NASA Technical Memorandum 89830

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Ceramic Matrix and Resin Matrix Composites: A Comparison

(hASA-Ttl-8983c) C B B A B I C E A 1 1 5 I X A h D R Z S I t i NE7- 1 E6 15 B A ' i R I X C O B E G S I 3 E S : A CCflEAfilSCL I h A S A ) 1U p

CSCL 11D Unclas

G3/24 43726

Frances I. Hurwitz Lewis Research Center Cleveland, Ohio

. Prepared for the 32nd National SAMPE Symposium and Exhibition Anaheim, California, April 6-9, 1987

https://ntrs.nasa.gov/search.jsp?R=19870009182 2018-06-09T19:03:42+00:00Z

.

CEM4IC MATRIX AND RESIN MATRIX COMPOSITES: A COMPARISON Frances I. Hurwitz

NASA Lewis Research Center Cleveland, OH 44135

Abstract

The underlying theory of continuous

fiber reinforcement of ceramic matrix and resin matrix composites,

their fabrication, microstructure, physical and mechanical properties are contrasted. The growing use of

organometallic polymers as pre- cursors to ceramic matrices is discussed as a means of providing

low temperature processing capabil-

ity without the fiber degradation encountered with more conventional

ceramic processing techniques. Examples of ceramic matrix com- posites derived from particulate-

filled, high char yield polymers

and silsesquioxane precursors are provided.

1. INTRODUCTION The increasing acceptance of resin matrix composites as structural

materials for aerospace applica- tions has led to a heightened interest in extending the use of

composites to higher temperatures, particularly in engine environ- ments. Ceramics, because of their

stability at elevated temperatures,

offer significant potential in

engine applications; however, their use has been limited by their

brittle fracture behavior and high degree of flaw sensitivity. The incorporation of particulates,

whiskers or continuous fibers into

a ceramic matrix could provide mechanisms for "toughening" the ceramic, increasing reliability by

decreasing flaw sensitivity, and,

in the case of continuous fibers, ameliorating the tendency toward catastropic failure.

The base of knowledge gained

with resin matrix composites can in

part, serve as a base for the development of structural ceramic

composites. However, there are

important differences between ceramic matrix and resin matrix

composites in fabrication, micro-

structure, physical properties and mechanical behavior. This paper

offers a comparison of the two, as well as a discussion of opportuni- ties for the utilization of

1

orgaiioi?e t a i l i c polymers as precur-

sors t o ceramic composites. The

focus w i l l be on cont inuous f i b e r

r e i n f o r c e d m a t e r i a l s .

2 . THEORY

I n r e s i n m a t r i x composi tes

t y p i c a l l y the f i b e r modulus and

s t r e n g t h a r e much g r e a t e r t han

those of the matrix. The f i b e r ,

t h e r e f o r e , p rov ides s t i f f n e s s and

load bear ing c a p a b i l i t y , wh i l e t h e

ma t r ix serves t o d i s t r i b u t e t h e

load among t h e f i b e r s . Of ten ,

e s p e c i a l l y i n t h e case of epoxy

ma t r ix m a t e r i a l s , a coupl ing agent

i s introduced t o enhance t h e

s t r e n g t h of t h e i n t e r f a c e , i nc reas -

i n g load t r a n s f e r c a p a b i l i t y from

matrix t o f i b e r . The s t r a i n capa-

b i l i t y of t he ma t r ix t y p i c a l l y

exceeds t h a t of t h e f i b e r ; hence,

i n i t i a t i o n of f r a c t u r e is from t h e

f a i l u r e of a f i b e r o r group of

f i b e r s . The l oad i s then r e d i s -

t r i b u t e d t o t h e remaining f i b e r s

u n t i l a d d i t i o n a l f i b e r breakage

occurs .

I n c c n t r a s t , i n a ceramic

m a t r i x composite t h e moduli of t h e

f i b e r and mat r ix t y p i c a l l y are ve ry

s imi l a r . The f i b e r t h e r e f o r e i s

n o t t h e source of s t i f f n e s s o r

n e c e s s a r i l y of s t r e n g t h f o r t h e s e

materials (1). Pionol i thic ceramics

o f f e r s u f f i c i e n t s t r e n g t h f o r

s t r u c t u r a l a p p l i c a t i o n s , so incor -

p o r a t i o n of f i b e r s i s n o t necessary

t o achieve load bea r ing c a p a b i l i t y .

Ra the r , the f i b e r is r e l i e d upon t o

b r i d g e cracks or f l aws i n t h e

mat r ix . Because of d i f f e r e n c e s i n

p rocess ing betweeii ~ e s - i i d n d ceramic

m a t r i c e s , t h e ceramic ma t r ix i s

a n t i c i p a t e d t o have a higher popu-

l a t i o n of p rocess ing in t roduced

f l aws and microcracks than i t s

r e s i n c o u n t e r p a r t . T i l t s ~ i a i i i Ldpd-

b i l i t y of t h e f i b e r i s expected t o

exceed t h a t of t h e mat r ix . Hence,

ma t r ix f a i l u r e p recedes f i b e r f r a c -

t u r e . Microcracking of t h e m a t r i ; r

may provide some toughening as w e l l

as some p s e u d o p l a s t i c i t y ( 2 - 4 ) .

The d e s i r e d mode of f a i l u r e i s

debondilig a t t h e i n t e r f a c e . Because

u l t i m a t e ma t r ix s t r a i n , E: i s less m y

than u l t i m a t e f i b e r s tLd in ,

f i b e r f r a c t u r e would result i n

b r i t t l e f a i l u r e of t h e material .

By b r i d g i n g microcracks o r f l aws ,

t h e f i b e r can i n c r e a s e t h e s t r a i n

t o f a i l u r e of t h e composi te over

t h a t of t h e mono l i th i c ( s e e F igu re

l ) , impar t ing non-ca tas t rophic

f a i l u r e . I f t h e f i b e r l m a t r i x i n t e r -

f a c e is no t s t r o n g , c r acks can

propogate around t h e t i b e r s r a t h e r

t han through them. Therefore , t h e

pr imary r o l e of t h e f i b e r s i n a

ceramic ma t r ix composi te i s crack

b r i d g i n g and toughening by ma t r ix

c rack b l u n t i n g and debaridiiig a t t h e

i n t e r f a c e , as opposed t o provid ing

s t i f f n e s s and s t r e n g t h i n t h e r e s i n

ma t r ix composite.

,

3. FABRICATION

Resin m a t r i x composi tes t yp i -

c a l l y are processed a t r e l a t i v e l y

low tempera ture (<35O"c ) . The

polymer r e s i n may b e e i t h e r a

the rmop las t i c or thermose t , bu t

t y p i c a l l y w i l l undergo v i s c o e l a s t i c

2

flow, providing a means of fiber

impregnation and matrix consolida- t ion.

Glasses do become molten, and

offer the advantage of forming glass matrix composites by hot

pressing or transfer molding. A

number of glass matrix composites have been studied (5-7). However,

the melting points of many glasses are sufficiently high so that reac- tion with the fiber can become a

problem. A s with resin matrix

composites, the use temperature of

glass matrix composites will be

well below their softening point.

Carbon reinforced glass composites have been reported which offer zero

coefficients of thermal expansion

( 7 ) Ceramics, by contrast, do not

melt flow at temperatures low

enough for composite fabrication. Conventional sintering and hot

pressing temperatures usually ex-

ceed those at which available fiber reinforcements degrade in strength.

Oxide ceramic matrices can be

formed by sol-gel processing (8-9). Ceramic composites can also be

fabricated by chemical vapor infil-

tration (CVI) or by chemical vapor deposition (0) (10-11). Silicon

nitride matrix composites have been

fabricated successfully by reaction bonding of silicon powder in a

nitrogen atmosphere (12-13). Organometallic polymers also might serve as precursors which can be

pyrolyzed to a ceramic char (14-24);

these are discussed in greater

detail below. In addition, porous

carbon chars can be formed by poly- merization around a pore former,

and the resulting porous carbon

reacted with silicon to form silicon carbide (25). All of these tech-

niques can be expected to produce matrices with rather high levels of porosity (up to 35 - 40 percent), as opposed to perhaps < 3 - 5 percent porosity typically found in resin matrix composites. How well this

porosity can be tolerated will

depend on the uniformity of pore size and distribution and the dia-

meter of bridging fibers.

4. FIBERS Resin matrix composites can be

fabricated from a wide variety of

fibers, offering a broad choice of

modulus and strength for tailoring

to particular applications. The

selection of fibers available for ceramic composites is much more

limited, however, due to the higher

temperatures required for composite processing and in use.

Carbon fibers have been used

with glass matrices to temperatures of about 600°C (7). However, at

higher temperatures carbon will

react to form carbides and will be extremely susceptible to oxidation

when used in a porous matrix.

Sic fibers would offer good modulus and high temperature proper-

ties. Large diameter (140 pm) Sic fibers produced by CVD are available with temperature capabilities to

1400°C (26), but their diameter pre-

cludes their use in woven structures

3

o r i n t h e format ion of complex

shapes. For s t r e n g t h e n i n g a ceramic

ma t r ix , i n t e r f i b e r spac ing less

than t h e c r i t i c a l f l aw s i z e i s

needed; t h i s r e q u i r e s f i b e r d i a -

meters t o be s m a l l e r t han t h e i n t e r -

f i b e r spacing. The u l t i m a t e ma t r ix

t e n s i l e s t r a i n a l s o i s in f luenced

by f i b e r diameter and t h e s t r e n g t h

of t h e f iber -mat r ix i n t e r f a c e (2 ) .

Ava i l ab le sma l l d iameter

f i b e r s i nc lude polymer-derived

Nica lon Sic and ox ide f i b e r s such

as Nex te l 312, 440 and 480 (27 ) ,

comprised of b o r i a , alumina and

s i l i ca , and FP-alumina (28 ) . How-

e v e r , a l l of t h e s e become the rma l ly

u n s t a b l e a t t empera tu res from 1000

t o 1200OC.

5. PHYSICAL AND

MECHANICAL PROPERTIES

I n any composi te s y s t e m where

t h e r e i s a d i f f e r e n c e i n c o e f f i c i e n t

of thermal expansion (CTE) between

f i b e r and ma t r ix , r e s i d u a l thermal

stresses w i l l be expec ted t o develop

du r ing both composi te f a b r i c a t i o n

and on thermal cyc l ing . For most

f i b e r s , expansion i s a n i s o t r o p i c ,

d i f f e r i n g a long t h e f i b e r ax is and

t h e r a d i a l d i r e c t i o n .

f i b e r might prove t o b e t h e excep-

t i o n .

A pure B-Sic

When g r a p h i t e f i b e r i s used as

t h e re inforcement , t h e CTE a long

t h e f i b e r ax i s i s s l i g h t l y nega t ive .

Res idua l stresses developed du r ing

f a b r i c a t i o n p l a c e t h e f i b e r i n

compression and t h e ma t r ix i n ten-

s ion . I n any ceramic composi te

f a b r i c a t i o n approach i n which

m a t r i x sh r inkage t a k e s p l a c e , and

e s p e c i a l l y i n polymer de r ived

m a t e r i a l s , t h e m a t r i x w i l l l i k e l y

b e i n t e n s i o n as w e l l . However,

s i n c e ceramic m a t r i x composi tes can

be expec ted t o see h i g h t empera tu res

bo th i n f a b r i c a t i o n and i n s e r v i c e ,

t h e r e s i d u a l stresses are expected

t o be much g r e a t e r t h a n i n t h e i r

r e s i n m a t r i x c o u n t e r p a r t s , and may

g i v e r ise t o microcracking on fab-

r i c a t i o n . I n terms of des ign ing

w i t h composi tes t h i s l e a d s t o a

s i g n i f i c a n t d i f f e r e n c e between

r e s i n and ceramic m a t r i x m a t e r i a l s .

Whereas r e s i n , g l a s s and probably

RBSN m a t r i x composi tes can be

des igned as a c r o s s p l y layup of uni-

d i r e c t i c n a l lamina, ceramic ma t r ix

composi tes i n which ma t r ix sh r ink -

age occur s i n p rocess ing l i k e l y

w i l l have t o be f a b r i c a t e d from 2 D

c l o t h o r as 3D woven s t r u c t u r e s t o

minimize sh r inkage c racking .

Some ceramic m a t r i x composi tes

might a l l e v i a t e t h e problem of

r e s i d u a l stresses by cho ice of

f i b e r s w i th h igh a x i a l CTE ' s rela-

t i ve t o t h e ma t r ix , which would

p l a c e t h e m a t r i x i n compression.

However, i f t h e CTE of t h e f i b e r i s

much g r e a t e r t han t h a t of t h e ma t r ix ,

debonding of t h e i n t e r f a c e w i l l

r e s u l t (3 ) .

Decreased matrix s t r a i n capa-

b i l i t y r e l a t i v e t o t h e f i b e r may

dec rease t e n s i l e s t r e n g t h , as t h e

m a t r i x cannot p l a s t i c l y deform t o

accommodate stress concen t r a t ions .

Behavior under f l e x u r a l l oad ing

might become more complex. Prewo

.

4

(6) has studied Nicalon/epoxy and Nicalon/lithium aluminosilicate

(LAS) composites in both tension and flexure, and notes that while the Nicalon f epoxy composites had

the higher tensile strength, in

flexure the LAS matrix material appeared stronger as it "yielded"

due to microcracking, shifting the

neutral axis of the beam toward the compressive side of the test speci-

men. Thus, ceramic composites may

appear to have greater strength in flexure than their resin matrix

counterparts as a result of both

their higher matrix compressive strength and the pseudoductility

imparted by microcracking.

Microcracking also can become

a problem to the environmental

stability of the fibers in air,

leading to fiber degradation and

fracture during use at elevated

temperature. 6. POLYMERIC PRECURSORS TO

CERAMIC MATRICES

The use of organometallic pre-

cursors which can be pyrolyzed to a ceramic char provides a means for

forming ceramic matrices utilizing the advantages of ease of fiber infiltration, control of rheology

and low temperature processing

typical of resin matrix composites. Choice of polymer is influenced by stoichiometry of the resulting char

and shrinkage on pyrolysis. Ideally, thr! closer to stoichio-

metric Sic or S i N the final 3 4 product, the more thermally stable it may be expected to be. Also,

the smaller the weight loss and volumetric shrinkage on pyrolysis,

the less the need for multiple re-

impregnation cycles and the less the likelihood of large cracks.

A number of organometallic

polymers have been studied; some of these are shown in Table I. The

polycarbosilane work of Yajima (14- 16) serves as the basis for Nicalon fibers. The char is rich in excess

carbon. It also contains oxygen intentionally introduced as a cross- linking agent to stabilize the fiber

structure prior to pyrolysis. The

deviation from stoichiometry results in thermal instability above 1200°C.

Most of the polysilanes and

polycarbosilanes listed undergo weight losses of 40-60 percent on

pyrolysis. Higher char yields have

been demonstrated for the polysila- zanes; 80-85 percent char yields

have been reported by Seyferth (23) .

The polysilazanes are, however, moisture sensitive, and require com-

posite processing to be carried out

in inert atmospheres. Recent work in our laboratory

( 2 9 ) has examined a group of silses-

quioxanes (Figure 2) having the

general structure RSiO 1-59 where R = methyl, phenyl, propyl, or

vinyl, as Sic precursors. The sil- sesquioxanes melt flow, thermally

crosslink, and then can be pyrolyzed

at nominally 50OoC. peratures they can be expected to

undergo a carbothermal reduction to

Sic with loss of CO (Figure 3 ) . By controlling the starting ratio of

At higher tem-

5

SifC we hope to control the stoich-

iometry of the end product. Microstructure of Nicalon/

silsesquioxane composites is shown

in Figures 4 and 5 . Initial im?reg- nation shows few voids. After

pyrolysis followed by heating to

1000°C in argon, matrix shrinkage and pore formation are evident.

(Figure 4 ) .

in the optical micrograph (Figure 4h) is epoxy which was vacuum in- filtrated into the composite after

pyrolysis for polishing, and shows the extent of matrix cracking. The

composite might be reimpregnated

with silsesquioxane and again pyro-

lyzed to increase matrix density.

The darkest phase seen

Matrix surface cracks both

parallel and perpendicular to the fibers also are observed (Figure 5 ) ,

with the more matrix rich surfaces

showing the higher extent of crack- ing. These cracks likely arise

from a combination of shrinkage on

pyrolysis, mismatch in CTE between fiber and matrix and anisotropic fiber expansion.

We also have studied a Sic particulate filled, high char yield

carbon resin (30) which introduced

the particulate third phase as a means of minimizing matrix shrink- age and cracking. We were able to

fabricate unidirectional composites

with very low pore volumes and few

cracks on a single impregnation and pyrolysis cycle, provided that we used small fiber tow sizes and fibers which were not uniformly

cylindrical (compare Figures 6 and

7). However, strain capability of

the matrix was low (%.3%), and ten- sile specimens failed in shear in the tab region. Stress-strain

behavior was linear to fracture. Cross-ply layups showed extensive

matrix cracking and delamination,

the result of residual thermal stresses.

Jamet et - al. (31) have shown

that additives of a BN particulate to a polyvinylsilane minimizes

linear shrinkage and weight loss on

pyrolysis. Thus, the concept of particulate filled precursors would

seem to warrant further study.

New polymers which yield more stoichiometric ceramic products, as well as small diameter, thermally

stable fibers are needed, as is a greater understanding of the pyro-

lysis process and resulting stresses

at the fiber-matrix interface, for major advances in new ceramic com- posite materials to be achieved.

7. ACKNOWLEDGMENTS The author wishes to thank

Drs. Donald R. Behrendt and James A. DiCarlo for heipfui discussions.

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c

T A B L E I

Some organometallic precursors to ceramics

Polymer Structure Reference

I t

I (-Si- $;-I,, 14-16 polycarbosilane

CH, poly vinylsilene 1

0, P Si- polyborosiloxane

B-O- 0,

#

polysilastyrene

R, F r R, H

- Si N - S i N-

polysilazane .Si-N, R Si-A, R --N Si-” Si -

I I 1 I

“N -Si\’ > -! I

$I R H

2 1

16

17

22-23

c

R = CH2

9

FIBER BRIDGING OF MATRIX CRACKS

2si01.6 + 5~ + 2 s i c + 3 C O t

SiO,, + C(0XCO.D) + sic + c + co t OU

01 STRESS

Sio,, + c (doflclont)- Sic + Si0

O m + SiOt + COT

4n El EU Figure 3 - P o s s i b l e carbothermal cs-u-,7w r e d u c t i o n p roduc t s of s i l s e s q u i o x -

STRAIN anes based on t h e S i / C r a t i o i n t h e s t a r t i n g polymer.

F igu re 1 - T h e o r e t i c a l stress s t r a i n behavior f o r cerar-Ep ma t r ix composi te f i b e r br idging of ma t r ix c racks . (F igure cour t e sy of D r . James A . DiCarlo) .

SILSESQUIOXANES

RSi0,,5

R = methyl, propyl, vinyl, phenyl

T Resin Ladder Polymer

Figure 2 - Si l se squ ioxanes may e x h i b i t e i t h e r an extended c ross - l i nked (T r e s i n ) o r double l a d d e r s t r u c t u r e , o r a mix tu re of t h e two. The extended s t r u c t u r e i s more preva len t when R = phenyl. F igu re 4 - Nica lon / s i l s e squ ioxane

composi te ( a ) as f a b r i c a t e d and (b) a f t e r p y r o l y s i s fol lowed by h e a t i n g t o 1000°C i n argon.

- 1 mm

Figure 5 - Sur face c rack ing of N ica lon / s i l s e squ ioxane composite (a) on p y r o l y s i s a t 525°C f o r 2 hours and (b) a f t e r h e a t i n g t o 1000°C in argon.

F igu re 6 - Cel ion r e i n f o r c e d h igh cha r y i e l d po ly (a ry lace ty1ene ) ma t r ix f i l l e d wi th S i c p a r t i c u l a t e showing ( a ) l a y e r i n g of f i b e r r i c h and ma t r ix r i c h r e g i o n s and (b) ma t r ix areas devoid of p a r t i c u l a t e (arrows) .

Figure 7 - ( a ) Nexte l and (b ) Nicalon r e i n f o r c e d S i c p a r t i c u l a t e f i l l e d poly(ary lace ty1ene) mat r ix . cannot close-pack (Compare Figure 6 ) .

Composite homogeneity i s enhanced when f i b e r s

11

____- . -

1 Report No

NASA '1M 89830

Ceramic M a t r i x and Resin M a t r i x Composites: A Comparison

___-- 2 Government Accession No 3 Recipient's Catalog No

506-43-1 1

4 Title and Subtitle 5 Report Date

Frances I . Hurw i t z I E 3481

-- 7 Author(s)

_____

8 Performing Organization Report No

- 9 Perlorming Organization Name and Address

N a t i o n a l Aeronaut ics and Space A d m i n i s t r a t i o n Lewis Research Center Cleveland, Ohio 44135

9 Security Classif (of this report) 20 Security Classif (of this page) 21 No of pages

U n c l a s s i f i e d U n c l a s s i f i e d 12

-I

22 Price'

A0 2

11. Contract or Grant No. c 13. Type of Report and Period Covered

2. Sponsoring Agency Name and Address Techn ica l Memorandum

R a t i o n a l Aeronaut ics and Space A d m i n i s t r a t i o n Washington, D.C. 20546

5 Supplementary Notes

Prepared f o r t h e 32nd N a t i o n a l SAMPE Symposium and E x h i b i t i o n , Anahejm, C a l i f o r n i a , A p r i l 6-9, 1987.

6 Abstract

l h e u n d e r l y i n g theo ry o f cont inuous f i b e r r e i n f o r c e m e n t o f ceramic m a t r i x and r e s i n m a t r i x composites, t h e i r f a b r i c a t i o n , m i c r o s t r u c t u r e , p h y s i c a l and mechan- i c a l p r o p e r t i e s are con t ras ted . The growing use o f o r g a n o m e t a l l i c polymers as p recu rso rs t o ceramic m a t r i c e s i s d i scussed as means o f p r o v i d i n g l o w temperature p rocess ing c a p a b i l i t y w i t h o u t t h e f i b e r d e g r a d a t i o n encountered w i t h more conven- t i o n a l ceramic p rocess ing techn iques . Examples o f ceramic m a t r l x composi tes d e r i v e d f r o m p a r t l c u l a t e ~ f i l l e d , ' h i g h char p recu rso rs a r e prov ided.

y i e l d polymers and s i l s e s q u i o x a n e

7. Key Words (Suggested by Author(s))

Ceramic m a t r i x comgosites; Polymer p r e c u r s o r s ; Phys ica l p r o p e r t i e s ; Mechanical p r o p e r t i e s

18. Distribution Statement

U n c l a s s i f i e d - u n l i m i t e d S l A R Category 24