvariant, this five- component reciprocal salt system consists of three quasiternary reciprocal salt systems (square planes) and two ternary systems (triangles). As shown in Fig. 3, the three reciprocal salt systems are Si3N4-SiOz-A1N-AI203, Si3N4-SiOz-Be3Nz-BeO, and AIN-A12Os-Be3Nz-BeO. The two ternary systems are Si3Na-A1N-Be3N 2 and SiO2-AlzO3-BeO. When they are plotted in equivalent percent, the system becomes a triangular prism. Below the subsolidus temperatures, the maximum number of phases that can coexist under equilibrium conditions will be four. Inside the triangular prism, the space will be filled with compatibility tetrahedrons. Each tetrahedron can be replotted in either weight, more, or atomic percent. An example of this system is given in Fig. 4 for the

4 STUDY OF SILICON NITRIDE CERAMICS 131

Y 203(Y ')

YN

3As(G)

/ / _ I - \ ~ ~ ~ A3S2(M)

Si3N 4

AIN(AN)

FIG. 4. The subsolidus phase relationships in the Si3N4-SiO2-AI203-A1N-Y203-YN system, showing compatibility tetrahedrons [3].

Si3N4-SiO2-A1N-A1203-YN-Y203 system determined by Naik and Tien [3] and Sun et al. [4]. Sixty-eight compatibility tetrahedrons in the S i -A1-Y-N-O system have been established by these authors.

This chapter is not intended to be a literature survey, since an excellent review paper on phase diagrams of silicon nitride-metal oxide systems has been written by Sorrell [5]. This chapter, is, however, intended to illustrate the use of phase diagrams to study silicon nitride ceramics.

III. Si3N4-Metal Oxide Systems

A. Silicon Nitride Systems with One Metal Oxide

1. Si3N4-MgO SYSTEM

Magnesium oxide was the first sintering additive ever used to densify silicon nitride ceramics [6]. There are three different versions of the phase diagram published in the literature [7]. A composite diagram is given in Fig. 5 [8], showing the phase boundaries determined by three independent researchers, Jack [9], Inomata [10], and Lange [11]. There is obvious disagreement between these authors as the location of the phase boundaries, illustrating the difficulty in determining the true phase relationships in the silicon nitride-metal oxide systems. The term phase diagram

132

FIG. 5.

TSENG-YING TIEN

----Inomata - - L a n g e .......... Jack

SiO 2 MgSiO 3 Mg 2 SiO 4 MgO

/,/i/ / / /

Si2N20?7/i/ / ! / /Mg.N20 I ! , t / I / i i.--- .....

i'.,' ill. ........................... Si3N 4 MgSiN 2 Mg4SiN 4 Mg3N2

A composite plot of the Si3N4-SiO2-MgO-Mg3N 2 system [8-11].

implies that the phases exist under equilibrium conditions. Because of the high vapor pressure of SiO, however, the vapor phase cannot be ignored, and therefore silicon nitride-metal oxide systems should not be considered as condensed systems. However, it is very difficult to take the vapor phase into consideration in determining the phase diagrams. Miiller et al. [8] have shown that when mixtures of Si3N 4 and MgO were heat treated in a furnace with a large, open volume, the Si3N4-MgO join exists. When these mixtures were heat treated in tightly covered crucibles, the MgzSiO4-MgSiN 2 join was observed. These results indicate that the phases diagram published by Jack [9] and Inomata [10] were determined in a near closed system while Lange [11] determined the diagram in a more open system. However, the Si3N4-MgzSiO 4 join, was reported by all of the authors.

For processing silicon nitride ceramics with MgO as a sintering additive, only the compositions at either side of the Si3N4-MgzSiO 4 join are of interest, because most ceramics have a composition that lies on this join. During sintering, a silica-rich liquid forms, which aids densification. After cooling, the liquid becomes glass and is located at the grain boundaries. During low-temperature heat treatment, forsterite crystallizes from the glass. Compositions lying to the left of the Si3N4-MgzSiO 4, join have a lower eutectic temperature. The compositions lying to the right of the Si3N4-MgzSiO 4 join fall in the triangle Si3N4-MgzSiO4-MgSiN 2. The

4 STUDY OF SILICON NITRIDE CERAMICS 1 3 3

eutectic temperature in this triangle is higher, hence the densification will be very difficult, but this refractory grain-boundary phases would be more desirable for high-temperature applications.

2. Si3N4-Y203 SYSTEM

Four different versions of the phase diagram for the Si sN4-S iOz-YzO3-YN system have been published. Fig. 6 shows the latest version of the diagram, which was determined by Hohnke et al. [12]. Among these diagrams, the key disagreements between them are the number of ternary oxynitride compounds and the compositions of these compounds. However, for the purpose of processing of silicon nitride ceramics, only compositions in the compatibility triangles on either side of the Si3N4-Y2Si20 7 join are essential. Fortunately, in this region, there are no differences among the diagrams published.

For compositions on the Si3N4-Y2Si20 7 join, liquid forms at the binary eutectic temperature below 1550~ and it crystallizes at 1500~ Composi- tions with higher Y203 contents than compositions on the join will fall in the compatibility triangle outlined by Si3N4-Y2SizOv-Y10(SiO4)6N 2 (apatite), which has a lower eutectic temperature than the triangle outlined by Si3N4-Y2SizOy-Y10(SiO4)6N 2. The liquid composition may not affect the microstructural development. However, the composition of the

Si02 Y2 Si207 u Y203

a 90IX ,' .7~" ~ " / l ~ Y l o ( S i O ~

,

To ~~ F ~ ~ / / / / / / f - - / / ~ . , o~ . .~ so ll

,o

3oV ///// ,2s,,o,,,,, ,,.,.,.,,,,n., s,.,.,o[ / / / / ,o~f

Si3N 4 10 20 30 40 50 60 '70 80 90 YN Eq.-% Y

"~ 70 N2

16~ h' I I I / / 1 1 " / / o I // / / / / / ,r

~oLr o.,::,..~

' ~ Eq.- %Y ----

FIG. 6. The Si3N4-SiO2-Y203-YN system [12].

(a) 1550~ isotherm. (b) 1500~ isotherm.

134 TSENG-YING TIEN

grain-boundary phase will be different, which will affect the properties of the ceramic. Compositions with lower Y203 contents fall in the compatibility triangle outlined by Si3N4-Y2SizOy-SizN20. Compositions in this triangle are more refractory than the compositions in the triangle formed by Si3N4-Y2SizOy-Ylo(SiO4)6N 2 (apatite).

3. Si3N4-A1203 SYSTEM

Oyama [13] and Jack [14] reported the solid solution formation of aluminum oxide in silicon nitride. In these solid solutions, aluminum ions replace silicon ions and nitrogen ions replace oxygen ions in the lattice. Oyama [13] and Jack [14] both presented their phase diagrams as triangular ternary systems. These solid solutions were named by Jack [14] as /3'-SiA1ON and he has suggested that these solid solutions contain lattice vacancies and hence an increase in the ionic nature of the chemical bond. If this is true, a high-density silicon nitride ceramic should be able to be made by solid-state sintering through lattice diffusion without the aid of a liquid phase, and thus without the need of sintering additives. A polycrystalline material without a grain-boundary liquid phase should therefore have better creep resistance. The inferior high-temperature properties of sintered and hot-pressed silicon nitride ceramics have been attributed to the glassy grain-boundary, which becomes soft at high temperatures [15]. Thus, it was believed that these single-phase silicon nitride solid solutions (/3-SiA1ONs) will not contain grain-boundary glass if properly made. This possibility has led to many investigations in an attempt to obtain silicon nitride ceramics with no grain-boundary glass and hence better high- temperature properties.

In a later study, Gauckler [2] demonstrated that these solid solutions are substitutional and contain no lattice vacancies. These solid solutions were described by Gauckler by the formula ~-Si6_xAlxN8_xO x. The limiting composition has an x value of 4. The phase diagram of the Si3N4-SiOz-A1N-AI203 system is shown in Fig. 2 as a reciprocal salt system. Gauckler [2] was the first to present the phase diagram in this manner.

Boskovic et al. [16] sintered compacts of presynthesized /3'-SiA1ON powder and found that no densification occurred at temperatures as high as 1900~ In these studies, Boskovic showed the sintering behavior of some selected compositions in the Si3Nn-SiOz-A1N-A120 3 system, using different combinations of the individual compounds as the starting materials, to produce a single-phase, ~-Si6_xAlxNs_xO x solid solution. Since the


Si0 2

/ /

Si 2 AI404N 4 / / ~=4.oo /

(I3 50)/

x:o.77 / ~ / ~ - \ \

0 10 Si3N4

AI2 03 lOO

-90

-80

-70

0 -60

-so

-4O ~

-30

-20

-10

0 20 30 40 50 60 70 80 go 100

Equivalent % AI A]N

FIG. 7. The Si3N4-SiO2-A1203-A1N system, showing the compositions that can be prepared from any combination of the starting materials in the triangles defined by either Si3N4-A1N-SiO 2 or Si3N4-AIN-AI203 [17].

system is a reciprocal salt system, the compositions as marked in Fig. 7 can be produced using different combinations of different groups of starting materials [17], i.e., combinations of SisN 4, SiO 2, and A1N, or SisN4, A1N, and A1203. Boskovic's results showed that all compositions can be sintered to high density as shown in Fig. 8. The sintering of these mixtures can be described as transient liquid phase sintering. When mixtures are heated at high temperatures, liquid forms at the contact points of the particles, which aids densification. During the later stages of sintering, the liquid reacts with the starting material to form crystalline /3'-SiA1ON. When the reaction is completed, the liquid is exhausted. Compositions using mixtures of Si3N4, SiO2, and A1N had a higher weight loss and could only be sintered at high temperatures where large amounts of liquid were formed, sealing the open pores in the beginning of the sintering process, and thus preventing vaporization. This can be explained using the phase diagram shown in Fig. 9, which was published by Naik et al. [18]. In the initial stages of the reaction, A1N reacts with SiO 2 to form a liquid with a composition near the SiO 2 corner of the diagram. This liquid decomposes easily and vaporizes, resulting in a high weight loss and low density. For mixtures using SisN4, A1N and A1203 as starting materials, SisN 4 and

136 TSENG-YING TIEN

I001 w I i i i I I

" " g o

~ 8 o ol c-

13 7 0

o

c - 6 0

5 0

1 1 1 l l ~ l ~00 1603 1'800 2000

sintering temperature [*C]

FIG. 8. Sintering rate curves by Boskovic [16] for the composition, f l -Si6_xAlxN8_xO x (x = 4, or 60 equivalent percent of aluminum) in the Si3Na-SiO2-AI203-A1N system using different starting materials. A, Si2A14OaN4; O, Si3N4; r-l, SiO 2 + AIN; O, SiO 2 + AIN.

A120 3 react to form a liquid with low SiO 2 content, which is more stable. This produces a lower weight loss and results in a higher density material at lower temperatures.

4. Si3N4-BeO SYSTEM

The 1780~ isotherm of the Si3N4-SiO2-Be3N2-BeO system is given in Fig. 10. Similar to the Si3N4-SiO2-AIN-A120 3 system, solid solutions with the /3-Si3N 4 structure were found [19, 20]. These solid solutions were found to be stoichiometric, containing no lattice defects, and having a general formula , Si3+xBe2xO4xN 4. Similar to the Si3N4-SiO2-A1N-A1203 system, polytypoids are also found at the Be3N 2 corner of this diagram. Note that a two-phase region,/3-Si3+xBe2xO4xN 4 solid solution plus BeSiN 2, exists in the phase diagram, but there is no phase field in which the solid solutions are in equilibrium with a liquid at


SiO 3Al203.2Si02 AI203 . . . . . , ' ' ' J - ' ~ 100

~ , , , " ~ Al,e/3,.13) O~_.N. " - ~ - . . / / ~ / j , / / / ? > . . I / / //11 1

~ ' , / / " / , , 80

Si,ON

. . . . . . . ~ 0 o ~o 20 30 ~'o sb 60 70 80 90 ~oo

Si~ N~ Equivelent - % A[ AIN

FIG. 9. The 1750~ isotherm (18) for the Si3N4-SiO2-AI203-A1N system.

1760~ A composition in this two-phase field should have superior high- temperature mechanical properties. Greskovich [21] has synthesized ceramics in this phase field, and high-temperature stress rupture tests showed that they are the most stable silicon nitride ceramics among all of the other systems studied.

B. Silicon Nitride Systems with Two Metal Oxides

Subsolidus phase relationships in the Si, A1, Y / N , O system have been established by Naik and Tien [3] and by Sun et al. [4]. Sixty- eight compatibility tetrahedrons have been established in the Si3Na-SiOz-A1N-AlzO3-YN-Y20 3 system. Subsolidus phase relationships in part of the Si, A1, Mg/N, O system bounded by the compounds Si3N4, SiO2, MgzSiO4, MgAlzO4A1203, and SizA1404N have also been studied by Nunn et al. [22]. Twenty compatibility tetrahedrons have been established in this subsystem. A number of compositions in these compati-

138

S i 0 2

T S E N G - Y I N G T I E N

B e O

f' I ~ I

>,, x

o o I

I 4--, C 2 1 I (3 > Ii (:D cr II

LLI I.

S i 2 N 2 0

2 0 5 3 K

60 BO S i 3 N ~ 2 0 2 B e S i N 2 B e 3 N 2

E q u i v a l e n t % B e

F]o. 10. The S i3N4-SiO2-BeO-Be3N2-BeO system showing f l-Si3N 4 solid solution formation. Note that the polytypoid in the Be3N 2 corner of the diagram is similar to those in

}2-A1203-AIN [19, 20].

B e 3 N 2 - - - ~qu~vo,~tO~o ~ - - B e O

\\

~ B e S i N 2

\

, , j , i ~ i i , , .

~ 2 J _. f BeO.Ab03 I

/ // . " " L3AI2 03 'Si02 I II -- BeO. 3AI203

- ~ . . . . - - . , , K A:I 89 67 5:6 4:5 3:4 "[AIfsi3_Xll404_xN x

A t N - - Eq . . . . lent% 0 - . - A t 2 0 ,

FIG. 11. The S i3N4-SiO2-AI203-AIN-BeO-Be3N2 system. A wide solid solution range was observed in this system. The single-phase/3-Si3N 4 solid solution region is restricted to the metal-to-nonmetal ratio of 3 :4 plane. This indicates that the solid-solution is substitutional and no lattice defects exist in the structure [23].


bility tetrahedrons can be used to develop silicon nitride ceramics. These are discussed in later sections.

The Si3Na-SiOz-A1N-AlzO3-Be3Nz-BeO system has been studied by Gauckler [23] and is shown in Fig. 11. This figure shows an extensive single-phase solid solution region with a /3-Si3N 4 structure. These solid solutions are restricted to the plane having a metal-to-nonmetal ratio of 3:4. The metal-to-nonmetal ratio of 3:4 indicates that these solid solutions are substitutional and contain no lattice defects in the structure. Because of the toxicity of the Be, no further development work was performed after his dissertation. Therefore, this system is not discussed further.

1. Si3N4-SiO2-A1N-A1203-Mg3N2-MgO SYSTEM

Among all of the subsolidus compatibility relationships in this system, only the tetrahedra containing the compound cordierite (2MgO:2A1203: 5SiO 2) and the compound MgA12SinO6N 4 (MgO:A1203:2Si2N20) are of interest. Both of these phases are compatible with silicon nitride and are glass forming [24], thus making it possible to design silicon nitride ceramics with either one of these compounds as the grain-boundary phase. The thermal expansion coefficient of cordierite is 2.5 x 10-6/K and the thermal expansion coefficient of the MgAlzSi406N 4 phase is 2.31 x 10-6K. The thermal expansion coefficient of/3-Si3N 4 is 3.4 x 10-6/K, which is higher than both of these compounds. Silicon nitride ceramics with either cordierite or the compound M g A l z S i 4 0 6 N 4 as the second phase would contain internal stresses at the grain-boundaries due to the thermal expansion mismatch between the matrix and the grain-boundary phase. The effect of thermal expansion mismatch of the grain-boundary phase and the/3-Si3N 4 grains is discussed in a later section.

2. Si3N4-SiO2-A1N-A1203-YN-Y203 SYSTEM

Fig. 12 shows the four major solid phases in the system Si3N4-SiO2-A1N- A1203-YN-Y203, as reported by Sun [4]. These four silicon nitride phases are /3-Si3N4, a-SiA1ON, A1N-polytypoids, and silicon oxynitride. The morphology of these phases is different. The /3-Si3N 4 grains are elongated hexagonal rods; the a'-SiA1ON grains are equiaxed; the A1N- polytypoids are platelets; and the silicon oxynitride grains are equiaxed. Using combination of these phases, the following ceramics with different

140 TSENG-YING TIEN

YN Y303N Y203 H (Apatite)=Y 10(SiO4)6N2 K (Wollastonite)=YSiO2N ? , M (Melilite)=Y2Si303N4 / r J (Wohleri te)=Y4Si207N2 / r \ Jss=Y4Si207N2-Y4A1209 S.S. f \ YAM=Y4A1209 / / YAG=Y3AI5012 / \ ~ ~ 2SIO5

I t~-slalon pl ~.~2Si3N6 .

YN:aA1N L _ ~ ~ SiO2 ~ 1 3 N 4

3A1203.2SIO2 2H8 27R 2'iR 12H i'5R AI203:AIN

A1N A1203

FIG. 12. The subsolidus phase relationships in the Si3N4-SiO2-AI203-AIN-Y203-YN system, showing four major silicon nitride phases [4].

microstructures can be designed:

1. Acicular fl-Si3N 4 grains with grain-boundary phases. 2. Acicular ]3-Si3N 4 grains with equiaxed a-SiA1ON.

3. Acicular/3-Si3N 4 grains with polytypoid platelets.

4. Equiaxed a'-SiA1ON with polytypoid platelets.

5. Acicular/3-Si3N 4 grains with equiaxed silicon oxynitride.

The next section of this paper uses the Si3N4-SiO2-A1N-AI203- YN-Y203 system to demonstrated the alloy design.

IV. Alloy Design

Silicon nitride ceramics should have high flexural strength and fracture toughness at room temperature and good creep resistance at elevated temperatures. Because these properties are determined by the microstructure of these ceramics, further improvement of these ceramics is possible through alloy design. The following sections will be devoted to micro-


structure design through the use of the phase relationships of silicon nitride-metal oxide systems.

A. Acicular f l -S i 3 N4 Grains with Grain-Boundary Phases

1. GLASSY GRAIN-BOUNDARY PHASES

The common additives used for sintering silicon nitride ceramics are MgO, A120 3, and Y203 or combinations of these oxides. When mixtures of silicon nitride starting material, usually a-Si3N 4 powders with oxide sintering aids, were mixed, compacted, and sintered in a nitrogen atmosphere, the oxide additives and the oxide layer on the surface of the silicon nitride particles form a eutectic liquid at the start of the sintering process. The liquid aids densification of the compact according to the liquid phase sintering theory. A sintering shrinkage curve [25] is given in Fig. 13, which allows one to follow the sintering process. The first stage of sintering produces reactive liquids, and the densification process is by particle rearrangement. After the first stage of sintering, the a-Si3N 4 starting powder will start to dissolve in the reactive liquid and reprecipitate from the liquid as/3-Si3N 4.

With prolonged heating at sintering temperatures, /3-Si3N 4 grains, which are in equilibrium with the reactive liquid, will grow anisotropically follow the empirical grain growth law [ 26 ]:

O n _ D' d = KDt

-3 130 (0

- 6 t - . n r -

u) - 9

O) t - "i -12

-15

!~atina rate' 3~

SIAION-YAG

SiAION-Cord.

0 300 600 900 1200 1500 Temperature (~

FIG. 13. Sintering curves for compositions in the SiAION-Y3AlsO12 system and the SiA1ON-cordierite system ( 2 4 ) . The heating rate was 3~ per minute [25] .

142 TSENG-YING TIEN

The rate constant K and growth exponent n are different for the length and width directions of the /3-Si3N 4 grains. During prolonged heating, the /3-Si3N 4 grains develop into an acicular morphology and the aspect ratio increases with time. Lai and Tien [26] show that the time exponent n equals 3 in the length direction and 5 in the width direction. Activation energies 686 and 772 J / m o l for length and width directions, respectively, when the compacts were sintered under a nitrogen pressure of 10 atm.

The aspect ratio is defined as:

A R = W = K ~ 5 t(1 - K ~ 5 t

/ ( 1 / 3 ) ~ L "~'0L 2/15

= i,.1/5 t exp ~0W

Qw t 5

RT

where L is the average length of the /3-Si3N 4 grains, W is the average length of the /3-Si3N 4 grains, K c and K w are the rate constant in the length and width direction of the /3-Si3N 4 grains, Qc and Qw are the activation energies, R is the gas constant, and T is temperature. Fig. 14 shows our results for grain-growth studies [26]. Fig. 14c shows the change in aspect ratio as a function of sintering time. All of the aspect ratio vs. time curves are parallel straight lines with a slope equal to 2/15. This slope is equivalent to the difference in reciprocal growth exponents in both directions, 1 /3 - 1 /5 = 2/15.

Grain-growth studies have also been carried out under 1 atm of nitrogen pressure [25]. The results were found to be different from those studies performed at 10 atm. The time exponent n for both directions was equal to 3. This will result in an equation in which the exponent of the time term is equal to one, and the aspect ratio is time independent during sintering.

Lai and Tien [26] (Fig. 15) have shown that the flexure strength and fracture toughness of silicon nitride ceramics depends on the aspect ratio

FIG. 14. Isothermal growth of fl-Si3N 4 grains in a composition containing 90 wt % Si3N 4 and 10 wt % Y3AI5012 as the sintering aid. Sintering at 10 atm pressure of nitrogen [ 26 ]. O, 1900~ o, 1800%; t2,1700~ II, 1600~ (a) Growth rate in the length direction. The growth exponent n equals 3 in the

empirical equation D n - D~ = Kt.

(b) Growth rate in the width direction. The growth exponent equals 5. (c) The aspect ratio of the acicular fl-Si3N 4 grains in specimens sintered at

different temperatures and time.

4 S T U D Y O F S I L I C O N N I T R I D E C E R A M I C S

100

10 v

1

.1 .1

10

v

. , . . .~

.1 .1

...-..--.o

r

I ~ - . ~

m-'-"

. . . . . . . !

1 10

e . . ' -"

. . . . . . . !

1 10

100

~ 10

O y = 8 .8462 * x A 0 . 1 M 2 0 RA2 - 0 .986

�9 y - 7 .0926 * XA0.16615 RA2 -- 0 . 9 8 0

D Y -- $..q403 * x A 0 . 1 6 9 8 8 RA2 = 0 . 9 8 8

�9 y = 4 . 4 4 4 9 * x A 0 . 1 ~ t 9 2 RA2 = 0 .963

C

. . . . . . i . . . . . .

1 I 0

Time (hr)

143

1100

13.

t -

r

( 9

03 . . . , .

L - 3 X ( 9

LL

1000

�9 i

I 9OO

800

700

600

i I

500 �9 ' 7 8

144

l . i . i .

9 10 11

Aspect RaUo

T S E N G - Y I N G TIEN

12

11

10

9

E

a. 8

0 m

7

I �9 i �9 i - n -

5 - I i I ~ I , L J

7 8 9 10 11 12

Aspect Ratio

FIG. 15. Microstructural dependence of flexural strength and fracture toughness of silicon nitride ceramics. The composition of these specimens was 90 wt % Si3N4-10 wt % Y3A15012 (27).

.<

0 Ln 0 t- O.

tn D 0

Q. L 0 E 0

10-

$

a �9 t r i v l e p o i n t s

�9 g r o i n k o u n O a r i e =

0 " �9 , w

7 8 9

j l im

/ i _.i__.__= .,

I I I I I

10 11 12 13 14

lU. 1.0

0.9 E O l...

0 8 -

0 .7 t/}

c 0.6 L.

0.5 7

b . F; o st. dev.

I I I I

,, + ,o ,', ,2 ,'3 ,',+

E q . % 0 2 - ( AI 3 ~ 11 E q . % }

=_.. - 0 . 4 o

D C3. Q

- 0 . 3 ca.

0.2 <

0.1 ~

I1-11 1673 K

E 0 B st ra in

FIG. 16. Greil and Weiss (28) studied the effect of glass content at the grain boundaries on the mechanical properties of silicon nitride ceramics. Compositions with different oxygen contents in the Si3Na-SiO2-AI203-AIN system were prepared. Results show that the composition containing the higher glass content deforms less under load at high temperatures. As shown in (b), the /3-Si3N 4 grains in the specimens have a higher glass content and a higher aspect ratio. The shape factor in (b) was defined as F = (4pA)/U 2, where A is the cross-sectional area and U is the circumference of the grain. The lower the shape factor, the higher the aspect ratio. When the shape factor equals one, the grains are equiaxed.

146 TSENG-YING TIEN

FIG. 17. Microstructure of silicon nitride ceramics sintered at 1900~ for 2 h under 10 atm of nitrogen (27).

of the j3-Si3N 4 grains [27]. Greil and Weiss [28] have shown similar results. Greil and Weiss formed /3-Si3N4 solid solution compositions containing various amounts of liquid phase in the Si, A1 /O,N system. Their results indicated that the compositions containing higher amounts of liquid had a smaller deformation rate under load at high temperatures than the specimens containing lower amounts of a glassy grain-boundary phase. The Greil et al. results are shown in Fig. 16 At the time, Greil suggested that the differences in deformation at high temperatures was due to the differences in viscosity of the grain-boundary glass. It was speculated that when the impurity contents in each composition are the same, the larger amount of glass would have a lower impurity content, and hence a higher SiO 2 content and, therefore, a higher viscosity. On closer observation, as shown in Fig. 16b, it was found that the /3-Si3N 4 solid solution grains grown in the presence of a larger amount of liquid showed a higher aspect ratio. These results led the development of silicon nitride ceramics in the direction of a microstructure having long interlocking /3-Si3N 4 grains, as shown in Fig. 17.


2. CRYSTALLINE GRAIN-BOUNDARY PHASES

The solid-liquid reactions in the /3-SiA1ON-Y3A15012 system (yt t r ium-aluminum-garnet-YAG) have been determined by Wisnudel [29] and the diagram is shown in Fig. 18. /3-SiA1ON with YAG as the sintering aid have been prepared by Hohnke and Tien [30]. These compositions after densification contain a crystalline/3-Si3N 4 solid solution and a glassy grain-boundary phase. However, the YAG crystallizes at the grain boundary on annealing.

Chen [31] has shown that the creep resistance increased after the grain-boundary glassy phase was crystallized. The results are shown in Fig. 19. The three curves in this figure each represents a different treatment. The first specimen was hot-pressed at 1750~ for 60 min under a pressure of 20 MPa. This specimen failed after 5 h under a load of 133 MPa at 1170~ The second specimen was hot-pressed under the same conditions as the first, but it was annealed at 1250~ for 50 h, allowing the garnet to crystallize at the grain boundaries. This sample failed after 150 h of creep testing. A third sample was hot-pressed under the same conditions but for 150 min and then annealed. This specimen did not fail after 150 h under load. This specimen has larger fi-Si3N 4 grains. These results suggest that the creep resistence of silicon nitride ceramics depends on their microstructure and the grain-boundary phase. The larger the/3-Si3N 4 grains, the slower the creep rate. Crystalline materials at the grain boundary resist creep at high temperatures.

YAG

Si 3 N 4 wt. %AIN :AI 2 O3"=1~ ~ 60

FIG. 18. Liquidus curves of the SiA1ON-Y3AI5012 system (29)._._,1550~ isotherm; __.., 1650~ i so therm;~ , 1750~ isotherm.

148 TSENG-YING TIEN

500.0

400.0

o o o o e 3 0 0 . 0 - , o - ' *

'-! gO v ooO, o c ..=,4,jo.oOO j

x ,,oO- _....~- ~, 200.0 .. .- ee ~ a ~

0 . 0 - I n i 0.0 30.0 60.0 90.0

1310-7G SIAl ON ,~,.- 1170 ~ 133 MPa ,..,r"

4r" i r . 4 r -B"

o,..m

i , i

120.0 150.0 180.0 Creep time (hour)

FIG. 19. Creep behavior of silicon nitride ceramics. The composition of this ceramic contains 7 vol % Y3A15012 as a sintering aid. The specimens were hot-pressed at 1750~ under a pressure of 20 Mpa. The holding times are marked on the curve. This curve indicates that crystallization of the grain-boundary phase improves the creep resistance. Garnet crystals were observed at the grain boundaries after annealing [31]. e, Sample A60; e, Sample B60; A, Sample B150, x, failure.

Fig. 20 relates the room-temperature fracture toughness to composition and state of the second phase (crystalline or amorphous) as shown by Bonnell [32] in the/3-SiA1ON-garnet and/3-SiA1ON-cordierite systems. A decrease in toughness with increasing volume of the grain-boundary glass was seen in the hot-pressed samples of both systems. More dramatic differences were observed between the amorphous and crystallized samples. After cordierite crystallized at the grain boundaries of the/3-SiA1ON ceramics, the toughness of the specimen increased, whereas the toughness decreased when the garnet crystallized. This phenomenon illustrates that the thermal expansion mismatch affects the mechanical properties of the silicon nitride ceramics. Crystalline garnet has a thermal expansion coefficient of 8 • 10 -6, which is higher than that of silicon nitride (3 • 10-6). While in the/3-SiA1ON-cordierite system, the grain-boundary crystalline phase has a thermal expansion coefficient of 2 • 10 -6, which is lower than that of silicon nitride. Hence the grain-boundary phase will be under tension in the case of the /3-SiA1ON-garnet system but will be under compression in the/3-SiA1ON-cordierite system.


Od

lel

E Z :E

u_ x 5

4 �9 �9

T T h x n n e o l e d ~

r - , " - - , , e ~ Cor~; .

; ,b ,? 2'0 .... 25 VOLUME PERCENT ADDITIVE

FIG. 20. R o o m - t e m p e r a t u r e f rac ture toughness as a funct ion of vo lume fract ion of the

second phase for S iAION composi t ions conta in ing garne t and cordier i te . The specimens were

ho t -pressed [ 32 ].

1000

800

~ L

600 ._m

~ 400

~., 200

I ---.- : RBSN-A (as-received) -..r : RBSN-A (1900~

O : Infil., 1600~ (RBSN-A,1600~ �9 ..a.... : Infil., 1600~176 �9 ,,o- : Infil., 1600~ cry. (RBSN-A, 1900~ 2H) -.o- : Infil., 1700~ (RBSN-A, 1900~ 1H �9 --o- : Infil., 1700~ cry. (RBSN-A, 1900eL, IH) ,,,o-- : Infil., 1800~ (RBSN-A,1800~ 2H)

b

0 400 800 1200 1600

Temperature (~

FIG. 21. F lexura l s t rength of R B S N infi l t rated with equi l ibr ium liquid. The liquid has the

eutect ic compos i t ion on the S i 3 N 4 - Y 2 S i 2 0 7 join sa tu ra t ed with silicon ni t r ide crystals [33].

150 TSENG-YING TIEN

3. INFILTRATION

In the S i 3 N 4 - S i O 2 - Y N - Y 2 0 3 system, on the Si3N4-Y2Si20 7 join the eutectic forms at 1500~ as shown in Fig. 6. Any composition on the join with higher Si3N 4 content then the eutectic liquid will have two phases, ~ - S i 3 N 4 and Y28i207, which coexist under equilibrium conditions. Based on this fact, the following experiment was carried out by Sheu [33]. In a BN crucible, a high-porosity reaction-bonded silicon nitride (RBSN) was immersed in a liquid of the eutectic composition supersaturated with /3-Si3N 4. The crucible was kept in a furnace for a period of a few hours

YN YSOSN Y 203 tt (Apadte)=Y 10(SiO4RiN2 /r - K (Wotlastonite)= YSiO2N / M (Melilite)=Y2Si303N4 / \ J (Wohlerite)=Y4Si207 N2 / Jss=Y4Si2OTN2-Y4AI209 S.S. / \ YAM=Y4AI209 / YAG=Y3AI5OI2 / ~\\ SiO5

/ a'\ Si3N4 o'

/ ~ - 3~2o3.2sio2 21'J6 2)R 2'iR i-2H I'SR

A1N A1203

AIN

ot'-sialon plane

YN:3AIN ~ . _ . . ~ 0t'

" " " ~" " . ~ i 3 N 4

o.p l . . ~ .s t j .,. ~ts~'~ 0

~ ~ ~ ....

A1203 :AIN

FIG. 22. Subsolidus phase relations in the Si, AI, Y / N , O system showing a series of compatibility tetrahedrons formed by a'-sialon,/3'-SiAION, and A1N polytypoids [34].

4 STUDY OF SILICON N I T R I D E CERAMICS 151

until the open pores of the RBSN were saturated with the liquid in equilibrium with / 3 - S i 3 N 4 [11]. In principle, the liquid will not react with the RBSN skeleton. Fig 21 shows the strength change after the infiltration. The strength dropped at high temperatures. These data indicate that the liquid had corroded the contact points of the / 3 - S i 3 N 4 grains in the RBSN.

B. Two-Phase Composites

Sun [34] has shown that / 3 -S i3N 4 solid solutions with a'-SiA1ON and AlN polytypoids form a series of compatibility triangles as shown in Fig. 22. Using different combinations of these phases, a series of ceramics with different microstructures can be designed. All of the melting points of these phases are very high. Compositions in these compatibility tetrahedrons do not form a liquid at very high temperatures. Ceramics with compositions in these tetrahedrons are expected to have superior high- temperature mechanical properties. Using individual oxides and nitrides as starting materials, high-density ceramics should be obtainable through transient liquid-phase sintering.

YN Y203

\\

�9 / "~2Y > t . . . . . . ~ . i " $ i 3 N 4 o'__ _ J __ .7--- - io:

3A1203.2SIO2

2H8 27R ~ A1203 :AIN A1N A1203

FIG. 23. The Si3N4-SiO2-AI203-AIN-Y203-YN system showing that a'-SiA1ON compositions can be formed on the Si3Na-(Y203:9AIN line intersecting the line between the triangles Si3N4-A1N-Y203 and Si3Na-(AleO3:AIN)-(Y203:3A1N) [35].

152 TSENG-YING TIEN

1400

1200

1000 e L

800

~ 600

~ 400

E 200

, , l l l . . l . . . l , , - * . , . i * . . l i l l l l , l

3

[] �9 30% o~', 1780~ �9 �9 30% a', 1780~ Lx �9 30% cx',1800~ x �9 30% o(, 1900~ �9 �9 40% o(, 17800C,0.5H a �9 40%a' , 1780~ 1H �9 " , i , , - i , , , i - - - i - , - i - , , i - , , i , , ,

0 400 800 1200 1600

Temperature (~

F~G. 24. Flexural strength of hot-pressed specimens containing two phases, a ' - S i A I O N plus flr-Si3N4. The numbers indicating volume percent of the a ' - S i A I O N are given on the line. The hot-pressing temperatures and times are given in the figure [36].

1. 3 - S i 3 N 4 GRAINS WITH E Q U I A X E D a ' - S i A 1 O N

As shown in Fig. 23, the line connecting the compound ~-Si3N 4 and the point Y203:9A1N intersect the single-phase solid solution region a'-SiA1ON as shown by Huang et al. [35]. Compositions on this line between the /3-Si3N 4 and a'-SiA1ON contain two phases at subsolidus temperatures. Since the melting points of both of these two compositions are very high, ceramics containing /3-Si3N 4 and a'-SiA1ON will not become soft until they are subjected to very high temperatures. Specimens containing these two phases were hot-pressed. The composition containing 70% /3-Si3N 4 and 30% a'-SiA1ON showed a flexural strength of 1200 MPa at room temperature and 800 MPa at 1400~ as shown in Fig. 24 [36].

2. a ' - S i A I O N WITH POLYTYPOID PLATELETS

As shown in Fig. 22, a'-SiA1ON is compatible with all of the A1N polytypoids. Ceramics with equiaxed a'-SiA1ON and polytypoid platelets should have high toughness [37]. Since a'-SiA1ON has a high hardness, it is anticipated that such ceramics would have superior mechanical properties.


YN

\

\

\

a'-sialon plan~

YN:3A1N/~. ~.

A1N

Y203

Si207

YAG ~ .

~ $102

3A1203.2SIO2

A1203

FIG. 25. The Si3N4-SiO2-AI203-AIN-Y203-YN system, showing that silicon oxynitride and a'-SiA1ON are not compatible phases [39].

3. fl-Si3N 4 GRAINS WITH POLYTYPOID PLATELETS

High-density ceramics containing /3-Si3N 4 grains with polytypoid platelets are very difficult to synthesize because the liquid-forming temperature in this part of the system is very high. Scientists at the Shanghai Institute of Ceramics [38] had synthesized two-phase ceramics consisting of/3-Si3N 4 grains and polytypoid platelets. La20 3 was used as a sintering aid. Their results showed a modest increase in toughness. However, because of the sintering aid used, a glassy grain-boundary phase was observed.

4. fl-Si3N 4 GRAINS WITH SILICON OXYNITRIDE

Trigg and Jack [39] prepared ceramics containing oxynitride and 13 -Si3N 4. As described before, the oxynitride has equiaxed grains and the /3-Si3N 4 grains are elongated hexagonal rods. It is anticipated that these microstructures should make ceramics with good mechanical properties.

Fig. 25 shows a compatibility relationships between oxynitride and /3-Si3N 4. Note that the tetrahedron /3-Si3N4-fl-SiA1ON-YzSi20 7- Y3A15O12 separates the silicon oxynitride from a'-SiA1ON and A1N

154 TSENG-YING TIEN

polytypoids. This indicates that ceramics with silicon oxynitride and a'-SiA1ON and A1N polytypoids cannot be formed.

V. Conclusion

Silicon nitride ceramics are one of the most promising materials for structural applications. The properties of the silicon nitride ceramics depend on the phases present and their microstructures. This paper describes how to approach microstructural design of these materials by using the phase equilibrium diagrams of silicon nitride-metal oxide systems.

There are four major silicon nitride phases in this system:/3-Si3N4, a'-SiA1ON, A1N-polytypoids, and silicon oxynitride. The morphology of each of these phases is different. The morphology of the/3-Si3N 4 grains is that of elongated hexagonal rods; the a'-SiA1ON grains are equiaxed; the A1N-polytypoids are platelets; and the silicon oxynitride grains are equiaxed. Using a combination of these phases, the following ceramics with different microstructures can be designed:

1. Acicular/~-Si3N 4 grains with grain-boundary phases.

2. Acicular/3-Si3N4 grains with equiaxed a'-SiA1ON.

3. Acicular/3-Si3N 4 grains with polytypoid platelets.

4. Equiaxed a'-SiA1ON with polytypoid platelets.

5. Acicular/3-Si3N4 grains with equiaxed silicon oxynitride.

The Si3N4-SiO2-A1N-A1203-YN-Y203 system has been used to demon- strate the principle of alloy design.

Acknowledgment

I would like to thank the U. S. Department of Energy and the National Science Foundation for their financial support. I would also like to thank Dr. Paul C. Becker, formally of Allied-Signal Corporation, for his stimulating discussions and his critical review of this manuscript.

4 S T U D Y O F S I L I C O N N I T R I D E C E R A M I C S 155

References

1. L. J. Gauckler and G. Petzow, Representation of multi-component silicon nitride based systems. In "Progress in Nitrogen Ceramics" (F. L. Riley, ed.), pp. 41-62. Voordhoff Int., Gr6ningen, Leyden, The Netherlands, 1977.

2. L. J. Gauckler, H. Lukas, and G. Petzow, Contribution to the phase diagram Si3N4-SiOz-A1N-AI20 3. J. Am. Ceram. Soc. 58 (7-8), 346-347 (1975).

3. I. K. Naik and T. Y. Tien, Subsolidus phase relations in part of the system Si, AI, Y/N, O. J. Am. Ceram. Soc. 62, 642 (1979).

4. W.-Y. Sun, T.-Y. Tien, and T.-S. Yen, Sub-solidus phase relationships in part of the system Si, A1, Y/N, O: The system Si3N4-AIN-YN-A1203-Y20 3. J. Am. Ceram. Soc. 74 (11), 2753-58 (1991).

5. C. C. Sorrell, Silicon nitride and related nitrogen ceramics. J. Aust. Ceram. Soc. 18(2), 22-34 (1982); 19(2), 48-67, 68 (1983).

6. G. G. Deeley, J. M. Herbert, and N. C. Moore, Dense silicon nitride. Powder Metall. 8, 145 (1961).

7. T. Y. Tien, G. Petzow, L. J. Gauckler, and J. Weiss, Phase equilibrium studies in Si3N4-metal oxide systems. In "Progress in Nitrogen Ceramics" (F. L. Riley, ed.), pp. 89-99. Martinus Nijhoff, Boston, 1983.

8. R. Miiller, E. E. Hucke, T. Y. Tien, and G. Petzow, The system Si3N4-SiOz-Mg3Nz-MgO. Am. Ceram. Soc. Bull. 58(9), 885 (1979).

9. K. H. Jack, The fabrication of dense nitrogen ceramics. In "Processing of Crystalline Ceramics" (H. Palmour, III, R. F. Davis, and T. M. Hare, eds.), pp. 561-578. Plenum, New York and London, 1976.

10. Y. Inomata, Y. Hasegawa, and T. Matsuyama, Reaction between Si3N 4 and MgO added as a hot-pressing aid. Yogyo Kyokaishi 85(1), 29-31 (1977).

11. F. F. Lange, Phase relations in the system Si3N4-SiOz-MgO and their interrelation with strength and oxidation. J. Am. Ceram. Soc. 61 (1-2), 53-56 (1978).

12. L. J. Gauckler, H. Hohnke, and T. Y. Tien, The system Si3N4-SiO 2 - Y203 . J. Am. Ceram. Soc. 74 (1-2), 35-37 (1980).

13. Y. Oyama and O. Kamigaito, Solid solubility of some oxide in Si3N 4. Jpn. J. Appl. Phys. 10, 1637 (1971).

14. K. H. Jack and W. I. Wilson, Ceramics based on the Si-AI-N-O and related systems. Nature (London), Phys. Sci. 238(80), 28-29 (1972).

15. D. R. Clarke, The microstructure of nitrogen ceramics. In "Progress in Nitrogen Ceramics" (F. L. Riley, ed.), pp. 341-358. Martinus Nijhoff, Boston, 1983.

16. S. Boskovic, L. J. Gauckler, G. Petzow, and T. Y. Tien, Reaction sintering forming b-Si3N 4 solid solutions in the system Si, A I / N , O Powder Metall. Int. 9 (4), 185-89 (1977); 10(4), 184-85 (1978); 11 (4), 169-171 (1979).

17. G. Petzow, L. J. Gauckler, T. Y. Tien, and S. Boskovic, b-Si3N 4 solid solutions formation by reaction sintering in the system Si, AI/N, O. In "Factors in Densification and Sintering of Oxide and Non-Oxide Ceramics" (S. Somiya and S. Saito, eds.), pp. 28-39. Gakujutsu Bunken Fukyu-Kai, Tokyo, 1979.

18. I. K. Naik, L. J. Gauckler, and T. Y. Tien, Solid-liquid Equilibrium in the system Si3Na-SiOz-AIN-AI20 3. J. Am. Ceram. Soc. 61(1-2), 332-335 (1978).

19. I. C. Huseby, H. L. Lukas, and G. Petzow, Phase equilibrium in the system Si3N4-SiO2-Be3N2-BeO. J. Am. Ceram. Soc. 58(1-2), 377 (1975).

20. D. P. Thompson and L. J. Gauckler, Further study of the Be-Si-O-N polytypes. J. Am. Ceram. Soc. 60 (9-10), 470-471 (1977).

156 T S E N G - Y I N G T I E N

21. C. Greskovich, W. D. Pasco, and G. D. Quinn, Thermomechanical properties of a new composition of sintered SiaN 4. Am. Ceram. Soc. Bull. 63(9), 1165-1170 (1984).

22. S. D. Nunn, H. Hohnke, L. J. Gauckler, sand T. Y. Tien, Subsolidus phase relationships in part of the system Si, A1, Mg/N, O. Am. Ceram Soc. Bull. 57, 321 (1978).

23. L. J. Gauckler, Phase equilibrium studies in the system Si, A1/N, O and Si, Al, Be/N, O. Doctoral Dissertation, University of Stuttgart (1975).

24. S. D. Nunn, Processing and properties of SiC whisker reinforced SiaN 4 ceramic matrix composite. Doctoral Dissertation, University of Michigan, Ann Arbor, (1991).

25. C. M. Hwang, The system SiAION-Y3AI5OI2 and SiA1ON-cordierite: Sintering and grain growth, Doctoral Dissertation, University of Michigan, Ann Arbor (1988).

26. K. R. Lai and T. Y. Tien, Kinetics of ~-SiaN 4 grain growth in silicon nitride ceramics sintered under high nitrogen pressure. J. Am. Ceram. Soc. (in press).

27. K. R. Lai and T. Y. Tien, "Optimization of Silicon Nitride Ceramics," Bimon. Rep. Ceramic Technology for Advanced Heat Engine Project, Oak Ridge National Labora- tory, Oack Ridge, TN, 1991.

28. P. Greil, and J. Weiss, Evaluation of the Microstructure of b-SiAION solid-solution materials containing different amount of amorphous grain boundary phase. J. Mat. Sci. 17 (6), 1571-1578 (1992).

29. M. Wisnudel, Solid-liquid reaction in the system SiAION-YaAIsOI2. MS Thesis, Univer- sity of Michigan, Ann Arbor (1991).

30. H. Hohnke and T. Y. Tien, Solid-liquid reactions in part of the system Si, Al, Y/N, O. IN "Progress in Nitrogen Ceramics"(F. L. Riley, ed.) pp. 101-110. Martinus Nijhoff, Boston, 1983.

31. C. F. Chert and T. Y. Tien, High temperature mechanical properties of SiAION ceramics: Microstructural effect. Ceram. Eng. and Sci Proc. 8 (7-8), 778-795 (1987).

32. D. A. Bonnell, T.-Y. Tien and M. Ruble, Controlled crystallization of the amorphous phase in silicon nitride ceramics. J. Am. Ceram. Soc. 70 (7), 460-465 (1987).

33. T. S. Sheu and T. Y. Tien, Infiltration of RBSN with equilibrium liquid in the system Si, Y/N, O. To be published.

34. W.-Y. Sun, T.-Y. Tien, and T.-S. Yen, Solubility limits of a-SiA1ON solid solutions in the system Si, Al, Y/N, O. J. Am. Ceram. Soc. 74 (10), 2547-2550 (1991).

35. Z.-K. Huang, T.-Y. Tien, and T.-S. Yen, Subsolidus phase relationships in Si3N4-AIN- rare-earth oxide systems. J. Am. Ceram. Soc. 69 (10), C241-242 (1986).

36. T. S. Sheu and T. Y. Tien, "Two-Phase Ceramics in the System Si, AI, Y /N ,O: I. fl-SiaN 4 and a'-SiAION," Bimon. Rep. Ceramic Technology for Advanced Heat Engine Project, Oak Ridge National Laboratory, Oak Ridge, TN, 1991.

37. I. Solomon and T. Y. Tien, "Two-Phase Ceramics in the System Si, AI, Y / N , O : II. a'-SiAION and AIN-Polytypoid," Bimon. Rep. Ceramic Technology for Advanced Heat Engine Project. Oak Ridge National Laboratory, Oak Ridge, TN, 1991.

38. H. R. Zhung, W. L. Li, J. W. Feng, Z. K. Huang, and D. S. Yan, SiaN4-AIN polytypoid composites by GPS. Eur. Ceram. Soc. (submitted for publication).

39. M. B. Trigg and K. H. Jack, Silicon Oxynitride and O'-SiAION ceramics. In "Ceramic Components for Engines" (S. Somiya, E. Kanai, and K. Ando, eds.), pp. 343-349. KTK Scientific Publishers, Tokyo, 1984.

The Use of Phase Studies M the Development of Whiskers and Whisker-Reinforced Ceramics

A L E K S A N D E R J. P Y Z I K A N D A L A N M. H A R T

Dow Chemical Company Central Research and Development

Advanced Ceramics Laboratory Midland, Michigan 48674

I. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 157 II. Whiskers and Ceramic Matrix Whisker-Reinforced Composites . . . . . . . . . . . 160

A. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 160 B. Silicon Carbide (SIC) Whiskers . . . . . . . . . . . . . . . . . . . . . . . . . . . 161 C. Silicon Nitride (Si3N 4) Whiskers . . . . . . . . . . . . . . . . . . . . . . . . . . 173 D. AI203-SiC Whisker Composites . . . . . . . . . . . . . . . . . . . . . . . . . . 181 E. Si3N4-SiC Whisker Composites . . . . . . . . . . . . . . . . . . . . . . . . . . . 187

III. Ceramics Toughened by in Situ Synthesized Whiskers . . . . . . . . . . . . . . . . 195 A. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 195 B. Si3N4-SiC Composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 195

IV. Self-Reinforced~Ceramics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 199 A. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 199 B. Silicon Nitride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 200 C. A1N-SiC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 217 D. Directly Solidified Eutectic Ceramics . . . . . . . . . . . . . . . . . . . . . . . . 218

V. Conclusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 221 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 222

I. Introduction

R e c e n t p r o g r e s s in p r o c e s s i n g a n d m a c h i n i n g o f c e r a m i c m a t e r i a l s has

c r e a t e d t e c h n o l o g i c a l b a s e s for m a n u f a c t u r i n g c e r a m i c c o m p o n e n t s wi th

i m p r o v e d f r a c t u r e s t r e n g t h s a n d i n c r e a s e d W e i b u l l m o d u l i . D e s p i t e t h e s e

a d v a n c e s , t he use o f l o w - t o u g h n e s s c e r a m i c m a t e r i a l s for s e v e r e - e n v i r o n -

m e n t s t r u c t u r a l a p p l i c a t i o n s is still l i m i t e d by t h e c e r a m i c ' s su scep t ib i l i t y

to a s u b s t a n t i a l s t r e n g t h loss c a u s e d by d a m a g e i n t r o d u c e d d u r i n g service .



158 ALEKSANDER J. PYZIK AND ALAN M. HART

A solution was found in the improvement of the fracture toughness, which leads to materials with higher reliability, more useful strength, and longer service lifetimes.

One of the successful approaches proven to increase the fracture toughness of ceramics is the application of single-crystal whiskers, introduced into the material or grown in situ during the densification stage. The chemically pure nature and highly ordered structure of ceramic whiskers can provide very high strengths, approaching 4 million psi [78].

The design of microstructures in "tough ceramics" is based on the understanding of the contributions coming from the main toughening mechanisms such as debonding, crack bridging, pullout and crack deflection [24]. The theoretical and experimental data show the importance of selecting a matrix material and whisker reinforcement that have the proper combination of elastic and thermal properties [4]. Toughness increases with increases (i) in the ratio of matrix to whisker Young's moduli, (ii) in the ratio of matrix to interface fracture energies, and (iii) of the whisker strength, diameter, and volume fraction [5]. When cracks

FIG. 1. Typical crack pattern in whisker-reinforced ceramics.

5 WHISKERS AND WHISKER-REINFORCED CERAMICS 159

propagate throughout the material, as shown in Fig. 1, the separation along whisker-matrix interfaces is desirable, independent of the predominant toughening mechanism. The debond length and size of the debond zone depend on several variables, such as stress applied, interfacial bonding, interface morphology, and whisker strength and diameter. The morphology and surfaces of ceramic whiskers can be tailored by controlling phase equilibria during whisker manufacture. Properties of the matrix-reinforcement interfaces, which are dependent on the interracial surface chemistry, can be changed by varying the chemical composition of starting materials and by controlling the phase equilibria established during the processing of composites. The nature of interactions between matrix and reinforcement is rather complex. Typically, at some processing stage, they can be related to the chemistry of the single component or the chemistry of the component systems, as illustrated schematically in Fig. 2. Although several excellent reviews analyze the effect of whiskers additions on the mechanical behavior of ceramic materials [5, 24], there is relatively less emphasis on the optimization of mechanical properties through controlling the chemistry of the reinforced ceramic systems.

This review focuses on the characterization of the whisker-matrix chemical interactions and on the relationships between the chemistry of the ceramic components, their grain morphologies, and their fracture

Chemistry of Reinforcing

Phase

DBfUlckts r s2Uor~thC:ss I Le:gteht &

Whisker Strength

Chemistry of Reinforcement-Matrix

Interface

Phase Compatibility

Interfacial Interface Bonding Morphology

Debond Length

FIG. 2. Parameters controlling the whisker debond length.


toughnesses. The phase relationships are discussed in terms of the three main toughening approaches:

1. Toughening of the ceramic matrix (phase A) by ceramic whiskers (phase B) formed in a separate process:

A + B , A B (1)

2. Toughening of the ceramic matrix (phase A) by whiskers or whisker- like grains (phase B) produced in situ during densification:

A + (B1, B2, B3) , A + B ) A B (2)

3. Toughening of the ceramic matrix (A) by controlling the micromor- phology of the matrix itself (self-reinforced ceramics). The process requires that an initial composition (A) form a liquid or solid solution (A 1) from which phases with different morphologies than those existing in A and A 1 are grown:

A ,A~ ,A2 (3)

II. Whiskers and Ceramic Matrix Whisker-Reinforced Composites

A. Introduction

The toughening of a ceramic by the addition of whiskers that are different from the matrix material is the most traditional and the best explored methodology. The advantage of this method is the broad spec- trum of ceramics that can be prepared in whisker form. Consequently, there is the potential for a broad range of ceramic matrix-ceramic whisker combinations. Since the whiskers are manufactured in a separate process, it is possible to control their size, morphology, and surface chemistry without affecting the matrix material. Additionally, whiskers can be chemically or thermally treated to change their future interactions with the matrix.

The disadvantage of this method is the high manufacturing cost of whiskers and the difficulties associated with processing them into homogeneous composites, e.g., mixing and densification. As a result, the whisker volume present in the composites is practically limited to 30 to 40 vol % when pressure-assisted densification techniques are employed and to be-


low 15% when pressureless sintering is applied. Despite these difficulties a remarkable two- to threefold increase in fracture resistance has been reported [44, 110].

Despite the fact that many different ceramic whiskers have been produced in research laboratories [93] only SiC and Si3N 4 have gained real commercial significance so far. Among whisker-reinforced ceramics the best results were obtained with AI203-SiC whiskers and Si3N4-SiC whiskers.

B. Silicon Carbide (SIC) Whiskers

1. PHASE EQUILIBRIA IN THE Si-C SYSTEM

The formation of silicon carbide from silicon and carbon is a well-known reaction and can occur at temperatures well below the melting point of

W T % C

5 15 25 30 50 3000 I I , -'1 I , , I , , " ] I I I

, "2830 + 40,, " [1.21 L

, , ' " 19 ,, 0 . . . . . . . [3] s S /

s e /

2600 - - ," t ~ L + C �9 _ _ 2540 + 40 .,,.,

0

z 2200 m

LLI 132 :=) I-- <: rr w 13_

1800 w I--

1400

70 90 I I I

L ~ 2 7

L + SiC

1402 + 5

SiC + C

SiC + C

~ooo I I I I I I I I I Si 20 40 50 60 80 C

A T % C

FIG. 3. Si-C phase diagram. After Elliot [22]. Reproduced with permission of McGraw- Hill Inc., New York.


0.5922

.=_ I - ~-- 0 .5924 O

0.5926

a t % C

0 4 10-3 10-2 10-1 . . . . . . . . I . . . . . . . . I , . . . . . . . . I / I / ' ' '

melt I [ melt+p-SiC

- liquidus . . . . . . I I (Scace & Slack' 1959) I ~- liquidus

- I / (Hall, 1958) liquidus . _ _ _ ] /

__%_:_ . . . . . . ,,,_

1 ~2_ measured ~ eutectic _ C in Si ~ solid solubility

I (Bean & Newman, 1971 )

- I C in Si + [J-SiC i

, , , , I . . . . , , , , I ~ ~ ~ , , ~ l ~ l , , , , . . . . I 4 " '

1017 1018 1019 102o

Carbon in a t o m s / c m 3

100 ! !!|

o bS_

FIG. 4. Si-C phase diagram for dilute carbon concentrations. After Voltmer and Padovi [126]. This paper was originally presented at the 1973 Spring Meeting of the Electrochemical Society held in Chicago, IL. Reproduced with permission of Electrochemical Society, Pennington.

silicon as demonstrated by van Konijnenburg [64]. A phase diagram for the silicon-carbon binary combined from other proposed diagrams by Elliot (22) is given in Fig. 3. The only condensed phases normally accepted to be present are silicon, silicon carbide, and graphite. Within the phase diagram, there is a eutectic point between silicon and silicon carbide at 1402~ and 0.75 atom % carbon. The liquidus curve between silicon and silicon carbide is shown up to 2600~ and 27 atom % carbon. A peritectic point is located at 2540~ and 27 atom % carbon under normal conditions. The effect of increased pressure of up to 300 atm would be to increase the peritectic point to about 3000~

The reaction to form silicon carbide whiskers, as described in a later section, will occur at or near the surface of carbon-containing materials. The dependence of the reaction and its kinetics on the concentration of carbon is critical. Figure 4 depicts the phase relationship of silicon and carbon at very dilute concentrations of carbon, which favor whisker growth [126]. The solubility of the carbon in silicon is greatly affected by the concentration of carbon present. For example, the solubility of carbon at 2000~ is estimated to be about 10 wt %.

2. PHASE EQUILIBRIA IN THE S i - C - O SYSTEM

Synthesis of most commercial silicon carbide whiskers is a result of the reaction of carbon with silicon dioxide so that the three-component


1.0

cf cO u_ 0.6 0 I'-- Z 53 o ,<

0.2

2000 2400 TEMPERATURE, K

2800 3200

FIG. 5. Phase diagram of the silicon-containing products of the reaction SiO 2 and C. After Gol'dshleger and Merzhanov [26]. Reproduced with permission of Plenum Publishing Corp., New York.

system, S i -C-O, is of great importance to this study. Krivsky and Schumann [65] showed that there is no evidence that silicon carbide is present as a single-phase condensed product at 1427~ (1700 K). At higher temperatures the S i - C - O equilibrium diagrams show that silicon carbide formation, as a single-phase condensed product, occurs only in very narrow regions. A clearer picture of these relationships by Gol'dshleger and Merzhanov [26] is depicted in Figs. 5 and 6. In the first figure, the equilibrium reaction products between SiO 2 and carbon with temperature demonstrates that silicon carbide formation begins above the melting point of SiO 2 and continues up to 2427~ (2700 K). In Fig. 6, the equilibrium compositions of the products from the reaction of the S i O - C - O phases are shown at 2127~ (2400 K). A maximum in silicon carbide concentration occurs around the 60% level of SiO 2 concentration. The silicon carbide level drops dramatically at high SiO 2 concentrations with the formation of SiO(g) and SiO 2 as the condensed phases.

In the presence of excess carbon concentration or low SiO2, silicon carbide is the major phase present, as shown in the Fig. 6. With decreasing carbon content, the reduction product of SiO 2 becomes gaseous SiO. In the presence of excess carbon, the SiC is synthesized via a gaseous intermediate according to:

C + S i O 2 , SiO + CO (4)

2C + SiO , SiC + CO (5)


1.0

O I 8

l c co , __.._ sioo (~ 0.6

O.4 8

O.2

O0 0.0 012 ...... 0.4 016 018 1.0

AMOUNT OF Si02

FIG. 6. Equilibrium composition of the products in the reaction of SiO 2 with carbon at 2400 K. After Gol'dshleger and Merzhanov [26]. Reproduced with permission of Plenum Publishing Corp., New York.

Other reactions that are possible in this system include S i ( s )+ CO(g) (favorable at low temperatures), or CO + SiO and C(s) + Si(s, l, or g). These reactions result in producing SiO 2 if the partial pressure of oxygen is too high. Figure 7 shows the relationship of SiO2 versus SiC formation for the maximum partial pressure of oxygen-containing species CO and SiO in the temperature range for typical SiC whisker syntheses as developed by van Konijnenberg [64]. The chemical reaction of most commercial significance is given below:

SiO 2 + 3C , SiC + 2CO (T > 1786 K) (6)

3. SILICON CARBIDE CRYSTAL STRUCTURE

Silicon carbide consists of many structurally different crystalline forms. The many variations result from slight shifting of the basic coordinating tetrahedra of SiC 4 and CSi 4. The fourfold coordination of both silicon and carbon atoms is satisfied by sharing of corners. The crystalline forms of silicon carbide have been identified during the past century and have been loosely grouped into two general types. The /3-SIC describes the cubic form of silicon carbide identical to the zinc-blend (ZnS) structure. The c~-SiC refers to the family of SiC structures related to the zinc-blend or wurtsite-type crystal structures. Notations describing the various types of SiC are given with respect to the stacking sequence of the repeating unit


10-1

E 10-2

.=_

6 " 1 0 - 3 O3 0..

1 0 .4 .,.-., o o .-- 10-5

10 .6

/ " .... ; J18". 10 -1

~ 1 0 -6 I I I I

1200 1400 1600 1800 2000

Tempera tu re in K

)2.0

FIG. 7. The maximum SiO or CO partial pressures above which formation of SiO2 is possible, whereas below SiC is formed. After van Konijneberg [64]. Reproduced with permission of the Institute of Ceramics, Shelton.

cell. There are more than 70 different polytypes, as opposed to polymorphs, with some of the repeating structures requiring several hundred layers. The most common polytypes of silicon carbide have been given the nomenclature of 2H, 3C, 4H, 6H, and 15R where the number refers to the number of layers in the unit cell and the letter designates the crystal structure: C = cubic, H = hexagonal, and R = rhombohedral. These five materials have short-range periodicity and are considered to be the most stable forms.

The form that silicon carbide takes depends on many factors including thermal history, impurity type and level, and environment. The/3 form is generally felt to be the stable phase at low temperatures, whereas the a form is the high-temperature form. There are many exceptions to the rule, as the conversion to a from /3 and the converse have been reported. The stability and transformations of the various polytypes vary among themselves and constitute a subject that is too broad for this effort. The basic a and/3 descriptors will be used for the remainder of this section.

4. PROCESSES TO PRODUCE WHISKERS

The formation of silicon carbide whiskers can occur through one of three general methods: vapor-solid (VS), chemical vapor deposition (CVD), vapor-liquid-solid (VLS). The processes vary widely in the raw


materials used, the environments, and process economics. The morphology of whiskers resulting from these processes can show more than one form of whisker morphology present in a single synthesis batch. Mechanisms that cause the growth of the acicular crystals are not fully understood. The following discussions briefly cover the whisker formation step for each type.

a. Vapor-Solid Method. The formation of all silicon carbide whiskers occurs with the presence of gaseous phases sometime during the process. A straightforward mechanism for whisker growth is the solid-vapor mechanism in which the SiO(g) reacts at a carbon substrate surface as described in the section on the S i - O - C phase diagram. Temperature affects the growth kinetics as do vapor concentrations and impurity levels. Metal impurities can increase the formation of whisker morphologies although there is argument as to the exact mechanism that occurs. Work by Joo and Kim [53] described the Acheson method, a derivative of the VS mechanism, for whisker generation.

An example of a commercial process that is generally felt to follow this VS mechanism is the thermal decomposition of rice hulls first reported by Lee and Cutler [70] and Sharma et al. [111]. They took waste rice hulls that contained cellulose as the major component, but had a very high residual ash content (approximately 15 to 30 wt %). The ash chemistry was > 95% SiO2, with small amounts of alkalies and trace impurities. The procedure to form particulate silicon carbide from rice hulls consists of driving the volatile organics from the material at about 900~ in the absence of air with further treatment in a CO atmosphere up to 1500~ Some fraction of the materials forms whiskers, but the addition of a catalyst such as Fe makes whisker formation dominant. The whiskers produced during the decomposition of rice hulls have submicron diameters with aspect ratios of a hundred or more. In the early work of Lee and Cutler and later Sharma et al., evidence of the catalytic effect similar to that described in a later section on the VLS mechanism of whisker growth was given. A variety of morphologies are also seen in the whisker product by the VS method so that other mechanisms may also play a role in the growth patterns. The morphology of some commercial SiC whiskers, Tateho SCW and American Matrix, is shown in Fig. 8.

b. Chemical Vapor Deposition (CVD). CVD techniques are well known for a variety of applications from coating to particulate synthesis. A myriad of materials can be utilized as feeds to deposit silicon carbide in a variety of forms. In all cases, the morphology of the deposited silicon carbide depends on two major factors: the degree of supersaturation and the temperature at which the deposition is made. The lower the supersaturation, the slower the growth and the greater the tendency to form larger crystals over time. The higher the temperature of deposition, the stronger

5 W H I S K E R S A N D W H I S K E R - R E I N F O R C E D C E R A M I C S 167

FIG. 8. Typical morphologies of commercial SiC whiskers produced by Tateho (top photograph) and American Matrix (bottom photograph). Magnified 5000 x .


FIG. 9. The CVD SiC whiskers (courtesy of L. Riester from ORNL).

the tendency to form larger crystals such as whiskers. The characteristic morphology of CVD whiskers is shown in Fig. 9. These whiskers were produced at Oak Ridge National Laboratory (ORNL) by heating precour- sors in presence of a Cu catalyst.

Chin et al. [16] developed the chemical composition-process relationships in CVD-SiC from Hz-CH3SiC13. The phase relationships given in Fig. 10 indicate that the two parameters of supersaturation and temperature are joined by the mass transport requirements in terms of the rate of deposition of silicon carbide. In this case, the rate is also dependent on the hydrogen concentration as it is required in the reaction to form SiC and is the carrier gas for the reactants. At low hydrogen levels, pyrolytic carbon is formed.


, ~ stoichiometric

14oo ;

/1- L l . . . . . . . . . . . r 3 ( 1 )

% A "> 10 20 30 40

H2 / CH3SiCI3 ratio

FIG. 10. Chemical composition-process relationships in CVD SiC from H2-CH3SiC13. After Chin et al. [16]. Reproduced with permission of Elsevier Science Publishing Co., New York.

c. Vapor-Liquid-Solid (VLS) Method. The VLS technique for growing /3-SIC whiskers was originally proposed by Wagner and Ellis (130). Re- finements of the process were done by Shyne and Milewski [112], Milewski et al. [79], and Knippenberg and Verspui [60]. The terms used in the process title refer to the physical states of the three phases that interact to form the SiC whisker. Vapor refers to the form of the gaseous silicon feed materials. Liquid refers to the intermediate catalyst that dissolves the silicon into the melt and reprecipitates out the silicon carbide product, which is the solid. The crystals grow on a substrate at the solid-liquid interface due to supersaturation in the liquid. The presence of the liquid catalyst is unique to this process and allows the crystallization of extremely uniform single-crystal silicon carbide whiskers.

Figure 11 gives a representation of the growth of silicon carbide whiskers by the VLS method. In this case, the silicon source is reacted with hydrogen to provide elemental silicon, which dissolves into the catalyst liquid. The catalyst can be selected from a variety of materials including transition metals or iron alloys. Crystals grow from the substrate carrying the catalyst liquid at the tip of the whisker, allowing for resupply of silicon from the atmosphere and continued growth. The size of the catalyst droplet controls the diameter of the resultant whisker.

Wagner et al. [131] describe their observations of whiskers as dislocation free and suggest that the whiskers form in a two-stage process: rapid growth, which forms a filament, with a second step of a slow thickening of the whisker. They conclude the main growth of the whisker is along the { 111} axis, which is the direction of slow growth. With this theory came the


Vapor Feed H2 CH4

Vapor Feed SiO (

VLS Liquid Catalyst Supersaturated With Si And C

VLS Solid Catalyst / , / [ '~ , . .~ ~ (e.g. Steel) 30 i~m

)

Solid SiC Crystals __. I

FIG. 11. Illustration of the VLS process for SiC whisker growth. After Shalek and Parkinson [110]. Reproduced with permission of Materials Research Society, Pittsburgh.

identification of a liquid layer between the vapor and solid, which is the major difference from the VS mechanism.

A variation to the supply of reactant SiO has been demonstrated by Milewski et al. [79] who used an SiO generator that is simply S i O 2 brick impregnated with silica and carbon to produce SiO at temperature. The SiO then dissolves into the liquid catalyst. Carbon was introduced as C H 4

carried by a mixture of H z, C O , and N 2. The kinetics of the growth rate are then controlled by the composition of the surrounding atmosphere, which controls the partial pressure of SiO to the liquid. The stoichiometry

TYPE AND COLOUR

OF GROWTh

D BALLS

BRANCHES

WOOL

NEEDLES

1

BLACK

BLACK 1D

BALLS

BLACK

BRANC~

BLACK 1B

WOOL

BLACK 1A

NEEDLES

2 BLACK & GRE

BROWN 2D

BALLS

BROWN 2C

BROWN 2B

WOOL

BROWN 2A

NEEDLES

CARBON RICH GASES

- - INCREASED CH4 CONTENT - - REDUCED N 2 FLOW = = = REDUCED CO FLOW

3 WHITE

WHITE 3D

BALLS

WHITE 3C

3RANCHESI

WOOL

WHITE 3A

'4EEDLES

4 LIGHT GREEN

LT. ~G~EEN ~.(~

BRAIIC~ES

LT. GREF~ 4B \

"~OL

LT. GREE-~N 4A

NEEDLES

i

5 6 GREEN BLUE

GREEN BLUE 5D 6D

BALLS BALLS

GREEN BLUE 5C 6C

3RANCHES BRANCHES

GREEN BLUE 5B 6B ~ OL WOOL

'C..-~REEN . _B_L.VE_ _ _j 5A 6A

qEEDLES NEEDLES

HIGH SUPER SATURATION~ ARE UP +

LOW SUPER SATURATIONI ARE DOWN +

MOST ALL FIBRE BALLS �9 *****, *

FIBRE TYPE SIDE BRANCHES ON NEEDLES O R / ~

WOOL ~____.Z~.__2~ ~

WOOL FIBRE d 1-3 MICRON DIAMETER

~:> FIBRE 4-8 " ~ MICRONS DIAMETER

SILICON RICH GASES

FIG. 12. Identification chart for SiC whisker growth conditions. After Milewski et al. [79]. Reproduced with permission of Chapman & Hall, London.

5 WHISKERS AND WHISKER-REINFORCED CERAMICS 17/

FIG. 13. The morphology of the VLS SiC whiskers (courtesy of T. Tiegs from ORNL).

and morphology of the product whisker are controlled by the ratio of carbon to silicon in the surrounding atmosphere. Evidence of these effects is seen in Fig. 12. The whiskers change colors from black to white to green and eventually to blue when going from highly carbon rich to highly silicon rich atmospheres. The VLS method is known to produce crystals having large diameters and high aspect ratios. VLS SiC whiskers fabricated at Los Alamos National Laboratory, with characteristic smooth surfaces, large diameters above 1/xm, and lengths of a few hundred microns are shown in Fig. 13.

5. CHARACTERIZATION OF WHISKERS

The bulk chemistries, surface chemistries, and morphologies of silicon carbide whiskers vary widely depending on the type of process used, the stage of the process development, and the postsynthesis treatments prac- ticed by the producer. The synthesis of the whiskers as described earlier is


done with different starting materials, atmospheres, temperatures, and growth mechanisms. The utility of the silicon carbide whiskers in composites depends on the ability of the composite fabrication process to accom- modate these variations and allow the proper toughening mechanisms to take place when under stress. Unfortunately, even some materials that come from the same process have significant lot-to-lot variations that can have negative results in the ease of fabrication or in the final composite properties.

The bulk analysis of /3-SIC whiskers shows the least variation in chemistry. In some whiskers, the residual metals content can vary, most likely, as a result of additives that used as catalysts during synthesis. These include iron, cobalt, and chromium. Studies by Karasek et al. [56] have shown that the physical properties of silicon carbide whisker-reinforced composites do not correlate to the bulk properties of the whiskers significantly. This lack of significant correlation is mainly due to the fact that the important phase chemistry of the whisker-matrix interface is controlled by the matrix chemistry and the surface chemistry of the whiskers. There seems to be little impact of the diffusion of materials into or out of the bulk whisker material.

The surface chemistry of the whisker is a very important parameter in defining the utility of the whisker to toughen materials. In a ceramic matrix composite, a tightly bonded whisker does not allow for a toughening mechanism such as pullout to occur. It has been shown that a loosely bonded whisker is desirable so that some energy is dispersed during the interracial rupture and more is dispersed by the frictional forces during the pullout. Taylor [120] argues that the surface chemistry can also affect predensification processing by controlling the stability of dispersions. The stability of the dispersions can, in turn, affect the amount of sintering aids required. The added oxide content present at the surface can change the phase chemistry during the heat treatment used during densification. If consistent, the oxides at the surface may be utilized to reduce the amount of the sintering required. If inconsistent, the whiskers will cause uncontrolled second phase chemistry, thereby limiting performance and /o r reliability of the final composite.

The surface chemistry of the whiskers varies from producer to producer and from production lot to production lot from a single producer. The desirability of low amounts of surface impurities is acknowledged and, in the later generations of products, the control of surface oxides becomes evident. Karasek et al. [57] have shown that the resulting surface phases tend to be more of a S i - O - C glass rather than thick coatings of SiO 2. Whiskers produced from the plasma- or gas-phase approach have a suboxide chemistry with the presence of graphitic carbon. In many cases,


the chemistry of the whisker surface contains metal impurities that are carried over from the production processes. It has been shown that these materials are responsible, at least in part, for the thermal decomposition of whiskers in the composite at high temperatures.

The morphology of whiskers ranges from straight, smooth surfaced whiskers to curved or branched whiskers with a considerable amount of defects present. Occasionally, a small percentage of hollow whiskers is formed, apparently by a secondary growth mechanism. Some of the products contain a high percentage of particulates that are detrimental to the final composite properties. The diameters can range from 1 /xm to 20 nm with lengths of greater than 100 /xm. The defect structure most prevalent in the irregular whiskers were grouped by Karasek et al. [57] as Iwanga type b, which describes stacking faults perpendicular to the growth axis. A lesser occurrence of the Iwanga type a defect with stacking faults at angles not equal to 90 ~ to the growth axis have also been seen. The physical properties and chemical compositions of some commercial SiC whiskers are presented in Table I.

C. Silicon Nitride (Si 3 N4 ) Whiskers

1. PHASE RELATIONSHIPS IN THE S i - O - N SYSTEM

Silicon nitride occurs in two forms, alpha (a) and beta (/3), with similar hexagonal crystal structures, but different cell dimensions. The a phase has almost twice the cell volume of/3 (a has an a axis = 0.775-0.777 nm and a c axis = 0.516-0.569 nm, and /3 has a = 0.759-0.761 nm and c = 0.271-0.292 nm) [121]. In contrast to SiC, these two forms cannot be transformed from one to the other without a reconstructive processing, usually via a liquid phase.

Both phases are built up of SiN 4 tetrahedra joined in a three-dimen- sional network by sharing corners. In a -S i3N 4 the bottom half of the unit cell is the same as /3, but the top half is related to the bottom by the operation of a c-glide plane. The bond lengths and angles in a are less uniform than in /3 [121]. Although there is good agreement on the structure of the /3 phase, there is ongoing discussion on the structure of a-Si3N4, and the a-/3 relationship. Blegen's calculations [8], based on temperature-pressure relationships, led to a conclusion that te-Si3N 4 is not a true nitride, but rather the oxynitride represented by the Si23N300 formula, with a range located between fl-Si3N 4 and SizN20. The determination of Wild et al. [136] pointed to a defect structure and also to the substitution of oxygen for nitrogen on N(1) sites. Jack's [50] determination

TABLE I

PROPERTIES OF COMMERCIAL SiC WHISKERS

Tateho Chemical Advanced Composite Tokai

Source Industries Materials Company (ACMC) Carbon American Matrix

Product name SCW SC-9 TOKAWHISKER AM Crystalline type /3-SIC /3-SIC /3-SIC /3-SIC Density, g / c m 3 3.18 3.2 Diameter, ~m 0.05-1.5 0.2-1.5 0.3-1.5 0.4-2.2

(0.25) (0.6) (1.1) (0.9) Length,/xm 5-200 10-100 10-80 10-100

(15) (20) (40) (15) Aspect ratio 20-200 20-100 10-50 10-50 Morphology Rods Rods Rods Rods

Chemical composition, wt %: C 28.5 28.6 28.7 29.6 Mg 0.2 0.02 0.02 0.2 Ca 0.5 0.11 0.05 0.6 A1 0.6 0.07 0.08 0.2 Fe 0.1 0.01 0.02 0.02 O 0.6 0.9 0.2 0.6

Sources: T. N. Tiegs and J. W. Geer, "SIC Whisker," Reinforced Ceramic Composites Handbook. Oak Ridge National Laboratory, Oak Ridge, TN, 1987. Karasek et al. [56]. Manufacturer 's literature data Own characterization

5 WHISKERS AND WHISKER-REINFORCED CERAMICS 17.5

18

20 o

g~

22

24

SiO2 Si

26 -6 -4

SiO2 ..."" """"

si

\ Si2N20 ", , , S 13-Si3N4

/

-2 0 2 log (PNJ

FIG. 14. Thermochemical diagram for the S i - O - N system. After Jack [51]. Reproduced with permission of Kluwer Academic Pubs, Norwell.

of stable phases at 1327~ and 1527~ shows that, at higher temperatures, the only stable form is /3-Si3N 4 (Fig. 14). At higher temperatures, more nitrogen is required in the system to produce silicon nitride. Low nitrogen partial pressure leads to the formation of deleterious SiO 2 or Si metal, depending on the partial pressure of oxygen.

The review of theories explaining the a-/3 relationship is given by Mitomo [81] and summarized in Table II. The most recent theories assume that ce-Si3N 4 is a metastable form whose formation is favored by low temperature and certain reaction routes in which nitrogen is available in the atomic form [52].

TABLE II

THEORIES EXPLAINING THE t~-/~ RELATIONSHIP

Theory a fl

Low temperature High temperature Unstable at any temperature Stable at any temperature

High po 2 Low po 2 ' Stable in low temperature and high P02 regions Stable at any temperature

Source." Mitomo [81].


Another phase of the S i - N - O system is silicon oxynitride (Si2N20) , is the only compound in the Si3N4-SiO 2 system. The formation of SizN20 from Si3N 4 and SiO 2 only occurs via a liquid phase provided by the eutectic of composition 98 mole % SiO 2 and 2 mole % Si3N 4. The investigations of Huang et al. [45] determined the eutectic temperature to be 1590~ The same authors found that below 1450~ no SizN20 was detected. Above 1750~ gradual decomposition of SizN20 to j3-Si3N4(s) , Si(/), SiO(g), and Nz(g)was observed. The kinetics of SizN20 formation from Si3N 4 and SiO 2 is highest at 1700~ [45]. The oxidation study conducted by Du et al. [20] showed that the duplex oxidation scale consisting of SiO 2 and SizN20 formed on the Si3N 4 surface. The intermediate SizN20 layer was identified as a single-phase material rather than a physical mixture of Si3N 4 and SiO 2. The SizN20 structure is built up of SiN30 tetrahedra joined together by sharing corners. The orthorhombic crystal structure consists of irregular, but parallel, sheets of covalently bonded silicon and nitrogen atoms linked by S i -O-Si bonds. Because of the structural similarities between nitride and oxynitride, there is some limited solubility and a solid solution can be formed without a change of structure.

Whenever silicon nitride is synthesized in the presence of aluminum- containing compounds (frequently used as a flux material in process of growing whiskers), there is a high probability of the formation of /3'-SiA1ONs. Up to two-thirds of the silicon in/3-Si3N 4 can be substituted by A1 without a change of structure. The/3'-SiA1ON has mechanical and physical properties similar to /3-Si3N 4. It is, however, thermodynamically more stable than silicon nitride.

2. PROCESS TO PRODUCE WHISKERS

The formation of silicon nitride whiskers was observed in several different reactions, including vapor deposition, CVD, and growth from a melt. However, only the following techniques are considered to have commercial significance: nitriding of metallic silicon or silicon-silica mixture, carbothermal reduction of silica with simultaneous nitridation, and thermal decomposition of silicon halides.

a. Direct Nitridation. The direct nitridation technique is based on reactions of Si or SiO 2 in a N 2 + H 2 atmosphere where

3 Si + 2 N 2 , Si3N 4 (7)


It was shown [29] that the growth of whiskers from a mixture of SiO 2 and Si requires specific substrate materials (mullite or graphite) or the presence of A1 and Fe catalysts in silicon at the growth zone. The a-type whiskers are synthesized in the range of 1200 to 1400~ while the/3-type tend to be formed above 1400~ The reaction is strongly exothermic, which causes problems in controlling the synthesis temperature. Besides temperature, which requires precise control to form pure Si3N4, the composition of the vapor phase, particularly the oxygen and nitrogen partial pressures present under equilibrium conditions, is of practical importance. Figure 15 shows the reaction products in the S i - N - O system depending on the oxygen and nitrogen partial pressures. At 1200~ the oxygen partial pressure above 10 -22.5 atm produces conditions under which SiO 2 is the only stable phase. At lower oxygen partial pressures, small variations in the nitrogen content change the phase stability. The thermodynamic condition for which it is possible to form only one phase, e.g., pure a or pure /3, are extremely limited. This explains why a-/3

1200~

-22

-24

-26

o -28

-30

-32 -

- Si02

_ Silicon ' ~

Si2N20 + e~ + 13

~:~ +13

-4t7 -4[6 -5"5 -41"4 log PN2

FIG. 15. Thermochemical diagram for silicon nitrided at 1200~ showing products of reaction, pressure in atmospheres. After Colquhun et al. [17]. Reproduced with permission of the Institute of Ceramics, Shelton.


mixtures are usually obtained during the formation of silicon nitride whiskers.

The mechanisms of whisker growth vary depending on the substrate used for the reaction. For example, the Si3N 4 whiskers grown on mullite show axial defects, characteristic of crystallization by an axial-screw-dislocation mechanism [29]. In the presence of a catalyst, whisker growth occurs by means of a VLS mechanism, involving crystallization of the nitride from drops of binary or ternary alloys of silicon with iron and aluminum. The iron and aluminum act as solvents for silicon.

b. Carbothermal Nitridation. Carbothermal reduction and simultaneous nitridation of silica according to

3 SiO(g) + 3 C O ( g ) + 2 Nz(g) = Si3N 4 + 3 CO2(g) (8)

is a potential method for producing high-purity silicon nitride whiskers. Its advantage over direct nitridation is better control of the processing temperatures (direct nitridation is strongly exothermic and thus disposed to thermal overshooting). Several processes have been developed and reported. According to Rahman and Riley [104], the decomposition of pyrolyzed rice hulls at a temperature range from 1260 to 1450~ under 95% nitrogen-5% hydrogen leads to formation of a-Si3N 4 whiskers. The ratio of whiskers to particulates in the final product increases with increasing size of pyrolyzed grains to above 350 /zm. Wang and Wada [133] formed a-Si3N 4 and /3'-sialon whiskers by reacting a mixture of silica, carbon, and halide (3NaF, A1F3, or A1F) in a graphite furnace with high purity (N 2 + 3% H 2) gas. The reactions were carried out at 1350~ for 10 h. Reaction products consist of/3 ' -s ialon and a-Si3N 4 whiskers. The fi '-sialon whiskers had a 0.7-/zm diameter and the a-Si3N 4 was characterized by the ribbon-like morphology with a 1-/zm width and about a 0.1-/zm thickness. Produced at Dow Central Research, a-Si3N 4 whiskers have larger sizes but the same ribbon morphology (see Fig. 16).

Hayashi et al. [36] have grown Si3N 4 whiskers by the nitridation of a SiO2-C-Na3A1F 6 mixture. This produced whiskers containing about 85% fi'-sialon and 15% a-Si3N 4. The /3'-sialon whiskers were 1 to 10 ~m in diameter and up to 10 mm in length.

Niwano [89] described an industrial-scale process where a -S i3N 4 whiskers could be obtained by nitriding a mixture of colloidal silica, silicon, and lampblack. Whiskers were approximately 50 ~m in length with a mean diameter of 0.9 ~m. It was observed that these whiskers have two preferential types of morphology: triangular prisms and ribbons. The morphology of commercially available a -Si3N 4 whiskers (Tateho SNW-1) is illustrated in Fig. 17. The experimental observations showed that the


FIG. 16. ce-Si3N 4 whiskers with characteristic ribbon-like morphology. Magnified 149 x .

most common mechanism of whisker growth occurring during carbothermal nitridation was a solid-vapor mechanism.

c. Diimide Decomposition. The diimide thermal decomposition method is based on reaction of SiC14 and NH 4 to synthesize silicon diimide, Si(NH) 2. In the following step, the Si(NH) 2 is thermally decomposed to produce amorphous silicon nitride powder. Amorphous powder can be then used as a raw material to produce ol-Si3N 4 powder, /3-SIC (in the presence of CO or CO2), or silicon nitride whiskers when metal catalysts such as Fe are present. The Si3N 4 manufactured by this process has diameters ranging from 0.1 to 0.4/~m and lengths of about 5 to 20 ~m. These whiskers grow along the c axis by a VLS mechanism [84].

The modified diimide decomposition process uses gas mixtures of Si2C12, NH3, and H 2. Amorphous whiskers are deposited on quartz or graphite substrates coated with a metal catalyst. The uniqueness of this method is the ability to form whiskers of a special shape, e.g., a coil. Whiskers with the best coil regularity have been obtained at 1200~ on a graphite substrate with Fe as the catalyst. All whiskers have a very smooth surface,


FIG. 17. Prismatic morphology of a -S i3N 4 Tateho SNW-1 whiskers. Magnified 5000 x .

are 2 to 5/zm in diameter, and have a coil diameter of about 10 to 15/xm. Spherical droplet-like nodules, rounded tips, smooth surfaces, and uniform diameters indicate growth by the VLS mechanism [48].

3. PROPERTIES OF S i 3 N 4 WHISKERS

To satisfy toughening criteria, Si3N 4 whiskers, like SiC whiskers, should be characterized by a lack of overgrowths, defects, and imperfections. Smooth surfaces with low levels of impurities are required to reduce interfacial reactions. The Si3N 4 whiskers obey a strength-size relationship where strength = constant x cross-sectional area -z (0.8 < z < 1.1) [3]. At the same time, experimental evidence shows that debond length increases as the whisker radius increases [4]. Since a small diameter is required to increase whisker strength and a large diameter increases the debond length of the matrix-whisker interface, the optimum whisker size will vary depending on composition of the matrix. A summary of physical,

5 W H I S K E R S A N D W H I S K E R - R E I N F O R C E D C E R A M I C S

TABLE III

PROPERTIES OF Si3N 4 WHISKERS

Tateho Chemical UBE Industries Industries Ltd.

181

UBE Industries Ltd.

Product name SNW UBE-SN-W UBE-SN-WB Crystalline type a-Si3N4 a-Si3N4 fl-Si 3N4 Density, g/cm 3 3.18 3.18 3.19 Surface area, m2/g 1.92 1.21 2.07 Diameter,/zm 0.1-1.6 0.1-0.4 0.1-1.5 Length,/zm 5-200 5-20 10-50 Aspect ratio 20-200 20-100 20-100 Morphology triangular prisms ribbon-like rods

Chemical composition, wt %: Mg 0.2 m Ca 0.5 0.01 0.01 A1 0.2 0.01 Fe 0.1 0.5 0.07 Y - - m 0.9 O 00 2-3 0.4

Sources: Niwano [89]; Homeny and Neergaard [41]; Homeny and Vaughn [43]. Product literature from Tateho and UBE

chemical, and structural properties for three commercial silicon nitride whiskers is presented in Table III.

D. A1203-SiC Whisker Composites

]. THERMAL STABILITY OF THE A 1 2 0 3 - S i C SYSTEM

Thermodynamic calculations conducted by Homeny et al. [44] predicted that no reactions would take place between A120 3 and SiC even at very high temperatures with the only melting phase being AleO ~. The alumina melts at 2015~ which is above the processing temperature required to densify this material. The presence of SiO 2 and C, however, changes this equilibrium. The SiO e from the SiC surface reacts with AleO 3 to form mullite (3A1203 2SIO2) , and carbon reacts with A120 3 to form AlaC 3.

The stability of these phases is governed by the total pressure of gases produced during the reaction. If the pressure at the A1203-SiC interface exceeds the ambient pressure, the interface can be considered as chemically unstable. According to Misra [80] the carbon-mullite interfaces become unstable above 1567~ (1840 K), whereas the SiC-mullite inter-


.30

.25

.20

.15

.10

.05

I I I

SiC-C-Mullite

/ I I I

1800 1900 2000

TEMPERATURE. K

SiC-Mullite-Si

0 I I 1700 2100 2200 2300

FIG. 18. Total pressure of gases corresponding to the SiC-Si-mullite and SiC-C-mullite equilibria as a function of temperature. After Misra [80]. Reproduced with permission of the American Ceramic Society, Westerville.

faces become unstable above 1867~ (2140 K) (Fig. 18). In composites having SiC whiskers with a carbon-rich surface (high C/SiO 2 ratio), all SiO 2 would be converted to SiC above 1567~ At these conditions the AI203-SiC-C interface is stable until an AlzO3-A14C 3 eutectic forms at 1947~ The solubility of SiC in the A1203-A14C3 melt is not known. However, the sintering experiments suggest that SiC must be soluble in A 1 2 0 3 - A 1 4 C 3 melt because densification occurs by a solution-reprecipitation mechanism. In A1203-SiC composites with oxidized and /o r carbon- free whiskers (low C/SiO 2 ratio), all free carbon would be converted to SiC above 1567~ Under these conditions, SiO 2 reacts with A 1 2 0 3 to

form mullite. Alumina saturated mullite melts above 1827~ [1] forming an aluminosilicate melt. At this temperature, SiC, A1203, and molten mullite are still chemically stable. At 1867~ the SiC/mullite becomes unstable, resulting in the formation of the liquid silicon. Although the solubility of SiC in aluminosilicate melt is doubtful [80], the solubility in molten Si reaches about 8 mole % at 1867~ The modification of interfacial bonding between A1203 and SiC can, therefore, be expected above 1827~ and dissolution of SiC whiskers can take place above 1867~ Both changes reduce the properties of A1203/SiC whisker composites.

The critical temperature above which the modification of the A1203-SiC interface occurs is often even lower than 1827~ because pure SiC whiskers and pure A]20 3 powders are not commonly used. Typically SiC contains


residual amounts of Fe, Co, Si, or Ni, which are used as catalysts in the whisker production. Alumina contains small amounts of sintering additives, such as CaO and MgO. SiO2-AlzO3-CaO and SiO2-AlzO3-MgO form low-temperature eutectics, which alter the interracial bonding. The presence of sintering additives can also change the crystallization products formed during cooling or during high-temperature use. Powell-Dogan and Heuer [95,96] studied the crystallization characteristics of an A120 3 ceramic containing MgO and CaO. They showed that the crystallization of high-magnesia grain-boundary glasses resulted in the formation of or- thoenstatite ((A1)MgSiO3), a-cordieri te (MgzA14SisOl8) , forsterite (MgzSiO4) , sapphirine (Mg4A110Si2023) , and spinel (MgAI204). The crystallization of high-calcia glasses produced equally complex phases consisting of anorthite (CaAlzSi208) , gehlenite (CazAlzS i307) , a solid solution of grossularite (Ca3AlzSi3012) and pyrope (Mg3AlzSi3Ol2) , calcium hexa- luminate (CaAll2019), and spinel.

The different fabrication routes result in A120 3 powders and SiC whiskers with different surface characteristics. The interracial chemical compositions vary depending on the combination of whiskers and A1203. This causes the formation of a liquid phase and the chemical reactions at the A1203-SiC interface to occur at different processing temperatures. Therefore, conditions selected to achieve full density also have a critical influence on interfaces and on material mechanical properties. Some combinations of A120 3 and SiC work better than others, but all require individual optimization of processing conditions.

2. A1203-SiC INTERFACES

The ceramic grains can interface directly, similar to metals, or via a grain-boundary phase. Such interfacial films often tend to be amorphous providing sufficient structural flexibility to link the different bond arrays on either side of the boundary. In AI203-SiC whisker composites, the presence of the interracial layer is well documented. Sarin and Ruhle [108] have discovered the presence of an amorphous layer less than 5 nm at the interface between the whisker and the matrix. Barrett and Page [2] showed the existence of an 0.8-nm-thick layer at the AI203-SiC interface and 0.7 nm at the SiC-SiC boundary. Ravichandran and Knowles [105] were able to show evidence of the presence of about a 2-nm-thick amorphous layer at these interfaces. The thickness of the interfacial layer at the whisker-matrix interface depends in large degree on the purity of the AI20 3 and SiC. However, even in AI203-SiC composites processed from u|traclean materials, a 1-nm layer was detected [94].

184 A L E K S A N D E R J. PYZIK AND ALAN M. H A R T

Since the behavior of interfaces is known to be crucial for toughening, their control is of special importance. Most attempts to develop weak matrix-whisker interfaces centered on modifying the whisker surface chemistry by subjecting the interfaces to thermal or chemical treatments. The presence of a silicon oxycarbide species and free carbon was associated with good debonding of whiskers and resulted in composites with a high fracture toughness. The formation of oxynitride glass and aluminosilicate glass evidently increased the extent of whisker-matrix interfacial bonding and resulted in a reduced fracture toughness [41, 43, 124, 125]. The investigations of the crack path in high fracture toughness AI203-SiC composites show that the matrix cracks are mostly transgranular and conchoidal, while cracks in the whiskers are planar and always parallel to the basal plane of SiC [13]. The length of the debonds ranged between 2R and 6R, R being the whisker radius, and the bridging zone near the crack tip was about 4 to 6 whisker diameters.

3. FABRICATION AND PROPERTIES

Properties of A1203-SiC composites depend to a great extent on the characteristics of the whiskers and on the whisker-matrix interfaces. It has been shown that while the maximum fracture toughness is achieved at 20 to 40 vol % of SiC whiskers, the maximum flexure strength is observed between 25 and 30 vol % [123-125]. The requirement to optimize both properties causes most SiC-reinforced aluminas to have 25% to 30% of whiskers. For A1203 with 20 to 30 vol % SiC whiskers, Becher and Wei [6], Becher et al. [7], and Homeny et al. [44] have achieved fracture toughness values between 8 and 10 MPa m 1/2. Tiegs [122] showed that at the same whisker content, the fracture toughness increases with increasing whisker diameter. Homeny and Vaughn [42] demonstrated that the fracture toughness of A1203-SiC could vary from 4.0 to 9.0 MPa m 1/2a when utilizing whiskers with similar aspect ratios, but different chemistries.

The effect of whisker-matrix interfaces on mechanical properties becomes clear when the same A1203-SiC w materials are hot-pressed at different temperatures. While maintaining similar density, similar sizes of alumina grains, and identical whisker volumes, these composites have different fracture toughnesses as shown in Fig. 19. The differences in Kic values between A1203 with ACMC SC-9 SiC and A1203 with Tateho SCW

aMeasured by chevron notch technique (CNT).


. 10

9 13_

8 o0 uJ z 7 "1- (.9 0 6 I-- uJ rr 5

I-- 0 4 13:: u_ 3

- o - SC-9 SiC whiskers SCW SiC whiskers

~,.. I I I 1 ! 1700 1750 1800 1850 1900

HOT PRESS TEMPERATURE. C

FIG. 19. Fracture toughness of AI203-SiC whisker composites as a function of the hot-pressing temperature.

SiC whiskers came mainly from the morphology and bulk properties of the whiskers themselves. However, the differences in toughness values for composites made from the same materials are due to the changes taking place at the whisker-matrix interfaces. All commercially available whiskers produce materials exhibiting similar behavior. The increase in fracture toughness as a function of the temperature reaches a maximum followed by a reduction. The differences are in the temperatures where the maximum is reached and in the breadth of the temperature ranges where high toughness is maintained. The maximum fracture toughness for A1203-30 vol % SCW SiC whiskers (5.2 MPa m 1/2 CNT) can be achieved between 1780~ and 1800~ The processing window is very narrow; and the fracture toughness at 1825~ is only 4.3 MPa m 1/2 CNT, which is very close to pure alumina (4.1 MPa m 1/2 CNT). In contrast, in the A1203-30 vol % SC-9 SiC whisker composite, the fracture toughness above 8 MPa m 1/2 CNT can be maintained between 1760 and 1875~ and its reduction at higher temperatures is slow. This behavior corresponds well with thermodynamic calculations (see Section II.D.1). The formation of a liquid phase in a mullite-rich interface (SCW SiC) is predicted to occur at about 1830~ and the silicon oxycarbide plus carbon-rich interface (SC-9 SiC) should be stable to about 1950~ A high-resolution transmission electron microscopic (TEM) photograph reveals an interface of a A1203-SCW SiC composite that has been hot-pressed at a temperature at which liquid


FIG. 20. A high-resolution TEM photograph reveals an interface of a AI203-SiC composite with a 2- to 3-nm interfacial layer and a large glass pocket at the top.

formation was expected. The amorphous 2- to 3-nm grain-boundary layer and larger glass pockets are present, as shown in Fig. 20.

A similar correlation is observed for the flexure strength. In this case, however, the maxima always occur at a temperature higher than the maxima of fracture toughness. Thus, in materials with chemically stable whisker-matrix interfaces that allow maintenance of high toughness over a broad temperature range, the maximum toughness and maximum strength can overlap. The AI203-SiC composites with unstable interfaces (interfaces where the chemical changes leading to increased bonding occur at


low temperatures) result in combinations of high toughness and low strength or low toughness and high strength.

E. Si 3 N4-SiC Whisker Composites

1. HIGH-TEMPERATURE STABILITY OF THE Si3N4-SiC SYSTEM

One of the necessary conditions to obtain efficient toughening of ceramic materials with whiskers is maintenance of the physical integrity of the reinforcement and weak interfaces. Therefore, chemical interactions between the system constituents, which are affected to a large degree by the processing conditions, are critical factors in controlling material properties. At the high temperatures required for densification of Si3N4-based composites, the system is under dynamic chemical nonequilibrium conditions. Depending on the N, O, SiO, and CO partial pressures, the surface chemistry of whiskers and the sintering additives, the chemical stability of the system changes.

The role of nitrogen pressure in thermal decomposition of Si3N 4 is illustrated in Fig. 21, which shows the stability diagram for Si3N 4 in equilibrium with Si and N 2. From a practical point of view, safe sintering processes require a nitrogen pressure that is higher than the equilibrium pressure of N 2 for the reaction

Si3N4(s ) , 3 Si(/) + 2 Nz(g ) (9)

At 1 atm of nitrogen, decomposition of Si3N 4 takes place above 1850~ This means that in order to avoid local decomposition, the processing should be conducted well below this temperature. If higher temperatures are required, the nitrogen pressure has to be increased as shown in the diagram. However, at high N 2 and low SiO pressures, SiC whiskers decompose to Si and C. This dissociation splits the stability region of Si3N 4.

At 1 atm of nitrogen pressure, the stable equilibrium, SiC-Si3N4-gas, exists between 1440 and 1877~ (Fig. 22). Above this range, Si3N 4 decomposition to liquid silicon and nitrogen gas is observed according to reaction (9). Below 1440~ nitridation of SiC takes place and SiC forms Si3N 4 and free carbon [74] according to:

3 SiC(s) + 2 N2(g ) , Si3N4(s ) + 3 C(s) (10)

188 A L E K S A N D E R J. P Y Z I K A N D A L A N M. H A R T

I I I I , I . I ~ " I I ' ' ' , - Sl3N4(g ) + Si(i)+ Si(v) + N2(g ) i

CONDmONS WHERE SPONTANEOUS ~ ~ "% 1' "; DECOMPOSITION OF SILICON ~" ~ ~ ~ ,z.,0 occu.s o

.... ..-,,,, - , %

o Si3N4(g) + Si(s)+ Si(,) + Si(v) + N2(g) 800 ~ ~_

Sils) + Si(v) + N2lg) ..(D. #

~OOo "'OOo o

-4 -3 -2 - 1 0 1 2 3 "7,

I I I PN2~ I I I 10 .4 10 .3 10 .2 10 -1 1 101 10 2 10 3

PRESSURE PN2 (ATMOSPHERES)

Fio. 21. Stability diagram for Si3N 4 in equilibrium with Si and N 2. After Greskovich [28]. Reproduced with permission of Kluwer Academic Pubs, Norwell.

2200

ooo \

o o i(0 18oo

Q. E 1600 I-- Si3N4 +C

1 4 0 0 - ~ ~" 4 + Si Si(s) '

1 2 0 0 101 10 -1 10 .3 10 -s 10 -7

p (N2) in M P a

Si(g)

FIG. 22. The Si3N4-SiC phase diagram. After Lorenz et al. [74]. Reproduced with permission of Elsevier Scientific Publishing Co., New York.


Temperature, ~ 1800 1600 1400 1200

" ~i '1 ' ' 1 ac=l "ll -15 - - PN2 - 1 atm-(0.10 MPa) I , -- PN2 = 10 atm (1.01 MPa)I

, __i_ ~

-25 -26

slo 6.o 7'.o 1/T X 10'. K"

Fio. 23. Phase relations in the S i - C - N - O system as a function of oxygen partial pressure and temperature at a c = 1 and nitrogen pressure equal to 1 and 10 atm. After Wada et al .

[128] . Reproduced with permission of the American Ceramic Society, Westerville.

The equilibrium range shown in Fig. 22 indicates that stable composite materials without whisker or matrix degradation can only be achieved when the nitrogen pressure and densification temperature are maintained between the two equilibrium curves for reactions (9) and (10). The boundary temperatures and ratio of Si3N4/SiC in the stability range change greatly as a function of the free carbon, nitrogen pressure, SiO pressure, and impurities. At nitrogen pressure = 1 atm and carbon activity equal 1, the silicon nitride cannot be formed above 1374~ without simultaneous formation of SiC, even at 10 -20 atm oxygen pressure (Fig. 23). However, if the nitrogen pressure is increased to 10 atm, the critical temperature to form Si3N 4 without SiC increases to 1536~ [128]. Thermodynamic calculations [21] show that the SiC/Si3N 4 phase boundary curve tends toward higher temperatures with increased pressure. At 1727~ and a molar ratio of N /S i = 100, nitrogen pressure of 1.6 MPa produced only Si3N4, while at 0.9 MPa nitrogen, both Si3N 4 and SiC are present.

The precise control of the SiC-Si3N 4 equilibrium is necessary for maintaining the stability of SiC whiskers in a Si3N 4 matrix and for forming SiC and Si3N 4 whiskers in situ from mixtures of the above powders. This process is discussed in Section III.

2. INTERFACES IN THE S i 3 N 4 - M e O - S i C SYSTEMS

Pure Si3N 4 powders are difficult to densify, even under high-pressure conditions. Therefore, sintering aids consisting of metal oxides (MeO) are


used. While the ternary S i -C-N system shows SiC, Si3N4, and gas as equilibrium phases, the presence of oxygen produces, in addition, SizN20 and SiO 2 (cristobalite). The SiO2, SizN20 , and other metal oxides (densification aids) form oxynitride glass, which interacts chemically with Si3N 4 and SiC. When SiC whiskers having a SiO 2 surface layer are sintered with Si3N4, the SiO 2 converts to SizN20. In the presence of SizN20 and nitrogen pressure of 0.1 MPa, Si3N 4 and SiC are stable only to 1838~ [87]. Above this temperature decomposition occurs according to the following reactions:

SiC(s) + Si2N20 (s) , 3 Si(/) + Nz(g) + CO(g) (11)

Si3Na(s) + SizN20 (s) , 4 Si(/) + SiO(g) + 3 N2(g ) (12)

The type and degree of these interactions affect Si3N4-SiC interfaces through modification of whisker surfaces and changes in the thickness and crystallinity of the interracial layer.

One of the most noticeable features of the whisker-matrix interface is the presence of microscopic facets on the whisker side [109]. The facets occur on {lll}-type planes and form step configurations. In studying SiC interfaces, Davis et al. [18] concluded that the steps tend to spiral up around whisker sides following the whisker curvature by incorporating short kinks along other {lll}-type planes. Their size depends on degree of SiC dissolution and typically varies from 1 to 3 nm in height. The stages tend to be 5 to 10 nm apart.

Layer facets were also reported to occur [11] as the amount of catalyst used for whisker growth increased. Apparently, the Fe20 3 additions, which react with Si to form a liquid iron silicide during nitridation, react similarly with the SiC whiskers, resulting in their dissolution and the formation of facets on the whisker surfaces. The increase of surface roughness reduces debonding ability and is one of the most significant strength and toughness limiting factors.

The formation of facets, characteristic for SiC whiskers containing Fe impurity, can vary from lot to lot as a result of postfabrication whisker acid washing, processing temperature, and degree of surface oxidation. The VLS SiC whiskers are stable in Si3N4-SiOz-MgO matrix at temperatures lower than 1600~ [18], produce strong interracial bonds when processed between 1600 and 1750~ [109], and reveal large steps (about 10 nm) indicating strong dissolution at 1850~ [18]. The presence of an amorphous silica layer may reduce interracial reactions, thereby maintaining


smooth whisker surfaces even at 1850~ In SCW SiC whisker, an SiO 2 layer initially present on whisker surfaces prevents reactions during densification and results in formation of an amorphous interracial layer between whisker and matrix. The interracial bonding provided by this amorphous layer is weak enough to permit whisker debonding during fracture. The crystallinity of this 1- to 10-nm-thick layer can be affected by the composition of the oxynitride glass. Thus, a silica coating is effective only in specific systems.

The rough morphology of the reacted whisker surfaces can be reduced by the application of a 25- to 200-nm coating of carbon. In composites containing carbon-coated whiskers, the propagating crack runs along the matrix-whisker interface. This is because the graphitized carbon coating, with its laminar structure parallel to the whisker axis, debonds easily. Both reduced roughness and increased debond length contribute to an increase in fracture resistance of the composite. However, carbon coating has its own deficiencies. A layer that is too thick results in separation of carbon particles forming soft agglomerates. Nonuniform carbon layers produce whiskers that exhibit an inhomogeneous surface chemistry with local carbon enrichments (pockets). Auger and HREM analyses conducted by Kleebe et al. [59] show that use of this type of whisker results in nonuniform interfaces in which continuous amorphous glassy phases as well as a crystalline film can be present along the same whisker-matrix interface. The explanation of this phenomenon offered by the authors of [59] is based on the assumption that local carbon enrichments promote formation of SiO and CO, thus removing the SiO 2 constituents via CO and SiO vapor phases during early stage densification. The glassy phase formation along the whisker-matrix interface is reduced or suppressed. When the crack propagates and intersects the area where Si3N 4 and SiC are directly bonded together, the crack kinks into the whisker causing low fracture resistance in the material.

3. Si3N4-Y203-SiC w SYSTEM

The Si3N4-4 wt % Y203 was developed for high-temperature applications because of its ability to crystallize most of the glass phase. Five binary phases (SizN20 , Y28i207, Y48i3012, Y2SiO5, and YzSi303N4) and three ternary phases (Y10SiyOz3N4-H, YSiOzN-K , and Y4SizOyNz-J) can be formed [69]. The phase diagram of the Si3N4-Y203-SiO 2 system is shown in Fig. 24. High-resolution microscopy by Campbell et al. [13] reveals that an intercrystalline phase, presumably amorphous, is present as


Si02

. ~ Y28i207 / J / \'~ y4si3012

/ / ~Y,oSi7023N4 (a phase) ~',,\ / /~_.. .~SiO~N ~K phase) ~ ~ _ / ~ ' - / ~ 2 ( J phase)

Si3N 4 Y2Si303N4 Y203 FIG. 24. The phase diagram of Si3N4-Y203-SiO 2. After Lange et al. [69]. Reproduced

with permission of the American Ceramic Society, Westerville.

a thin continuous film at all Si3N 4 grain boundaries and at the interface between Si3N 4 grains and SiC whiskers. The predominant crystalline grain-boundary phase was found to be a -Y2Si20 7. Mitomo [81] reported formation of Y2SiaO3N4 and a liquid phase with composition similar to YSiO2N. The composition of the grain-boundary phase depends not only on the amount of Y203, but, critically, on the amount of surface silica. With 4 wt % SiO2, the overall composition lies along the SiaN4-Y2Si20 7 tie line; but some deviations can be expected because of the closeness of the tie lines in the SiaN 4 co rne r . Kleebe et al. [59] reported the fracture toughness of Si3N4-4 wt % Y203 to be 5.5 MPa m l/2b. Addition of SCW-Tateho SiC whiskers results in matrix-whisker debond lengths ranging from one to three times the whisker radius [13], with a slightly increased toughness. Addition of 30 vol % Tokai SiC whiskers increases Kic to 6.4 MPa m ~/2 CFT; and addition of 30 vol % American Matrix SiC reduces toughness to 4.8 MPa m 1/2 CFT [59]. Application of the above whiskers with carbon coating increases toughness to 6.0 and 5.7 MPa m 1/2 CFT, respectively. Lower toughness in the case of AM SiC is explained by Ca and C1 impurities present at the whisker surface as well as pockets of graphitized carbon. The reduction of the surface impurities and elimination of the extra graphite in the AM SiC resulted in materials with toughness of 6.5 MPa m 1/2 CFT [59]. The use of 20 vol % SC-9 ARCO SiC increases the toughness to 7.1 MPa m ~/2 CFT.

bMeasured by controlled flaw technique (CFT).


4. Si3N4-Y203-A1203-SiC SYSTEM

Microstructural characterization of Si3N4-6 wt % Y203-2 wt % A1203-SiC w composites show the presence of 1- to 2-nm grain-boundary layers between Si3N 4 grains and 5- to 10-nm layers separating SiC whiskers from Si3N 4 matrix. A similar glass composition (14 atom % A1 and 15 atom % Y)was found in the glass pockets independent of whether a grain boundary separates two Si3N 4 grains or a Si3N 4 grain and SiC whisker [11]. The Y203-m1203-SiO2 phase diagram is shown in Fig. 25 and the Si3Nn-Y203-A1203 phase diagram is shown in Section IV.B.5 in Fig. 33.

Experiments conducted with two types of whiskers, SC-9 SiC and SCW Tateho SiC, show preferential degradation of the SCW SiC. Degradation is attributed to the catalytic nitridation and oxidation of whiskers at 1400~ caused by the surface iron impurities [11]. The 20 vol % SC-9 SiC increased the toughness from 5.9 MPa m 1/2 to 6.5 MPa m 1/2 CNT. The use of 30 vol % of SC-9 SiC in the Si3Na-6Y203-1.5A1203 matrix increased the fracture toughness from 4.7 to 6.4 MPa m 1/2 CFT [12]. At 30 vol % SCW SiC, the toughness was reduced to 5.1 MPa m 1/2 CNT. The visible degradation of Tokai SiC whiskers in the Si3Na-Y203-AI203

SiO2

Two L,qu,ds~, . ' i ' , " / i - - ~

J" s s" s / , . ' , " ~,.'18 ~ \

Y O .3s,o v n . ~ n ./', i ,..d,,4, . . . . . . \ . . , ,' / -2,--3 ,-,,'.~2 f,, .' 'k"/ ~ - -" K. ; ,' .- \ ....... - \ " ," ,,, ,,.o, \

/ ! s176 . . . . \ \ .,.~900 "-..N, ' , ,; , \ Y203 T20 -~ 40~ 60' dO AI203

2Y203"AI203 3Y203"5AI203 Y203"AI203

FIG. 25. The phase diagram of the Si3N4-Y203-AI203 system. After Bondar et al. [10]. Reproduced with permission of Academy of Science, USSR.


matrix was also reported [35]. This degradation was significantly reduced when a presintering step was eliminated and the hot pressing time was shortened from 2 to 1 h. A fracture toughness approaching 6.8 MPa m 1/2 CNT was obtained for both Tateho and Tokai whiskers [97]. For Si3N4-6 % Y203-3 % A1203, the fracture toughness increased from 5 to 7.8 MPam 1/2 CFT in the presence of 20% SC-9 whiskers. A fracture toughness of 10.5 MPa m 1/2 CFT was reported for composites containing mixtures of SC-9 SiC whiskers and 10% large (40-/xm) particulates [63]. Surprisingly, a large increase in fracture toughness was also observed in Si3N4-SiC composites prepared from amorphous Si-C-N powder precur- sors [88]. Even though SiC whiskers were not used, the fracture toughness increased from about 5 to 7 MPa m 1/2 CFT. TEM observations indicated that the SiC particles were dispersed not only at the grain boundaries, but also within the Si3N 4 grains. The SiC dispersions (below 10 vol %) accelerated the growth of elongated Si3N 4 grains, thus improving toughness and strength.

5. Si3N4-MgO-SiC SYSTEM

In silicon nitride with 5 wt % MgO, initial liquid formation occurs at 1390~ [30]. The phase diagram proposed by Lange [66] shows only three secondary phases: forsterite (MgzSiO4) , enstatite (MgSiO3), and SizN20. Perera [91] investigated the whole quaternary system and shows the additional formation of Mg4N20, Mg4SiN 4, and MgSiN 2 with a tie line between MgzSiO 4 and MgSiN 2. On cooling, a magnesium silicon oxynitride glass is formed, which devitrifies on heating at 1340~ in nitrogen atmosphere to give forsterite, silicon nitride, and a trace of crystobalite. The glass composition lies within this three-phase region. Compositional analyses conducted by Davis et al. [18] showed that the interface phase contained Si and Mg in approximately a 5:3 ratio of oxygen to nitrogen. Experiments conducted by Shalek et al. [110] show relatively high fracture toughness for a Si3N4-5 % MgO matrix (7.1 MPa m 1/2 CNT). The addition of VLS Los Alamos SiC whiskers increased the toughness to 10.5 MPa m ~/2 CNT when hot-pressed at 1750~ and to 12.5 MPa m 1/2 when hot-pressed at 1850~ To the authors' best knowledge, these excellent results could not be consistently reproduced later, which suggests that, in this particular case, the matrix material could influence the toughness to a larger degree than the SiC whiskers. Indeed, Shalek et al. [110] describe the presence of highly fibrous/3-Si3N 4 grains. The Si3N4-5 % MgO material, studied by Schoenlein et al. [109] showed a K~c equal


to 5.8 MPa m 1/2 CNT. The addition of 10% of SCW Tateho SiC increased the toughness to 7.8 MPa m 1/2 CNT.

llI. Ceramics Toughened by in Situ Synthesized Whiskers

A. Introduction

In addition to the high manufacturing cost of whiskers, their inhalation poses a potential industrial hygiene concern. This concern, based primarily on indirect experience with asbestos, has significantly slowed the development of whisker-reinforced materials. Solutions to this perceived problem include either the use of larger whiskers (those beyond respiratory limits) or the development of in situ grown whiskers. The second approach is deserving of attention because it not only reduces contact with individual whiskers, and therefore the health concerns of whisker handling opera- tions, but can also reduce the cost of the entire process.

In theory, the combination of ceramic powders (A) with components (B~,B2, B 3) that form ceramic whiskers (B) at elevated temperatures should make it possible to produce ceramics with in situ whiskers. After the whisker formation is completed, the greenware is densified at a higher temperature. The disadvantages of this approach are: potential reactions between A, B1, B2, . . . , Bn; the difficulties in removing the metal catalysts required to grow whiskers; and the space constraints that limit the ability of the whiskers to form high aspect ratio crystals.

B. Si 3 N4-SiC Composites

1. PHASE RELATIONSHIPS IN THE S i - C - O - N SYSTEM

Even though several different families of composites can be produced by this approach, the Si3N4-SiC system is one of the best examples. This is because both Si3N 4 and SiC whiskers can be formed in the same S i - N - O - C system by control|ing processing conditions. The formation of silicon nitride and silicon oxynitride depends on the partial pressures of oxygen, SiO, and nitrogen. Experiments with nitrogen flowing through the mixture of silica and carbon have shown [127] that concurrent formation of SiC, Si3N4, and SizN20 may occur depending on gas flow rate, temperature, and C/Si ratio in the starting mixture. Since the gas phase

196 A L E K S A N D E R J. PYZIK AND ALAN M. H A R T

reaction is the route for Si3N 4 and SiC formation, the growth form of these phases is closely dependent on the partial pressures of N 2, SiO, and C O / C O 2. A stability diagram showing the relationship between the gas phase composition and the stable solid phases was constructed as a function of Po2 and (Psio/PN2) ratios and is shown in Fig. 26. Each boundary on this diagram represents the equilibrium Po2 and (Psio/PN2) of the corresponding two solid phases at 1350~ and carbon activity a c = 1. The Si3N 4 is stable with respect to SiC when the PSio/PN2 ratio is below line A. This equilibrium is controlled by CO and moves up with lower CO contents, while the (Psio/PN2) ratio decreases. Similarly, the equilibrium between Si3N 4 and S i z N 2 0 is determined by a specific PSiO/PN2 ratio at a given Po2" If the PSiO/PN2 ratio of gas phase is higher than the specific value for the given Po2, S i z N 2 0 is the stable phase, whereas Si3N 4 becomes stable if the Psio/PN: ratio is lower than this specific value [134]. To suppress the S i z N 2 0 formation, the Psio/PN2 ratio and the oxygen pressure must be low. The calculations of equilibrium conditions for S i C / S i z N 2 0 [133] show that PSiO/PN2 is constant as long as carbon activity a c remains constant, and increases with decreasing ac. Therefore, to suppress the S i z N 2 0 formation, the Psio/PN2 ratio should be higher than in the S i C / S i z N 2 0 equilibrium; but keeping a c higher is more effective to ensure the formation of only/3-SIC. A partial pressure of oxygen below 10 -2~ atm is sufficient to suppress the formation of

-2 z

t3..

.0 -4 co

12. ,,_4

__o -6

-10

SiC ([3) Si02 (C)

20

I I \

-25 -20 -15

log P 02 (atm)

FIG. 26. Phase stability diagram as a function of PO2 and PSiO/PN2 at T = 1350~ and a c = 1. After Wang and Wada [134]. Reproduced with permission of the American Ceramic Society, Westerville.


Si2N20. The carbon present in the system tends to maintain a partial pressure of oxygen below 10 -20 atm and virtually eliminates the formation of Si2N20.

Although the existence of the boundary temperature between the formation of SiC and Si3N 4 in the SiO2-C-N 2 system is well established (see earlier discussion in Section II.D.2), the boundary temperature varies depending on the nature of the reactants as well as the products. The porosity and size of the sample also affect this critical temperature due to the CO concentration produced by reaction. Thus, different boundary temperatures have been reported, ranging from 1374 to 1550~ [118, 128]. When the carbon-to-silica molar ratio exceeds 3, the boundary temperature is about 1450~ [71], above which silicon carbide forms and below which silicon nitride forms. When, however, the carbon-to-silica molar ratio is less than 3, silicon carbide forms more readily than silicon nitride, but reverts to silicon nitride (at temperatures ranging from about 1370 to 1550~ Under these conditions Si3N 4 and SiO form through the following two-step reactions even at temperatures higher than the boundary temperature:

SiO2 + 3 C , SiC + 2 CO (13)

2SiC + SiO2 + N2 , Si3N 4 + 2 CO (14)

SiC + 2SIO2 ,3SiO + CO (15)

Once SiC is formed, all three reactions can take place either simultaneously or sequentially such as reaction (14) after reaction (13) or reaction (15) after reaction (13) [71].

In the absence of impurities in the reactants, formation of/3-Si3N 4 and oxynitride can be completely suppressed. Experimental evidence [113] shows that when high-purity silica and carbon are reacted in nitrogen, the stable phase is a - S i 3 N 4 up to 1500~ Above that temperature, SiC is formed. However, in the presence of impurities, liquid formation takes place that alters the phase equilibria by changing the thermodynamic activity of the phases and altering the rates of the competing reactions. Iron additions to SiO2 �9 C powders promote SiC formation during nitriding at the expense of Si3N 4. Iron results in the formation of pure SiC at 1500~ in a SiO2 : C mixture, which, when nitrided in the absence of iron, gives pure o~-Si3N 4 [113].


2. FABRICATION METHODS

In the in situ synthesis of ceramic whiskers, the source materials required for whisker growth are combined with Si3N 4 powder. Often Si3N 4 itself becomes a source for growing SiC. In this method, raw materials are mixed together with densification additives and whisker-forming agents. The mixtures are shaped and subjected to a heat treatment during which whiskers are formed. In the following step, the temperature is increased to obtain densification. The experiments conducted by Siddiqi and Hendry [113] using 1 atm of nitrogen, a 4C:SiO 2 mix, and different amounts of FeC13 resulted in the formation of Si3N 4 whiskers in the absence of the iron catalyst and /3-SIC whiskers in presence of this catalyst. The Si3N 4 whiskers were grown by a VS reaction, where SiO formed by reduction of SiO 2 was further reduced by carbon in the presence of nitrogen to produce a-Si3N 4. The reaction products involving iron contained/3-SIC in decreasing proportion from 1500 to 1350~ High aspect ratio /3-SIC observed at 1500~ was formed by the VLS mechanism. The whiskers were covered with a thin layer of amorphous carbon.

Kamijo et al. [54] discovered that other metals and alloys consisting of Ni, Co, Cr, Ti, Ta, W, and Mo can accelerate whisker formation similarly to Fe. For SiC or Si3N 4 powders mixed with sintering additives, whisker formation agents containing Si, SiO2, C, and one of the metal catalysts, resulted in composites with SiC or Si3N 4 whiskers when heated between 1300 and 1750~

Besides the thermodynamic and kinetic requirements, the formation of whiskers depends on the homogeneous distribution of constituents and the availability of space for whiskers to grow in the shaped greenware. Both requirements are addressed in the chemical mixing process. In this method, Si3N 4 and the source materials for producing SiC whiskers (SiO2, C, NaC1, and CoC12 added as a catalyzer) are mixed in methyl alcohol. The NaC1 is used as a space-forming agent for the growth of whiskers because its boiling point is at the whisker formation temperature. This provides the required porosity in the Si3N 4 at the time of reaction [141]. The reaction product consists of SiC whiskers and a-Si3N 4 particles. The SiC forms from the SiO 2 :C mix as well as from Si3N4. The route of synthesizing SiC whiskers in situ by direct carbothermal reduction of silicon nitride with graphite eliminates the use of the catalyzer, thus minimizing the amount of impurities in the final product.

Conducting the reactions just discussed in an argon atmosphere produces high aspect rat io/3-SIC 5 to 40 ~m in length and 0.2 to 0.8 ~m in diameter. Formation of /3-SIC was initiated at 1400 to 1450~ Above ]650~ silicon was formed when carbon levels were deficient. The Si3N 4


could be converted completely to SiC with a molar ratio of Si3N4 :C = 1:3 at 1650~ [132].

IV. Self-Reinforced Ceramics

A. Introduction

In this review, self-reinforced ceramics are defined as materials or material systems having the ability to produce microstructures in situ with unique micromorphologies that allow toughening mechanisms to take place. Self-reinforced ceramics can be formed by (i) growth from the transition liquid phase (e.g., Si3N4); (ii) by decomposition of single solid solution (e.g., A1N-SiC); and by (iii) simultaneous crystallization of two solids from a eutectic composition liquid (eutectically solidified ceramics). In all cases, the initial composition (A) forms a liquid or solid solution (A1) from which phases (A2) with a morphology different than those existing in A and A 1 are grown.

In this section, three self-reinforced ceramics, Si3N4, A1N-SiC, and ZrC-ZrB 2, are discussed. Heating of a-Si3N 4 powder with sintering additives produces a supersaturated oxynitride glass from which ]3-Si3N 4 grows. With a proper glass composition and proper processing conditions, elongated /3 grains with high aspect ratios can be produced. Similar growth from the liquid phase has been observed in other ceramics, such as A1203 [114], A1N [27], and mullite [83]. Duplex microstructures of equiaxed and elongated grains characterized by clean grain boundaries can be produced by the annealing of solid solutions, such as A1N-SiC [72]. Finally, the large diversity of microstructures, ranging from plate-like to rod-like phases can be achieved in the eutectic solidification of ceramics. In this approach, the solidification process occurs simultaneously with crystallization of two solids from one liquid. Several oxide (AlzO3-TiO2, A1203-ZrO2, MgO-CaO) and nonoxide (SiC-BaC, ZrC-TiB2, TiC-TiB 2) eutectic systems can be fabricated in this way [116].

The preferred crack propagation along interface boundaries makes self-reinforced ceramics more flaw tolerant than monoliths. The other advantages are the availability of raw materials (the same as for monolithic ceramics), lower cost than fabricated whisker-reinforced composites, the elimination of a potential health hazard associated with whisker handling, and good control of phase morphology. The microstructures of these materials can be tailored to a specific application by controlling the nucleation, growth, and crystallization of the ceramic.


Having unquestionable advantages, the self-reinforced ceramics cannot provide the solution to all of the problems facing ceramics. The self-reinforced ceramics are limited to specific systems and the improvement in fracture toughness ranges typically from 1.5 to 2 times, as compared to monoliths. This change greatly improves the material's mechanical response, but is not always sufficient to provide long life reliability. Perhaps the best potential for commercialization is in ceramics such as Si3N4,/kiN, or A120 3 where rod- or plate-like grains are grown from the liquid phase. However, even here, a better understanding of the phase relations and their interactions with the resulting microstructures is required before one can take full advantage of the microstructure designing capabilities characteristic of these materials.

B. Silicon Nitride

1. THE a- -~ TRANSFORMATION AND GRAIN GROWTH

IN SILICON NITRIDE CERAMICS

Si3N 4 exists in two crystallographic forms, a and/3, as was discussed in Section II. At high temperatures a transforms into/3. The significance of this transformation is in the rod-like morphology of the /3 phase. High aspect ratio silicon nitride grains act similarly to SiC or Si3N 4 whiskers in deflecting cracks.

Even though the tendency of Si3N 4 to grow long hexagonal grains has been observed for a long time [19], the control of this process toward the maximization of the fracture toughness is relatively recent. In 1979, Lange [67] reported an improvement in flexure strength and in the critical stress intensity factor K c (up to 6 MPa m 1/2) when /3-Si3N 4 grains with high aspect ratio were formed during sintering. Wotting et al. [140] produced Si3N 4 + 15% Y203 + 3.4% A120 3 having Kic of 8.2 MPa m 1/2c and a strength of 750 MPa. Tani et al. [119] obtained elongated silicon nitride grains in Si3Nn-Y203-AI203, Si3Na-La203, and Si3Na-CeO 2 systems having a toughness of about 6 MPa m 1/2 CNT. Matsuhiro and Takahashi [77] have reported fracture toughnesses of 9.7 MPa m 1/2 CFT and flexure strengths of 900 MPa; Li and Yamanis [73] reported a fracture toughness of 10.6 MPa m 1/2 (double cantilever beam) and a strength of 550 MPa; Pyzik et al. [99] produced silicon nitride with fracture toughnesses of 8 to 11 MPa m 1/2 CNT and strengths of 900 to 1250 MPa.

C Measured by microindentation technique (MT).


The growth of elongated grains in silicon nitride ceramics is a complex phenomenon in which the mechanism responsible is still not fully understood. However, experimental observations indicate that the grain morphology is affected by (i) the crystallographic modification, size, and surface chemistry of the silicon nitride powder; (ii) the chemistry of oxynitride glass formed at high temperatures; and (iii) the processing conditions.

2. THE EFFECT OF SILICON NITRIDE POWDERS

ON GRAIN MORPHOLOGY

In a-rich powders, the solution of material occurs at the a-Si3N 4 contact points. As a result of a concentration gradient, diffusion of Si and N takes place throughout liquid phase, followed by precipitation of /3 phase in the form of small 13 nuclei and on already existing/3 grains. The /3 phase is a more thermodynamically stable phase, but the magnitude of the difference is small (only 1.3% of the absolute value) [85]. Jennings [52] suggested that the formation of a-SiaN 4 requires molecular nitrogen, whereas fl-SiaN 4 is formed from atomic nitrogen. Dissolution of a-SiaN 4 leads to the formation of atomic nitrogen and therefore to the precipitation of fl-SiaN 4. With little or no fl in the starting powder, a certain supersaturation has to be reached before nucleation takes place. In fl silicon nitride powders, dissolution of small fl grains and precipitation of material on the large grains takes place instead. In the first stage, grain growth is mainly interface controlled and grains grow faster in the c direction than in the width. In the following stage, the driving force for grain growth is the reduction of the free surface energy, which results in a decrease of the aspect ratio. The effect of a-/~ transformation on microstructure development is shown schematically in Fig. 27.

It has been observed that the amount of/3 phase in the starting silicon nitride powder and the morphology of this phase significantly affects the microstructure of the SiaN 4 part. Lange [68] has reported a correlation between high aspect ratio grains in dense material and high-a content in the starting powder. Knoch and Gazza [61] showed that high-/5 silicon nitride powder has coarser and more equiaxed grains than the low-fl powder. Hoffmann et al. [40] observed that high aspect ratio /3 grains grow from/3 needles present in silicon nitride powder and, in fact, limited amounts of /3 particles with small diameters and high aspect ratios are desirable. Natansohn and Sarin [85] succeeded in fabricating acicular grains from 100% fl-Si3N 4 powder. The aspect ratio was, however, smaller compared to those produced from a rich powders.


�9 TIME I I S~176176 Dissolution and CONSOLIDATION Particle Rearrangement i (L iquid) ~ Precipitation Coalescence i Precipitation of 13

STARTING "LTquid Phase Formation 'll ~-13 Phase Transformation 11 Complete ~-13 Preferential " POWDER I Nucleation I Transformant Growth of 13

I (x and I 1 ~ ~ ~ f Grov~t I h

I I

Dissolution of oL ~m Precipitation .. Ostwald , , , 0 - . , . T 0e o~ + 13 ~ ' Ripening

. . . . . ' __Precipi tat ion on ....~ulssoluuon , ~ . - . of S m a l l e r f Large 13 Grains

13 13 Grains ~ ' , . t - - - I J ! I_

I Ltquld Encapsulated-Planar MICROSTRUCTURE Solubility of(, and or | Dissolut, on Rate ofl5 I AIII~-St3N4 | Growth ofl5 Grams 13 Jn Ltquid Phase. I Proporhonal to Super I Complete Conversion | Dependent on Amount. I Saturahon Limits of i Dependent on Continuous ComposttJon. and Viscosity I ,, and or IS m Liquid | Availabthty of Liquid Flux | of Liquid Phase [Temp.} I I I

L__ CONCURRENT STAGES ..._l

F[o. 27. Outline of Si3N 4 consolidation mechanisms. After Natanson and Sarin [85]. Reproduced with permission of the Deutche Keramische Gesellschaft E. V., Koln.

Besides the amount and morphology of /3 grains in the silicon nitride powder (commercial powders have typically 3% to 10% /3), the size and the distribution of a particles plays a critical role. Mitomo et al. [82] concluded that the large difference in solubility between grains of different sizes causes the grain size distribution in starting powders to determine the microstructure of the sintered material. Since the a-/3 transformation occurs by a liquid-phase mechanism, the surface chemistry and impurities greatly affect the temperature of the liquid formation, rate of dissolution, and, as a result, the a-/3 transformation. Wada et al. [129] have shown that the transformation of low oxygen content ce-Si3N 4 prepared by a CVD technique takes place at 2100~ whereas in the presence of surface silica, the transformation starts at about 1700~ [85]. In the presence of a liquid phase formed by silica and sintering additives (especially when powders have impurities of low-melting-point metal oxides), the formation of the /3 phase may start as low as 1350~ In silicon nitride powder containing 3% sintering additives (based on the glass composition discussed in Section IV.B.6) the transformation is completed at about 1750~ Elongated silicon nitride grains with diameters ranging from 0.1 to 0.7/zm and lengths approaching 50/zm are shown in Fig. 28a. The same silicon nitride containing, instead of Y203-MgO-CaO glass, only SiO 2 and heat-treated at identical conditions produces the grain morphology illustrated in Fig. 28b.

5 WHISKERS AND W H I S K E R - R E I N F O R C E D CERAMICS 203

FIG. 28. Morphology of "-Si3N 4 grains after heat treatment at 1750~ for 12 h depend on composition of glass phase. (Top) With Y203-MgO-CaO-based glass. (Bottom) With SiO2-based glass.


3. EFFECT OF GLASS CHEMISTRY ON GRAIN MORPHOLOGY

The a-13 transformation, nucleation, and growth of silicon nitride grains depend in large degree on the liquid phase. The chemistry of the liquid phase, along with the characteristics of the starting powders and processing conditions, are the key factors in determining the microstructure and properties of silicon nitride parts.

The observation that a high-viscosity liquid produces high aspect ratio grains, whereas the formation of equiaxed grains takes place in a low-viscosity glass, was first made by Wotting and Ziegler [137] and Hampshire et al. [32]. These authors reported that the higher the viscosity of the liquid phase, the longer the period of supersaturation, with the result that elongated 13 grains are formed. In the case of low viscosities, the diffusion is rapid and a reduction in the surface energy becomes the main driving force, resulting in microstructures with equiaxed grains. Hwang and Tien [46] explained the anisotropy of Si3N 4 grain growth by the differences in the structure of the solid-liquid interface and the local solute diffusivity of the liquid at the growth fronts in the two directions. The proposed model, based on Jackson's crystal growth model, explains the fast growth at the rough growth front in the [001] direction and the slow growth at the smooth fronts in the [210] directions. The model also assumes the existence of the higher solute diffusivity (according to the Einstein-Stokes relationships, a higher solute diffusivity implies low liquid viscosity) in liquid present in the multigrain junctions than in the two-grain boundary. Due to the prismatic configuration of /3 grains, the growth in the c direction is controlled mainly by the solute diffusion through the multigrain junctions, while growth in the [210] direction is controlled by the diffusion along grain boundaries. As a result, the growth rate is slower in the [210] direction. Kang et al. [55] based their explanation of SiaN 4 anisotropic growth on the fact that if local supersaturation is sufficiently high, the reaction kinetics at the edge regions can be changed from reaction controlled to diffusion controlled, resulting in a dramatically different growth rate.

Pyzik and Beaman [98] showed the effect of glass chemistry on the growth rate and the microstructure. This correlation, based on experimental observations, is schematically illustrated in Fig. 29. The growth rate increases with supersaturation and then decreases as the viscosity increases. In low saturated glass where the rate of nucleation, rate of grain growth, and viscosity are presumably low, the resulting microstructure consisted of large grains. The formation of the large prismatic grains in the SiaN4-MgO system has been observed also by Knoch and Gazza [61], Knoch and Ziegler [62], and Heinrich et al. [38]. With increase in


*****,,,,**"" . . . . . . . ~ / [02 ' " / /

< r

I

SUPERSATURATION FIG. 29. Schematic illustration of the correlation between glass chemistry, grain growth,

and microstructure. After Pyzik and Beaman [98]. Reproduced with permission of the American Ceramic Society, Westerville.

saturation and viscosity, the growth rate decreases and small grains with high aspect ratios are formed. Several oxides, including Y203, Nd203, Sm203, Gd203, etc., can be used to produce this type of microstructure [39,92,107].

The best results in controlling grain morphology were reported for multicomponent systems, such as S i 3 N 4 - Y 2 0 3 - A I 2 0 3 , S i 3 N 4 - Y 2 0 3 - M g O , and SiaN4-Nd2Oa-MgO. Recently, a new family of self-reinforced silicon nitrides has been developed based on the Si3N4-Y203-MgO-CaO system where the yttria acts as a conversion aid, magnesia as a densification aid, and calcia as a whisker growth-enhancing agent [99]. It was found that Y can be replaced by 7 other elements, Mg by 6 other elements, and Ca by 19. These compositions yield fine-grained, high aspect ratio silicon nitride ceramics with a unique combination of high flexure strength and fracture toughness [ 101].

4. EFFECT OF PROCESSING CONDITIONS ON MORPHOLOGY

OF SILICON NITRIDE GRAINS

Higher processing temperatures and/or longer processing times lead to formation of large grains. The effect of grain growth on properties will vary, however, depending on the growth mechanisms. In a Si3N4-Y203-A1203 system, the maximum strength and toughness are


achieved when the a-/3 transformation is just completed [138]. Longer processing times lead to grain coarsening and result in a reduction of propert ies. In contrast, the processing window in the Si3N4-YzO3-MgO-CaO system is much broader and grain growth in the c direction is observed long after transformation is finished [15]. The latter case gives more processing freedom for achieving microstructures tailored for needs of specific applications. In the Si3N4-SiOz-MgO system, the aspect ratio increases with increasing temperature from 1750 to 1850~ According to Richter et al. [106], who studied phase equilibria in the S i -Mg-N-O system, the increase is caused by the higher solubility of Si3N4 in the liquid phase at higher temperatures. At 1750~ the solubility is 42.4 wt %, and at 1850~ it is 52.2 wt %, for the molar ratio MgO/SiO 2 = 2. The viscosity, which is increased by higher concentration of Si3N 4 in the melt and decreased by higher temperatures, remains approximately constant, but the total amount of liquid phase increases.

Fabrication of dense structures requires the kinetics of densification to be faster than the kinetics of grain growth. This, however, is often associated with formation of equiaxed grains. The process of self-reinforcing is highly complex due to many variables affecting silicon nitride grain morphology and silicon nitride-glass interfaces. Even though the number of elongated grains is always present in any Si3N 4 ceramic, the low-temperature formation of highly fibrous structures is, so far, limited to a relatively few systems.

5. Si3N4-Y203-A1203 SYSTEM

Even though high aspect ratio /3-Si3N 4 grains are observed in silicon nitride ceramics densified with yttria, the fracture toughness of these materials is typically below 6 MPa m 1/2 CNT. This is due to the nature of the interfaces and the characteristics of the secondary phases formed in this system. The addition of A1203 results in improvement of the density and the room-temperature mechanical properties. Phase relationships and liquidus temperatures in the Y203-A1203-SiO2 system are shown in Fig. 30 [139]. Line A1-A 5 represents a composition of Si3N 4 with 15 wt % Y203 and different amounts of A1203. As the amount of A1203 increases, the composition shifts through several compatibility triangles characterized by decreasing liquidus temperatures. The density increases, but the aspect ratio decreases. With an increase in the additive concentration (line El-E3) , the secondary phase consists of more Y203 . According to Wotting and Ziegler [139], these compositional changes result in higher liquidus temperatures and viscosities at the sintering temperatures. This leads to


SiO2

Y203.2SIO2 2Y203-3SIO2 ,,z ," o-~-o o,., ,~u ,,., ~ " 3AI2032SIO2

Y203 SiO2 ' 1385oc

/ Y203 2Y~Oa'AI2Oa Y203"AI2Oa 3Y2Oa'5AI2Oa AI2Oa

FIG. 30. Phase relationships and liquidus temperatures in the Y203-SiO2-A1203 system. After Wotting and Zegler [139]. Reproduced with permission of Elsevier Science Publishing Co., New York. Dashed line, added by authors, shows reported in the literature compositional range at which the presence of elongated silicon nitride grains has been observed. The open circles show compositions where the authors have observed elongated silicon nitride grains.

the growth of crystals with higher aspect ratios, as indicated by t h e microstructural characteristics in Fig. 31a and b. The dashed line in Fig. 30 shows the compositional range of various Y203-AlzO3-SiOz-based silicon nitrides where the presence of elongated grains has been reported. The open points indicate compositions of some commercial materials. Most of them are characterized by a fracture toughness between 5.5 and 7 MPa m 1/2 CNT. The pressureless sintered compositions produced by Wotting and Ziegler contained 15 wt % Y203 plus 3.4 wt % A120 3 (E3) and had a fracture toughness of 7.8 MPa m 1/2 MT. Compositions containing less glass (5% Y203 + 1.13% A120 3) and densified by a sinter-hot isostatic pressing technique achieved a fracture toughness of 8.2 MPa m 1/2 MT. Similar compositions (5% Y203-2% A120 3) were investigated by Mitomo et al. [82]. These materials heated in 980 kPa of nitrogen achieved the highest aspect ratio (4.2) at 1950~ The fracture toughness of materials fabricated from a and /3 powders by heating at 1950~ for 1 h ranged from 5.8 MPa rn 1/2 to 6.3 MPa m 1/2, respectively. Both toughnesses were measured by a single-edge-precracked beam method. Hecht et al. [37]


a lO0 ,..-.,

e~

if) z m 90 121 _ J < _o I--- W n- O LLI -1- 80 I---

b loo-

o~ v

>..

r z LU 90 r .._1 < _o F - U.I n -

O u J -r 80 I--

10

, 9

F- ~7 . ...I 8 6 rr

5

i";

z ~

8 r r

P . . . + , - - " - - ~,,..x + ~ - -,. - - " " I , + , , , , , , = , , , = ' '

, , x f a

,.....K:),.---.-. . ' ' ' ' ' ' o "

_ r t _ - - ~ 1 | i i~iiiii!iiiiiiiiiiii~|- dmin ~

I--1[ I I Y203 5.0 10.0 15.0 AI203 + 1.13 +2.26 +3.40

A D D I T I V E C O N C E N T R A T I O N (wt. %)

- 9

- 8'~, It 'tl O

_ 7 t...- < rr" F--

- 6 ~ I I I

< - 5

A1 A2 A3 A,, As

0 r r ~ r ~1o

..if x 9 ~ x . . . . ,.+ . . . . . . . . . . . . ~ 9 . . . . . . . . . , . - .r v . . . . . . - ~ , , ~ �9 =-

8 I"~1 - - 8 x a ir

i i 51"- I ~ L ~ ~ - 5 <

T I It m,~ I I I I 0 1 2 3 4 5 6 7

AI203 C O N C E N T R A T I O N (wt. %)

FIG. 31. Microstructural parameters and density of silicon nitride sintered at 1800~ for 2 h + 1850~ for 2 h as a function of (a) the additive concentration and (b) the amount of AI20 3 added to Si3N 4 + 15 % Y203. After Wotting and Ziegler [139]. Reproduced with permission of Elsevier Science Publishing Co., New York.

measured the Kic of several commercial silicon nitride powders containing Y203 and A1203 additions using both controlled flaw and microindentation techniques. Data showed up to a 60% difference, giving higher numbers with the microindentation method. Neil et al. [86] reported 5.5 MPa m ]/2 CNT for a hot-pressed composition containing 6% Y203 + 1.5% A1203 and characterized by the presence of large elongated grains.

The phase diagram of the Si3N4-Y203-A1203 system shown in Fig. 32 illustrates that even a small change in the Y203-to-A1203 ratio can easily


Y203

MOL %

SiaN4 AI2Oa

FIC. 32. The Si3N4-Y203-A120 3 behavior diagram showing glass formation after heating at ]800 to ]900~ After Jack [50]. Reproduced with permission of Plenum Press, New York.

shift the composition from one compatibility triangle to another. Figure 32 shows the presence of a wide compositional range expanding with increasing temperature, where the cooled products are /3'-SiAION and Y-SiAION glass of a composition near Y10Si12AI17057N5 [51]. Small changes in the starting composition and variations in processing temperature will result in different products. Therefore, comparison of toughness data obtained for similar, but different, systems obtained under different processing conditions and measured by different techniques can only be used to indicate general trends.

Material properties in the Si3Na-Y203-A1203 system are determined by the amount of Si3N 4 (SiO2), the Y203/AI203 ratio, and the existence of crystalline phases, such as Y4A1209 (2: 1)-YAM, YAIO 3 (1 : 1)-YAP, Y3A15O12 (3:5)-YAG, and the nitrogen-containing phase N-melilite (Y2SizO3N4). YAM is stable over a wide range of temperatures and melts without decomposition at 2020~ YAM is isostructural and forms a complete solid solution with the yttrium silicon oxynitride N-YAM (Y4SizOyN2; also called J phase) [30]. The orthoaluminate (YAP) is a metastable compound that has two forms. The low-temperature form has the same structure as YSiNO 2, but is only stable up to about ll00~ The high-temperature form is stable above 1825~ [14]. Having a simpler structure, YAIO 3 is energetically more favorable to form than Y3A15O12 . The formation of the Y3AIsO~2 with garnet structure is possible at 900~


Si ,N4

~ " SiO2 \

/J_ H, .,'e~ Si;b'~2- ", N \ N ~,p

'a YAM YAP YAG AI203

FIG. 33. Tentative Si3N4-Y203-AI203-SiO 2 diagram.

YAG melts without decomposition at 1950~ [10]. A certain amount of A120 3 can enter the crystal structure of SizN20 , forming a partial solid solution, the so-called O'-sialon. The limiting solubility of A120 3 in SizN20 solid solution was determined to be about 15 mole % [14]. The fact that SiO 2 is always present in the AlzO3-Y203-Si3N 4 system introduces several additional phases. The "H" phase accompanies the formation of the yttrium silicon oxynitride and has the appatite structure. This nitrogen-yttrium appatite is represented by the general formula Y5 Si3 OlzN; but it has an appreciable homogeneity range near (Y4Si)(Si3OllN)N [49]. The locations of 10Y20 3 9SiO 2 Si3N 4 (H), 2Y20 3 SizN20 (J), Y203 SizN20 (K), Y203 Si3N 4 (M), 3Y20 3 AlzO 3 28i2N20 (N), Y203 28iO2, 2Y203 3SiO 2, Y203 SiO2, 2Y203 A1203 (YAM), Y203 A1203 (YAP), 3Y203 5A1203 (YAG), and 3A1203 2SiO 2 (mullite) in the Si3N4-SiO2-Y203-A1203 phase diagram are presented in Fig. 33. The dotted line in the Y203-A1203-SiO2 plane represents the approximate range of glass compositions resulting in silicon nitride ceramics with elongated grains and improved fracture toughness.

6. Si3N4-Y203-MgO SYSTEM

In Si3N 4 with only Y203 additions, the a-/3 transformation starts at low temperatures. Since diffusion through the highly viscous liquid is slow, the


transformation occurs with little transport of material and hence with little densification [30]. With a MgO addition, the transformation is slower, but rapid transport of material can occur. Because of these characteristics, both additives give markedly different microstructures in terms of silicon nitride grain morphology. Mixed additions of both MgO and Y203 allow one to combine a high rate of densification and microstructures with the desired grain size [25]. Hampshire and Pomeroy [31] investigated the microstructural development in Si3N 4 with 10% Y203 and MgO mixed in various ratios. The maximum aspect ratio of/3 grains was obtained at a Y203-to-MgO weight ratio close to 6: 1. Below the ratio of 2.3 : 1, reduction of the aspect ratio was observed. The fracture toughness measured for the materials from the above range was between 6 and 6.7 MPa m 1/2 CNT [86,98].

The addition of small amounts of several metal oxides to the YzO3-MgO glass enhances the growth of elongated grains and allows changes in the silicon nitride-glass interface. The fracture toughness is increased, depending on composition, from about 8 to 11 MPa m 1/2 CNT[IO1]. When CaO is used as a grain-enhancing compound, the maximum aspect ratio is obtained at about 2:1 Y203-to-MgO ratio (in wt %). In materials containing 14.7% glass, the fracture surface contains about 60% glass. Auger analysis of in situ fractured surfaces showed that the Ca content is 2.2 to 4 times higher than the average Ca content of the glass phase, suggesting preferential crack propagation through Ca-rich regions. Thus, not only does Ca promote the formation of high aspect ratio whiskers, the presence of Ca in the glass effects the crack propagation. Both influence the properties of the serf-reinforced Si3N4-YzO3-MgO-CaO ceramic.

The glass phase in the Si3N4-YzO3-MgO-CaO ceramic contains yttrium, magnesium, calcium, silicon, and nitrogen. The amount of nitrogen increases toward the higher content of Si3N 4 and the higher SiO2/MgO + Y203 ratio. Assuming that nitrogen is always present in the glass and the amount of catalyst is very small, the simplified glass composition can be expressed by the YzO3-MgO-SiO2 phase diagram. The compositional range (in wt %) required to improve fracture toughness is presented in Fig. 34. The replacement of CaO with other elements often shifts the compositional range toward the left side of the diagram, beyond a 4:1 ratio. For 0.2% CaO, the aspect ratio decreases outside the range indicated on line A by points A 0 and A 4 (for low SiO 2 content) and on line B by B 0 and B 4 (for high SiO 2 content). Line C represents compositions with the same YzO3-to-MgO ratio, but different SiO 2 contents, varied through the amount of silicon nitride. The aspect ratio (a95) , average diameter (d), and number of grains in a 100-/xm 2 area for these compositions is shown in Fig. 35. Between 4% glass (C 5) and 15% glass


SiO2

WT%

MgO.SiO2

Y~O3"2SiO2

2MgO-SiO2

Y203"SiO2

Y203 4.t 3MgO'Y203 1= MgO

F]c. 34. The effect of glass composition in the Si3N4-Y203-MgO-SiO2-CaO system on silicon nitride aspect ratio. Open circles indicate glass composition producing low aspect ratio grains. Solid circles represent compositions resulting in silicon nitride with a high aspect ratio.

(Co), the average length of/3 grains is similar, but the average diameter decreases from 0.35 to 0.21 /~m. As the amount of Si3N 4 increases, the number of 13 grains per 100/zm 2 increases from 250 to 410. The aspect ratio increases from 7.5 for the C o composition to 11 for the composition C 5. The fracture toughness measured for compositions C O to C 5 ranges from 10 to 10.8 MPa m ]/2 CNT. The microstructures of SiaN4-Y2Oa-MgO-SiOE-CaO materials containing 4% glass and 15% glass are shown in Fig. 36.

The compositions represented by points 0 to 5 on line C (Fig. 34) are positioned at the different compatibility triangles of the SiO2-MgO-Y20 3 system. If no postdensification heat treatment is applied, all of these materials consist mainly of/3 silicon nitride and amorphous glass. Upon heat treatment, the formation of new phases increases from C 5 (least) to C O (most). Glass in silica-rich compositions remains mostly amorphous


/ 11 0.80 i- / g 11 070~

-"m ~.0.60 I" 400 o o +! (o.~oi_I m + +o+o I +oo+0 '+ +o.3+p ] ":'+,

] z ~,,?, c,, c~ , c~ ?, , Co , r

4 6 8 10 12 14 16

GLASS CONTENT, wt %

FIG. 35. Microstructural parameters in the hot-pressed Si3N4-Y203-MgO-SiO2-CaO material as a function of glass content. After Pyzik and Beaman [98]. Reproduced with permission of the American Ceramic Society, Westerville.

even after 50 h of the heat treatment [47]. Crystallization of C 0, C2, and C 5 compositions produces several phases, but only two, MgsY6SisO24 (G) and MgSiO 3, without nitrogen. Below 1350~ the major phases are MgYSi2NO 5 (N), (Y,Mg)sSi3(NO)13 (A), and phase G. This is in good agreement with Patel's and Thompson's [90] work showing the presence of N, G, and A phases at 1250~ Above 1450~ the formation of phase A, Y2SiNO 2 (J), YSiNO2(K), and Si2N20 was observed [47]. All of these phases, except (Y,Mg)sSi3(NO)t3 , are located in the Y203-MgSiO3-Si2N20 plane of the Si3Nn-Y203-MgO-SiO 2 phase diagram. The tentative locations of the forming phases and range of glass compositions are shown in Fig. 37.

7. Si3N4-MgO-Nd203 SYSTEM

With Y203-MgO additives, the range of compositions giving both good sintering behavior and a beneficial microstructure is limited. With MgO-Nd20 3 additives, the processing window is much broader. The

214 A L E K S A N D E R J. PYZIK AND A L A N M. H A R T

FIG. 36. The microstructures of Si3N4-Y203-MgO-SiO2-CaO ceramics with (a) 4% of glass and (b) 15% of glass. After Pyzik and Beaman [98]. Reproduced with permission of the American Ceramic Society, Westerville.


Si.~N4

MOL % D

M

K

. . . . . . . H t t "~

Y203 Y2dga06 MgO

FIG. 37. The tentative locations of crystallization products in the Si3N4-Y203-MgO- SiO z system and range of glass compositions (dashed line) used in crystallization experiments.

MgO-SiO2-Nd20 3 ternary system has five binary phases and one ternary phase as shown in Fig. 38. With the 4:1 molar ratio of MgO to Nd20 3, a trace of Nd-N-mel i l i te (NdzSi3OlzN) is detected, while with 1:1 molar ratio, M g - N d - N apatite is observed [34]. The presence of Nd-N-mel i l i te as a grain-boundary phase is not desirable. However, additional heat treatment above 1075~ leads to the formation of the apatite. The Nd-N-apat i te , existing as a solid solution over a range of MgNdaSi3013 to NdsSi3OlzN , is a beneficial secondary phase because it can be crystallized into a wide range of compositions with different MgO : Nd20 3 ratios [49]. At high temperatures, the NdsSiO12N phase oxidizes to produce X1-NdzSiO 5 and Nd2Si20 7 phases with very small specific volume changes [33].

The materials with 10 mole % of M g O / N d 2 0 3 in the 1 : 1 ratio have an aspect ratio of silicon nitride grains equal to 3.3. With the same molar ratio, but a 17 mole % additive level, the aspect ratios are about 6. With the same total additive level (17 mole %), but a higher molar ratio (4: 1), grain size is similar (about 0.4 ~m), but the aspect ratio is slightly lower at 5.4 [34].

216 A L E K S A N D E R J. PYZIK AND A L A N M. H A R T

Nd2Oa / k

LINE OF CONSTANT / ~ , , ,ooc

N d 2 S i 2 O T ~ MgaNd206

~0.~.~,~o,~- / . . - ~ \ \ . . . . . . . 7.--\

Si02 MgSi03 Mg2Si04 MgO

FIG. 38. The Mg-Nd-Si-O phase diagram showing compositional lines corresponding to mole ratios of 4 and 1 used as additive compositions to silicon nitride containing SiO 2. After Hampshire et al. [34]. Reproduced with permission of the American Ceramic Society, Westerville.

2800 2600

2400

= 2200

2000 E

1800

1600 i ' . .~ . , ' , ~13,c l

14000 10 20 SiC

wt. % AIN 0 10 20 30 40 50 60 70 80 90 100

i--~i~..~_ . I , , I I I I I I I I

5H~ ..... ~-"--:'-""""---'-----'''"""-'-'---g.'.'S 4H ~H+~

) / 8(2H) solid solution

o / t o

so S . ts~,..,.- o . . . . . . . . . . . . . . . . . . . . . . . . ~-" .... ,,- " 81 + 82 "'"',,

�9

�9 "~,," metastable 13 ", \ - - , , , , , 304O 50 60 70 80 90 100

mol. % AIN AIN

FIG. 39. Tentative SiC-AIN phase diagram. After Zangvil and Ruh [142]. Reproduced with permission of the American Ceramic Society, Westerville.


C. AIN-SiC

1. PHASE EQUILIBRIA

The AIN-SiC alloys can achieve a single-phase solid solution by heat treatment at a solid solution range, as shown in Fig. 39. A single solid solution, 3, with the 2H structure extends, according to work by Zangvil and Ruh [142], from about 23% A1N (at 2100~ to 100% AIN. Below 23% A1N, a 3 + 4H two-phase field exists. When SiC-A1N compositions are heated within the single solid solution range, equiaxed grains with a 2H structure are formed. The lattice constants of solid solution vary linearly as the SiC/A1N ratio changes. Below 1900~ a miscibility gap exists, which is shown in a tentative phase diagram by a double dotted line. The miscibility gap was first suggested by Rafaniello et al. [103] who conducted annealing experiments of solid solutions obtained by a carbothermal reduction reaction. They found that the solid solutions, which were obtained at low temperatures, were in fact metastable and tended to separate into SiC-rich and A1N-rich phases, denoted in the Zangvil and Ruh diagram as 61 and 62.

2. MICROSTRUCTURE EVOLUTION

The significance of this phase separation is in the morphological changes associated with this separation. Lee [72] found that solid solution treated specimens, after annealing at 1860~ for 96 h, show duplex microstructures consisting of elongated grains and modulated grains within the equiaxed grains. The amount of elongated grains increased with higher SiC contents. The equiaxed grains were found to have a 2H structure and contained more/kiN. The elongated grains had other polytypes, such as 6H and 3C, that contained less than 10% A1N and were heavily faulted. Suzuki [117] produced sintered silicon carbide articles composed of elongated grains of a SiC-A1N solid solution consisting of 2 to 20% A1, from 0.2 to 10% N, from 0.2 to 5% O, from 0 to 15% of Group Ilia element, and the rest being Si and C (all in wt %). The highest aspect ratio was obtained with 70% to 97% of /3-SIC and an average particle size for all components below 1/zm.

Pure/kiN has a fracture toughness of 3.4 MPa m 1/2 MT and SiC has a fracture toughness of 3 MPa m 1/2 MT. The SiC-A1N ceramic alloys, after solid solution treatment, have a toughness of about 3.5 MPa m 1/2 MT. After the alloys were annealed in the miscibility region, the fracture


toughness of the SiC-rich alloys was increased to about 6 MPa m 1/2 MT. The fracture toughness of the alloys rich in A1N showed no increase. The alloys containing higher SiC content were found to have more elongated grains and a higher toughness [72].

D. Directly Solidified Eutectic Ceramics

1. EUTECTIC GROWTH

The eutectically solidified ceramics represent a large group of ceramic systems with microstructures that can be tailored during cooling. Since the eutectic solidification is based on the simultaneous crystallization of two solids from one liquid, many complex arrangements of the eutectic phases have been observed [ 93]. A schematic representation of a eutectic reaction is as follows:

eutectic liquid , a solid solution +/~ solid solution (16)

The eutectic growth takes place in the coupled zone region, which is determined by the temperature and composition within which lamellar or fibrous eutectic crystallization occurs at a greater velocity than the crystallization of either primary phase. The shape and position of the coupled growth region are governed by the shape of the liquidus lines of the phase diagrams and the relative crystallization velocities of the pure constituents [116]. Examples of a symmetrical phase diagram and coupled zone associated with a normal structure and an asymmetrical phase diagram and coupled zone associated with an anomalous structure are shown in Fig. 40 [23]. The essential feature of a normal structure is that the phases crystallize simultaneously by the advance of a common macrointerface into the melt. In an anomalous eutectic, the two phases do not grow at equal velocities to form a common crystallization interface. A mixture of the two phases is made possible by repeated nucleation of one or both of the phases [ 116].

The extensive morphology studies conducted by Smith and his co-workers [23] enabled him to classify eutectic microstructures in terms of the parameters that influence the growth process. Some of the main features of this classification scheme are illustrated in Fig. 41, which shows the influence of solution entropy and volume fraction on the microstructure at a growth velocity of 5 • 10 -4 cm/sec. The vertical line at h S a - 23 J /mol K -~ divides normal and anomalous structures. Whisker-like


(a)

a n o / an~ eutectic eutectic I.-

rr . . . . . . . / . - w (b) 13..

U.l F-

. . . . . . .

A and. ~ eutectic eutectic ~

A COMPOSITION B

FIG. 40. Phase diagrams for normal and anomalous structures. (a) Symmetrical phase diagram and coupled zone associated with a normal structure. (b) Asymmetrical phase diagram and coupled zone associated with an anomalous structure. After Elliot [23]. Reproduced with permission of the ASM International, Metals Park.

structures, regular rods (region b), and irregular fibers (region g) form at low volume fractions and at low solution entropy.

2. FABRICATION METHODS

If an eutectic composition is solidified using traditional techniques, the two phases a and/3 constituting the eutectic show different morphologies consisting of randomly oriented inhomogeneous aggregates. On the other hand, if the solidification is controlled and occurs with heat extraction in a single direction, the two phases of the eutectic structure are present in the form of lamellae aligned parallel to each other, or the second phase appears in the form of rods immersed and aligned in the first or matrix phase. Several methods can be used to solidify ceramic eutectics directionally. All of them are, however, characterized by a high degree of complexity, especially when large or net-shape parts are required to be fabricated.


/

6O

5

3 o

anomalous I I

. , . . . C

g l ~ q I 40

%

!

I /

s

o / , I I o 20 60

ENTROPY OF SOLUTION &S{~. J/mol/K

FIG. 41. Classification of eutectic microstructures in terms of volume fractions and entropy of solution for an anomalous growth velocity of 5 • 10 -4 cm/sec: a, regular lamellar; b, regular rods; c, broken lamellar; d, irregular; e, complex regular; f, quasiregular; and g, irregular fibrous. After Elliot [23]. Reproduced with permission of the ASM Interna- tional, Metals Park.

Nevertheless, the potential of this type of material for self-reinforcement has been proven.

Sorrell [115] studied the directional solidification of the ZrC-ZrB 2 eutectic. The ZrC-ZrB 2 eutectic consisted of columnar grains with parallel lamellae within the grains. The fracture toughness of the transverse and longitudinal sections increased with decreasing interlaminar spacing, and the Kic values were higher than those of individual components (ZrC = 1.65 MPa m 1/2 MT, ZrB 2 = 1.87 MPa m 1/2 MT, composite = 5 MPa m 1/2 MT). The average toughness of directionally solidified TiC-TiB 2 also showed greater values than the toughness of either mem- ber [116]. Matson et al. [76] fabricated alumina-YAG and alumina-YAP eutectics with Chinese script and rod/lamellae morphologies, respectively. Both eutectics had elongated grains and colonies aligned with the solidification direction. The metastable alumina-YAP eutectic was converted by a solid-state reaction into its equilibrium phase alumina-YAG while maintaining the rod/lamellae morphology. The toughness value for the


alumina-YAG composite is 4.33 MPa m 1/2 MT, which is approximately the same as the highest values measured for sapphire and higher than single-crystal YAG (1.36 MPa m 1/2 MT) [75].

V. Conclusion

Although several excellent reviews analyze the effect of whisker addition on the mechanical behavior of ceramics [4, 5, 24], this review focused on the effect of phase chemistry on whisker-matrix interactions and on the relationships among the chemistry of the ceramics components, their grain morphologies, and their fracture toughnesses.

Several whisker-reinforced ceramics with fracture toughness up to three times higher than monoliths have been developed in research laboratories; but, except for A1203-SiC, their broad commercial applications are still very limited. This is due to the lack of stable commercial whisker suppli- ers, commercial unavailability of whiskers other than SiC and Si3N4, the chemistry variations from lot to lot, high manufacturing costs, and industrial hygiene concerns. Even small differences in the whisker surface chemistry can easily alter whisker-matrix interfaces and change material properties. Therefore, a better understanding of the phase equilibria in the whisker-matrix interface is essential to make processing reproducible and to optimize interface debonding. Further progress also requires the development of a new generation of ceramic whiskers with larger diameters (exceeding 1/xm) and tailorable surface chemistries.

The process of in situ whisker growth offers an alternative way for manufacturing whisker-reinforced ceramic composites without the neces- sity of whisker handling. The most suitable systems are those similar to Si3N4-SiC where both phases can form whiskers. The in situ growth requires a fundamental understanding of the system's phase equilibria since small differences in processing conditions and/or the presence of impurities are sufficient to change the phase formation.

The advantage of self-reinforcing is the ability to control microstructures and their potential applicability to many ceramic systems. Ceramics belonging to this group have unique characteristics. The development of directly solidified eutectics is based on a solid theoretical understanding of phase equilibria that allows one to design a variety of microstructures. Theoretical understanding, however, often exceeds manufacturing capabilities and limits practical utilization of this technology. On the other hand, self-reinforced ceramics such as Si3N4, where grain growth takes place in the presence of a liquid phase, can be fabricated with existing processing

222 A L E K S A N D E R J. P Y Z I K A N D A L A N M. H A R T

technologies. Broader utilization depends, to a large extent, on the development of a better fundamental understanding. The formation of grains with controlled morphology and size depends on several mechanisms that occur simultaneously at high temperatures. The control of the chemical composition, the essential part of this process, is often difficult due to the complexity associated with multicomponent systems. The availability of phase diagrams and thermodynamic data for five-, six-, and seven-component ceramic systems is also limited. Despite all of these difficulties, the development of technology to allow tailoring of the microstructure and the designing of interfaces is absolutely essential for the future growth of advanced ceramics.

The experimental data indicate that, with a proper system chemistry and proper processing conditions, many ceramic materials can produce elongated grains, thus increasing the flexure strength and fracture toughness. Self-reinforcing appears to be an emerging technology with great commercial potential. The ability to control grain morphology and grain interfaces together with a better understanding of the phase equilibria in self-reinforced materials could allow for microstructure design on a nanoscale. This would, in turn, create a basis for a new family of advanced ceramics.

Acknowledgment

The authors wish to thank Dan Beaman and Dave Susnitzky for SEM and TEM characterization; colleagues from Dow Chemical Central Research and Development Ad- vanced Ceramics Laboratory and Ceramics and Advanced Materials Laboratory for helpful discussions; Terry Tiegs from Oak Ridge National Laboratory for help in obtaining some whiskers; and Laura Riester from Oak Ridge National Laboratory for photographs of CVD SiC.

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127. H. Wada and L. Wang, Ceram. Eng. Sci. Proc. 11(9-10), 1463 (1990). 128. H. Wada, M.-J. Wang, and T. Y. Tien, J. Am. Ceram. Soc. 71, 837 (1988). 129. N. Wada, S. A. Solin, J. Wong, and S. Prochazka, J. Non-Cryst. Solids 43, 7 (1981). 130. R. S. Wagner and W. C. Ellis, Appl. Phys. Lett. 4, 39 (1964). 131. R. S. Wagner, W. C. Ellis, K. A. Jackson, and S. M. Arnold, J. Appl. Phys. 35, 2993

(1964). 132. H. Wang and S. Fischman, J. Am. Ceram. Soc. 74(7), 1519 (1991). 133. M. J. Wang and H. Wada, Mater. Sci. Forum 47, 267 (1989). 134. M. J. Wang and H. Wada, J. Mater. Sci. 25, 1690 (1990). 135. A. J. Whitehead and T. F. Page, Ceram. Eng. Sci. Proc. 10(7-8), 986 (1989). 136. S. Wild, P. Grieveson, and K. H. Jack, Spec. Ceram. 5, 385 (1972). 137. G. Wotting and G. Ziegler, in "Ceramic Powders" (P. Vincenzini, ed.), p. 951. Elsevier,

Amsterdam, 1983. 138. G. Wotting and G. Ziegler, Sci. Ceram. 12, 361 (1984). 139. G. Wotting and G. Ziegler, Ceram. Int. 10(1), 18 (1984). 140. G. Wotting, B. Kanka, and G. Ziegler, (1986),. in "Non-Oxide Technical and Engineer-

ing Ceramics" (S. Hampshire, ed.), p. 83. Elsevier, London and New York, 1986. 141. S. Yamada, S. Kimura, E. Yasuda, Y. Tanabe, and Y. Asami, J. Mater. Res. 3(3), 538

(1988). 142. A. Zangvil and R. Ruh, J. Am. Ceram. Soc. 71(10), 884 (1988).

Index

A

Acetic acid leaching, 36 Acheson method, 166 Acicular crystals, 166 Acicular grains, 201 Acicular grains-aspect ratio, 142 Activation energies, 142 Activator-essential, 8 Activity coefficient, 103 Activity gradient, 116 Adiabatic reaction temperatures, 112 AI-Si binary alloy, 106 AI-Si-C ternary diagram, 106 AI-Si-Mg alloy, 92 AI20 3, 95 AI matrix, 89 A14C4, 182

ceramic matrix, 88 flux, 46 AI203-MgAI204, 18 AI203-MgO-Y20 3, 20 oxide additive, 130 precipitates, 28 A1203-SiC, 182 AI203-SiC whiskers, 161 AI203-SiC-C, 182 A1203-SIO2 system, 45 AI203-Y203, solid solubility, 23 AIzO3/AI matrix, 89

Albite, 50 Albite-crystallization, 64 Alighned lamellae, 219

Alkali carbonate, 79 Alkali concentration gradient, 75 Alkali penetration, 72 Alkali bearing minerals, 71 Alkaline flux, 66 Alloy chemistry, 91 Alloy design principle, 154 Alloying elements, 106 Alloys, 87 AIN-AI20 3, phase diagram, 29 ALON, 2

ALON-aluminum oxinitride spinel, 28 entrapped porosity, 30 in situ forming, 30 microstructure, 30 ALON-pseudo binary, 28 ALON-Y20 3 sintering aid, 30

a-Alumina, 62, 91 a-/3 relationships, 175 Alumina (see also A1203), 15

high fired, 60 refractories, 79

Alumina grains properties, 184

Alumino-silicate refractory, 75 Aluminum

activity, 101 boride, 104 ions, 134

Aluminum saturated MgO, 100 Amorphous carbon layer, 198 Amorphous phase, 145 Annealing, 147

2 2 7

2 2 8 I N D E X

Anomalous eutectic, 218 Anorthite, 183 ARTs, see Adiabatic reaction temperatures Asbestos, 195 Ash chemistry, 166 Aspect ratio, 146, 171, 211 Asymmetrical phase diagram, 218 Atomic nitrogen, 201 Autoradiography, 9

B.

BaC , 93, 114 Be3N4, 136

interfaces behavior of, 184

BeO-oxide additive, 130 /3-alumina, 46 Bimodal grain distribution, 21 Binary phases, 215 Blast furnace, 69 Bond depletion, 61 Bond reaction, 61 Boride composites, 121 Boron-rich liquid, 115 Boundary temperature, 197 Bridging zone, 184 Brittle failure, 86 Bubble porosity, 113 Bulk analysis, 172

C

C-fiber reinforced Si3N 4 matrix, 111 Calcining, 6 Carbide composites, 121 Carbon coating, 192 Carbon deposit, 69 Carbon-rich interface, 185 Carnegieite, 46 Carnegieite-sodium aluminate, 50 Catalytic nitridation of whiskers, 193 Catalytic oxidation of whiskers, 193 Ceramic matrix composite, 86

growth schematic, 89 manufacture, 90

Ceramic skin, 92

Ceramic whiskers, 198 Chemical attack, soda, 79 Chemical changes, 186 Chemical potential gradient, 99 Chemical treatments, 184 Chemical vapor deposition, 165 Chemical vapor infiltration, 91 Chinese script morphology, 220 Clay refractories, 65 CMC, see Ceramic matrix composite Coal gasification, 43 Coated carbon fiber, 111 Colloidal silica, 178 Color emission, 15 Coloration, 3 Columnar growth, 92 Commercial

applications, 221 availability, 221 SiC whiskers, 167

Compatibility tetrahedrons, 131, 137 Compatibility triangle, 130, 206, 209 Composite growth surface, 98 Composite

material, 121 microstructure, 105 processing, 122 properties, 172

Compositional changes, 206 Condensed phases, 163 Congruent melting point, 55 Consolidation mechanism, 202 Constant volume system, 113 Conversion aid, 205 Cooling rate, 19 Cordierite, 139, 183 Corrosion resistance, 103 Corrosion-reaction rates, 44 Corrosive reactions, 80 Corundum, 49 Cost, 160 Covalent bonds, 127 Crack

bridging, 158 deflection, 158 propagation, 199, 211 tip blunting, 86

Cracking, 80 Creep, 85, 148

resistance, 140

I N D E X 2 2 9

Cristobalite, 48 Crucible cover, 79 Crystal growth, 207 Crystalline film, 191 Crystalline phase, 75 Crystallization of ceramic, 199 Crystallization velocities, 218 Crystallographic modification, 201 Crystobalite, 194 CVD, see Chemical vapor deposition CVI, see Chemical vapor infiltration Curie temperature, 38

Debond length, 159, 192 Debonding, 158 Decomposition temperatures, 85 Degree of flexure, 73 Degree of supersaturation, 166 Dense Y203, 9 Densification aid, 128, 205 Densification kinetics, 3 Dielectric constant, 35 Differential thermal analysis, 112 Diffusion, 210 Diffusion rates, 8 Diffusion

lattice, 7 vacancy, 7

Directly solidified eutectics, 221 Disilicate (Na20.2SiO2), 47 Disilicate

quartz eutectic, 49 Dislocation free whiskers, 169 Dissociation temperatures, 48 Dissolution of whiskers, 190 Double cantilever beam, 200 Dow Central Research, 178 Dry ball mill, 12 DTA, see Differential thermal analysis Duplex microstructures, 217 Dusts, effect on refractories, 68

Efficient toughening, 187 Electronic transport, 101

Elemental silicon, 169 Elevated temperature stability, 85 Elongated grains, 201, 222 Elongated Si3N 4 grains, 194 Emittance, 2 Empirical grain growth law, 141 Energy of formation, 209 Enstatite, 194 Entrapped porosity, 22 Equiaxed grains, 201 Equilibrium

conditions, 177 constant, 100 diagrams, 122 phases, 190 range, 189

Erosion, 67 Erosion resistance, 103 Euhedral crystals, KAISiO4, 57 Eutectic

composition, 22 formation, 182 solidification, 199

Eutectically solidifed ceramics, 218 Exothermic reaction, 177

Fabrication, 39 Factors

material properties, 187 Fe, as catalyst, 179 Feldspar laths, 57 Ferroelectric phase-anisotropy, 35 Fiber

degradation, 110 reinforcement, 91

Firebrick, 67 Fireclay, 45 Flexural strength, 140, 186, 205 Flue openings, 68 Fluxes, 44 Fluxing action, 68 Forsterite, 183, 194 Fossil fuel, 44 Fracture

resistance, 191 toughness, 86 toughness increase, 194

230 I N D E X

Free energy, 100 Free Zr, 93 Frequencies

infrared and visible, 4

fuel cells, 8 Fundamental understanding, 222 Furnace campaign, 71 Furnace fumes, 66

G

Gas flow rate, 195 Gas-phase approach, 172 Gaseous SiO, 163 Gehlenite, 183 Glass, 2, 184

borate-rich, 15 composition, 199 fusion point, 61 hygroscopic, 76 matrix bond, 70 melting, 59 preparation, 46 regenerators, 67

Glassy phase, 73 Glazed surface, 67 Glazing effect, 76 Grain boundary

glass, 145 phase, 12, 23, 141 pinning phase, 38 YAG crystallization, 147

Grain coarsening, 206 diameter, 212 diameters, 202 formation, 222 growth, 2

inhibitor, 22 kinetics, 7 retardants, 8

length, 142, 212 morphology, 159, 211 shape parameter, 145

Grain-boundary glass volume, 148 phase, 146, 192

precipitate, 22 segregation, 9

Grain-enhancing compound, 211 Graphitic carbon, 172 Graphitized carbon, 192 Grossularite, 183 Group-four additives, 8 Growth

barrier, 91 conditions, 170 exponent, 142 interface, 101 kinetics, 128 nucleation, 97 of ceramic, 199 rate, 88, 100, 204

Growth-transition liquid phase, 199

I-I

H phase, 210 Hardness, 85 Hearth zone, 69 Heat

of reaction, 114 treatment, 215

Heat treatments, 63 Hexagonal rods, 153 Hf-B-C, 120 HfB phase, 120 HfB precipitate, 121 High aspect ratio crystals, 195 High aspect ratios, 198, 205 High hardness, 152 High temperature corrosion, 81 High toughness, 152, 187 High-pressure, 189 High-purity silicon nitride whiskers, 178 High-resolution microscopy, 191 High-temperature mechanical

properties, 151 High-viscosity liquid, 204 Homogeneity range, 210 Hot isostatic pressing, 2 Hot-pressed SiA1ON, 149 Hot-pressing, 147 Hydrothermal methods, 54 Hygroscopic crust, 73 Hystereses behavior, 34

I N D E X 231

Impurities, reactants, 197 Incongruent melting, 55 Incrustation, 64 Infiltration temperatures, 119 Infrared applications, 2 Inhomogeneous aggregates, 219 Interaction coefficients, 102 Interface

debonding, 221 pressure, 181

Interfacial bonds, 190 layer

amorphous, 191 Intergranular phase, 128 Interlocking grains, 146 Intermediate catalyst, 169 Intermetallic phases, 103 Intrinsic grain growth, 26 Inversion temperatures, 48 Ionic conductors, 8 Iron catalyst, 198 Isostatic pressing, 6 Isothermal growth, 142 Isothermal ternary, 115

J

J phase, 209 Jackson's crystal growth model, 204

KAISiO4, orthorhombic, 59 K20 corrosion, 80 K20, freezing, 69 K20.AI203.2SiO 2, orthorhombic, 57 K20.A1203 compound, 53 Kaliophilite, 58 Kalsilite, 58 Kinetic effects, 122 Kinetic requirements, 198 Kinetics, 44

model, 102 of densification, 206 of growth, 170

Kinking of crack, 191

L

La203, 23 La203-Y203, 2, 4, 39 La203, addition, 35 Lamp envelop, 14 Lanthanide additives, 4 Lanthanides, 7 Large diameter crystals, 171 Large elongated grains, 208 Lattice

constants, 217 defects, 136 diffusion, 134 parameter, 8, 19

Lead-lanthanum zirconium-titanate, 32 Leucite, 51 Leucite zone, 71 LiF

grain boundary phase, 26 postpressing anneal, 27 solid solubility, 27

Light scattering, 1 Limiting solubility, 210 Liquid-phase mechanism, 202 Liquidus temperatures, 206 Local decomposition, 187 Long life reliability, 200 Los Alamos National Laboratory, 171 Low-toughness ceramics, 157 Low-viscosity glass, 204

M

Machining of ceramics, progress, 157 Macrointerface, 218 Mass transfer, 101 Material properties, 209 Matrix degradation, 189 Mechanical behavior of ceramics, 221 Mechanical properties, 222 Melting phase, 181 Melting points, 151 Metal impurities, 173 Metal matrix composites, 86 Metal oxidation, 87

2 3 2 I N D E X

Mg-Nd-Si-O phase diagram, 216 Mg-Si system, 100 MgAI204, 2, 95 MgA120 4

A1203, 25 LiF, 26

MgO, 95 A1203, 18 densification rate, 17 grain boundary pinning, 17 metastable, 97 properties, 101 segregation, 17 stabilizer, 17 surface diffusivity, 17

Mg2SiOa-MgSiN2 join, 132 MgTiO3, 18 Microhardness, 9 Microindentation techniques, 208 Microstructural

characterization, 193 design, 193

Microstructure, 204 Microstructure

A1203, 16 control, 221

Milling, 6 Mineralogical composition, 63 Miscibility gap, 50, 58, 217 MMC, see Metal matrix composites Molar ratio, 197, 215 Molten boron, 114 Monolith ceramics, 221 Morphology

of phases, 154 of silicon nitride grains, 128 of whiskers, 166

Mullite, 44, 178 decomposition, 63 needles, 65 refractories, 77

Multicomponent systems, 205 Multigrain junctions, 204

Na20-AI203-SiO 2 system, 46 Na20

fireclay, 78 mullite, 77

NaA1SiO 4, 50, 58 NaCl-space forming agent, 198 Nanoscale, 222 Nd203-doped lasers, 9 Near-net-shape, 6

processes, 87 Nepheline, 58 Nephelite, 66 Net-shape parts, 219 Ni doping, 104 Nitrogen ions, 134 Nitrogen partial pressures, 107, 177 Nucleation, ceramic, 199

O

Oak Ridge National Laboratory, 168 Optical

anisotropy, 38 applications, 1, 2, 15 ceramics polycrystalline, 11 properties, 14 transparency, 38

ORNL, see Oak Ridge National Laboratory Orthodontic braces, 15 Orthorhombic crystal structure, 176 Ostwald ripening, 117 Oxalates, 6 Oxidation process, 108 Oxygen

activity, 101 diffusion, 97 gradient, 98, 101 solubility, 97 transport, 90

Oxynitride glass, 201

P

Parameters, debond length, 159 Partial pressure, oxygen, 164 Partial pressure-gas phases, 196 Partial solid solution, 210 Particulate reinforcement, 86 Particulate SiC, 111 (Pb,La)(Zr,Ti)O3, 2 PbTiO3-PbZrO 3, tie-line, 32 Peeling, 68

I N D E X 2 3 3

Peeling of fireclay, 72 Peeling stresses, 73 Perovskite structure, 34 Perovskites, 2 Petrographic examination, 64 Phase

boundary curve, 189 diagram, 4, 6, 87

PbO-ZrO 2, 34 PbTiO3-PbZrO3-La, 34 PbTiO3-TiO2-ZrO2, 33 square, 128 ternary oxide, 81 trapezium, 128

equilibria, 109, 221 field diagram, 4 formation, 221 relations, 2, 39, 200 relationships, 168 stability diagram, 196 stoichiometry, 116

Phases N, A and G, 213 PbTiO3-PbZrO 3, 32

Phonon edge, 2 Photodiode detection, 8 Planar coupling, 35 Platelet-reinforced, 87 PLZT, see lead-lanthanum zirconium-

titanate Polycrystalline ceramic, 8 Polycrystals, pore-free, 23 Polymorphs

of Na20.2SiO2, 49 Polytypoid platelets, 153 Polytypoids, 136 Pore

clusters, 20 elimination, 14 entrapment, 7 sealing, 135

Porosity, 1, 7, 11 within grains, 25

Postdensification heat treatment, 212 Postsynthesis treatments, 171 Potash feldspar, 57 Potassia

glasses, 53 vapor, 54

Potassium carbonate vapors, 59

Potential reactions, 87 Powder properties, 14 Power generation, 43 Preferential degradation, 193 Preform reinforcement, 88 Presintering step, 194 Press forging, 2 Pressure assisted densification

techniques, 160 Pressureless sintering, 3, 161 Presynthesized powder, 134 Prismatic configuration of grains, 204 Processing, 213

conditions, 201 methods, 86 of ceramics

progress, 157 temperature, 205, 209 times, 205

Product whisker, 171 Properties

of commercial SiC whiskers, 174 of Si3N 4 whiskers, 181

Protective coating, 67 Pseudo-protective, 76 Pyrope, 183

Quarternary system, 129 Quartz, from melts, 56 Quasiternary reciprocal salt systems, 130 Quantitative modeling, 102 Quenching experiments, 52

Radiative heat losses, 113 Random orientation, 219 Rare-earth

ions, 3 oxides, 8

Rate constant, 142 Rates of reactions, 197 RBSN, see Reaction-bonded

silicon nitride

2 3 4 I N D E X

Reactant materials, 60 transport, 117

Reaction kinetics, 204 products

crystalline, 81 rates, 114

Reaction-bonded silicon nitride, 150 Reactive test, 60 Refractories, 43, 78

burned, 62 Refractoriness, 51 Refractory disintegration, 70 Refractory glass, discoloration, 74 Refractory grain boundary phases, 133 Regular rods, 219 Reversible inversions, 55 Ribbons, 178

S

Salts of sodium and potassium, 43 Sapphirine, 183 Scattering defect, 11 Second phase chemistry, 172 Self-propagating high-temperature

synthesis, 114 Self-reinforced ceramics, 200, 221 Shanghai Institute of Ceramics, 153 Short-range periodicity, 165 SHS, see Self-propogating high-temperature

synthesis Shrinkage, 63 Si

-AI-Be-N-O system, 130 -AI-N-O system, 128 -A1-Y-N-O system, 131 buildup, 102 -C phase diagram, 161

dilute concentrations, 162 -C-O system, 162 diffusion, 102 -Mg-N-O system, 206 -O-Si bonds, 176 -SiC

liquidus curve, 162

SiAION, 134 cordierite, 148 garnet, 148 grains

equiaxed, 139 Y3A15012 system, 147

SiC, 109 -AIN phase diagram, 216 whiskers, 185

SiN 4 tetrahedra, 173 Si3N4

-AI20 3 system, 134 -BeO system, 136 grain morphology, 203 grains-acicular, 140 -MgO, 132 production, 107 -SiC whisker interface, 192 -SiC whiskers, 161 -SiO2-AI203-A1N system, 129 -SiO 2-AI 20 3-AIN-Be O-Be 3 N2

system, 138 -SiO2-BeO-Be3Nz-BeO system, 138 -SiO2-YzO3-YN system, 133 -YzO3-AI203-SiO2 diagram, 210 -Y203-MgO-SiO2-CaO, 212 -Y2Si207 join, 149

Si2N20, 194 formation, 196

Si23N300 , oxynitride, 173 Silica

-alumina fabrication, 63

brick, 51 compositions, 76

Silicate chemistry, 76 depletion, 72 -meta (Na20.SiO2), 47 -ortho (2Na20.SiO2), 47

Silicon carbide formation, 163 polytypes, 165 crystalline forms, 164 factors determining form, 165

Silicon nitride ceramics flexural strength, 144 fracture toughness, 144 microstructure, 146

I N D E X 235

Silicon nitride forms, 173 glass interface, 206 mechanical properties, 127 systems, 137 thermal properties, 127

Silicon, thermochemical effect, 99 Silica, 45 Single-crystal whiskers, 158 Single-edge-precracked beam

method, 207 Single-phase field, 2 Sinter, 1 Sinter-hot isostatic pressing

technique, 207 Sintered

compositions, 207 silicon carbide, 217 Y203, 9

Sintering, 23, 107 additives, 183, 202 aid, 4, 8, 141

AI20 3, 12 MgO, 15

doped solid-state, 3 liquid-phase, 3 mechanism, 8, 24 processes, 187 shrinkage curve, 141 temperature, 13 transient liquid phase, 13, 135, 151 transient second-phase, 3

SiO generator, 170 SiO-C-O phases, 163 SiO 2, liquid phase, 176 Slag attack, 65 Slag, lime rich, 70 Slagging, 68 Slip casting, 90 Small grains, 205 Soda attack, 68 Soda-poor phase, 52 Sodium lamps, 1, 15 Solid-liquid interface, 204 Solid oxidant, 93 Solid solution, 7

formation, 138 gradients, 11 range, 217

Solidification direction, 220 microstructure, 117

Solubility, 202 gradient, 97, 103 limit, 8, 19

Solute diffusivity, 204 Solution

entropy, 218 precipitation, 127

Space constraints, 195 Specific volume changes, 215 Spinel, 23, 183

transparent, 26 Stability diagram, 188 Stable equilibrium, 187 Stacking faults, 173 Step configurations, 190 Stiffness, 85 Strength, 85

refractory, 72 to size relationship, 180

Stress corrosion, 9 Stress field, 86 Stress intensity factor, 200 Stress rupture tests, 137 Structural applications, 154 Structural flexibility, 183 Structure

fluorite, 4, 8 hexagonal, 4 monoclinic, 4 spinel, 29

Subsolidus phase relationships, 137 temperatures, 152 materials, 177

Surface chemistry, 171 Surface iron impurities, 193 Surface silica, 192 Synthesis, in situ, 198

T

Techniques deagglomeration, 12 synthesis, 12

236 I N D E X

Ternary diffusion couple, 117 eutectic, 51 eutectic temperature, 22 phase diagrams, 122 system, 215

Tetrasilicate (KeO �9 4SiOe), 55 TGA, see Thermat gravimetric analysis Theoretical understanding, 221 Thermal decomposition, 173

method, 179 Thermal expansion, 69

coefficient, 76 mismatch, 139

Thermal history, 165 Thermal shock resistant, 44 Thermo-gravimetric analysis, 114 Thermochemical

analysis, 107, 122 diagram, 177

of Si-O-N system, 175 techniques, 121

Thermodynamic requirements, 198 stability, 176

Thermodynamically stable phase, 201 Thermodynamics, 96 Thickness, 97 ThO2-Y20 3, 9 Thoenstatite, 183 Ti-B-C system, 118 Ti-fiux, 120 TiB-TiC-Ti, 118 TiB e particulates, 104 TiB2-C pseudobinary, 119 Ti3B4, 118 TiC solid phases, 120 Time exponent, 142 TiO x, 18 Titanium aluminide, 104 Toughening

criteria, 180 mechanism, 87 mechanisms, 158

Toughness data, 209 increase, 158

Toxicity of Be, 139 Transformation, 211

-a to/3, 200

Transition metal, 105 Translucent A1203, 1, 18 Transmission electron micrograph, 22 Transparency, 3, 8

LiF, 27 optimum, 24

Transparent lamps, 11 Transportation, 43 Triangular prisms, 178 Tridymite, 48

stability range, 56 Tuyere, 69 Two-phase

ceramics, 153 field, 2

Two-step reactions, 197

U

Unstable interfaces, 186

V

Vapor pressure of SiO, 132 Vapor transport, 18 Vapor-liquid-solid method, 165 Vapor-phase transport, 28 Vapor-solid method, 165 Vesicular layer, 65 Vitreous appearance, 70 Void volume, 113 Volatalization, 36, 52 Volatile metals, 69 Volume expansion, 80

W

Wear resistance, 85, 95 Weibull modulus, 157 Wetting characteristics, 103 Whisker

growth, 162 growth-enhancing agent, 205 handling, 195 like grains produced in situ, 160 -like structures, 218 -matrix interactions, 221 reinforcement, 107

INDEX 237

strength, 180 surface chemistry, 221 treatment, 160

Whiskers, 86 Whiskers-special shape, 179 Wurstite type structures, 164

X-ray diffraction, 92 peak intensity, 61

Y

YAG, see also Yttrium-aluminum- garnet, 209

Y203, 2 content, 133 -A1203, 12 -Gd20 3, 7 -MgO-SiO 2 phase diagram, 211

YAG, YAM, YAP, 209 YSiO2 N, 192 Yttria, 2 Yttrium-aluminum-garnet, 147

Zircon refractories, 60 Zirconium carbide, 87

matrix, 93 Zirconium diboride-microstructure, 94 ZnA1204, 92 Zr system, 113, 119 Zr-B-C phase diagrams, 115 Zr-B0.sC0.2, 116 ZrB 2 reinforced ZrC, 112 ZrB2-ZrC, 88 ZrB2-ZrC-Zr system, 112 ZrC-ZrB 2 eutectic, 220 ZrO 2, 8

Phase Diagrams in Ceramics

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