Fiber Reinforced Ceramic Composites

Fiber_Reinforced_Ceramic_Composites/Fiber Reinforced Ceramic Composites/12333_01.pdf1

Mechanical Behavior of Ceramic Matrix Composites

Anthony G. Evans and David B. Marshall

INTRODUCTION

Practical ceramic matrix composites reinforced with continuous fibers exhibit important failure/damage behaviors in mode I, mode II, and mixed mode 1/11, as well as in compression.ThefaiIuresequence depends on the applied stress state as well as on whether the reinforcement is uniaxial, laminated, or woven. However, the underlying failure processes are con- veniently illustrated by the behavior of uniaxially reinforced systems. The basic features are sketched in Figure 1. The intent of this chapter is to provide an assessment of relationships between the properties of the constituents (fiber, matrix, interface) and the overall mechanical perforrn- ance of the composite. At the outset, it is recognized that the composite properties are dominated by the interface and that upper bounds must be placed on the interface debond and sliding resistance in order to have a composite with attractive mechanical properties. A major emphasis here thus concerns the definition of optimum propertiesfor coating sand interphases between the fibers and the matrix, subject to high temperature stability and integrity. Residual stresses in the composite caused by thermal expansion differences are also important and are confronted throughout.

The strong dependence of ceramic matrix composite properties on the mechanical properties of the interface generally demands consideration of fiber coatings and/or reaction product layers, at least for high temperature use. Thus, while low temperature matrix infiltration procedures, such as chemical vapor infiltration (CVI), can create composites that exhibit limited interface bonding and, therefore, have acceptable ambient temperature properties, experience indicates that moderate temperature exposure

1

2 Fiber Reinforced Ceramic Composites

0 t t

Fibers

Matrix (Mode I ) Cracks

Initial Notch

Delamination

Crack (Mixed Mode) 0

a) Tensile Damage

P/2

ixed Mode) Crack Matrix (Mode I ) Crack

b) Flexural Damage

Figure 1. A schematic illustrating the failure modes observed in high toughness uniaxially reinforced ceramic matrix composites. (a) Tension. (b) Flexure.

Mechanical Behavior of Ceramic Matrix Composites 3

causes diffusion, coupled with the ingress of O,, N,, etc., from the environ- ment, resulting in chemical bonding across the interface. The resultant interphases consisting of oxides, nitrides, and carbides (either separately or in combination) invariably havesufficiently high fracture resistance that desirable composite properties are not retained. Consequently, a major objective of continuing research on ceramic matrix composites is the identification of interphases that are both stable at high temperature and bond poorly to either the fiber orthe matrix. Certain refractory metals and intermetallicsseem to have theseattributes, as elaborated in the following chapters.

The basic philosophy of this chapter is that the overall mechanical behavior is sufficientlycomplexand involvesasufficiently large number of independent variables that empiricism is an inefficient approach to microstructural optimization. Instead, optimization only becomes practical when each of the important damage and failure modes has been described bya rigorous model, validated by experiment. The coupling between experiment and theory is thus a prevalent theme. It is also noted that this objective can only be realized if the models are based on homogenized properties that describe representative composite elements, while also taking into account the constituent properties of the fibers, matrix, and interface. Models that attempt to discretize microstructural details have little merit in the context of the above objective. In this regard, the present philosophy is analogous to that usedsuccessfullytodescribe processzone phenomena such as transformation and microcrack toughening,l-S as well as ductile fracfure,W wherein the behavior of individual particles, dislocations, etc., provides input to the derivation of constitutive properties that describe the continuum behavior.

The behaviorof thecomposite is intimatelycoupled tothe basicfeatures of crack propagation and sliding along interfaces. This is demonstrated first by examining the damage and fracture processes that occur in each of the important modes depicted in Figure 1. The results of these studies indicate the need for studying interface responses in judiciously selected test specimens. The basic mechanics and the implicationsof tests used to study interface debonding and sliding are then presented. Finally, implica- tions for the choice of matrices, fibers, and coatings that provide good mechanical properties are discussed. The tensile properties are discussed initially in qualitative terms, involving consideration of both debonding and sliding at fiber/matrix interfaces, as well as pullout, in accordance with the sequence depicted in Figure 2. Then, the special but important case of unbonded fibers will be given quantitative attention.

TENSILE BEHAVIOR OF UNIAXIAL COMPOSITES

Debonding and Sliding Present understanding of toughening of ceramics by brittle fibers is

consistent with the debonding and sliding events illustrated in Figure 2. To allow crack bridging by the fibers, debonding at the fiber/matrix interface


ITT K 1

4 Figure 2. A schematic illustrating the initial debonding of fibers at the crack front and fiber debonding in the crack wake.

must occur in preference to fiberfailure at the matrixcrackfront. When this condition is satisfied, the sliding resistance, s, of the debonded interface hastheimportant roleof governing the rateof loadtransferfromthefiberto the matrix. Specifically, large s enhances load transfer, causing the axial stress in the fiber to decay rapidly with distance from the matrix crack plane. Consequently, weakest link statistical arguments dictate that the fibers fail at locations close to the crack plane, thus diminishing the vitally important pull-out contribution to the mechanical properties. A small sliding resistance along the debond thus promotes high toughness.

The extent of debonding at the crack tip is typically small when residual compression exists at the interface, but can be extensive when the interface is in residual tension. However, more importantly, further debonding is typically induced in the crack wake? The extent of debonding is again governed largely by the residual field. Residual radial tension encourages extensive bonding, whereas in the presence of residual compression, debonding is stable. The extent of debonding is determined by the friction coefficient and morphology of the debonded interface.

Analysis of fiber debonding at the tip of a matrixcrack(discussed later in this chapter) indicates that debonding rather than fiber failure occurs,


Interlace

I I I -0 5 0 05 1 0

Elastic Mismatch. (Y

-. N

I I -0 5 0 05 1 0

Elastic Mismatch. (Y

Figure 3. The critical energy release rate required for crack front debonding.

provided the fracture energy of the interface, G,,, is sufficiently small compared with that of the fiber, Gfc (Figure 3),9

In this chapter G IS represented by G. G,< I G k 5 I/4

There is no direct experimental validation of this requirement. However, various observations of crack interactions with fibers and whiskers support the general features.lQl1 In particular, experiments on LASISIC composites reveal that as-processed materials with acarbon interlayerdebond readily and exhibit extensive pullout (Figure 4a), whereas composites heat treated in air to create a continuous SiO, layer between the matrix and fiber exhibit matrix crack extension through the fiber without debonding (Figure 5). Furthermore, compositeswith a thin interface layer of SiO, having apartial circumferential gap exhibit intermediate pullout characteristics(Figure 6). The associated constituent properties are summarized in Table 1. Based on these properties, the preceding arguments would indicate that crack


Figure 4. Interfaces and pullout in a composite consisting of LAS matrix and Sic (Nicalon) fibers: As-processed indicating C interlayer, thermal debonding, and extensive pullout.

Figure 5. LAS matrix/SiC fiber composite heat treated in air for 16 hours at 800C indicating a complete SiOp layer and no pullout.


Figure 6. LAS matrix/SiC fiber composite heat treated in air for 4 hours at 800C indicating a partial SiO, layer-with gap-and variable pullout.

front debonding should not occurwhen acompleteSi0, layerexistsatthe interface, whereas, appreciable crack front debonding should occur when the C layer is present, in accord with the observations.lo911 Thecomposites with only a partial SiO, interface layer are also interesting. For these materials, Gi, is related to the fraction of the circumference that bonds the fiber to the matrix: typically 1/3. Reference to Table 1 and to the initial debonding requirement (Figure 3) would thus indicate that debonding, while marginal, is possible.

Fiber (Nicalon)

Matrix (LAS)

Interface Amorphous C Amorphous Si02

E(GPa) S;e Urn-2) a (K-1)

200 4-8' 4 x 10-6

85 40 1 x 10-6

-- < 1 * -- 80 8 1 x 10-6


INTERFACES)

DISPLACEMENT

Figure 7. A tensile stress-strain curve for a tough ceramic composite.

ultimate strength is determined by fiber bundle failure, with the tail of the curve corresponding to pullout of broken fibers. This type of failure mechanism has beenobserved inawiderangeofcomposite materials, including reinforced cements, glasses, and glass-ceramics.15-18 If, on theother hand, a substantial proportion of the fibers break in the wake of the first matrix crack as it extends, then failure of the composite is catastrophic. In this case, the ultimate strength is limited by the growth of a single dominant crackand isdetermined byafracture resistancecurve.19The natureof the resistance curve and the magnitude of the steady-state toughness is governed by the zone of bridging fibers behind the crack tip. The composite properties associated with the above failure mechanisms are discussed below, along with the criteria that dictate the transition between mechanisms.

Some Basic Mechanics The opening of a crack bridged by fibers involves stretching of fibers

between the crack surfaces. This stretching may be characterized by a relation between the stress, t, in the fibers and an average local crack opening displacement, u, as depicted in Figure 8. The form of this relation depends on the details of the bridging mechanism and reflects properties such as fiber/matrix debonding and frictional sliding, as well as elastic stretching of the fiber. The peak value, t = S, represents the strength of the fiber, whereas the decreasing portion depends on the nature and location of fiber failure.

Thet(u) relations in Figure8 represent the rangeof behaviorexhibited by brittle reinforcing fibers. At one extreme, for a fiber that is sufficiently well


bonded that no debonding occurswhen the crackcircumscribes it, the t(u) relation is linear to failure. At the other extreme, for a fiber that is not bonded at all to the matrix, frictional forces resist pullout. Initially, the t(u) relation is an increasing function of u until the fibers break, then t(u) decreases as the broken fibers pull out of the matrix. Intermediate t(u) relations result from partial debonding and frictional sliding over the debonded crack surfaces.

t

BRITTLE REINFORCEMENT

DEBONDING

UNBONDED INTERFACE,

INTERMEDIATE T

U

Figure 8. A schematic illustrating various trends in crack opening with bridging stress

The influenceof bridingfibersonfractureof thecompositecan beevalu- ated by two equivalent approaches, both of which model the composite around the crack as a continuum and employ the t(u) relation as the link between the constitutive properties of the composite and its macroscopic continuum behavior. In one approach, the stresses in the bridging fibers are viewed ascracksurfaceclosure tractionswhich reduce thestressesat the crack tip.14 The corresponding reduction in crack tip stress intensity factor is calculated from these surface tractions using a standard Greens function. The criterion for crackgrowth is obtained by setting the resultant crack tip stress intensityfactor in the matrix equal to the toughnessof the unreinforced matrix. The alternative approach is to use the J-integral to evaluate the effect of the bridging tractions on the energy flux20.21 In general, both of these approaches require numerical solution of an integral equation to calculate the crack opening displacements in order to specify the distribution of closure tractions over the crack surface. However, for steady-state configurations, the J-integral approach provides simple analytical results and is thus more useful.

g A n n d

*c

STEADY-STATE MATRIX CRACKING STRESS

-=- I I I I I I


Resistance Curves and Toughening If the steady-state matrix cracking stress given by Equation (1) exceeds

the stress, fS, that can be supported by the fibers, then the fibers within a fully-bridged crack break before the crack extends in the matrix. The consequent reduction in bridging forces causes the crack tip stresses to increase (at constant applied stress), resulting in unstable crack extension in the matrix, accompanied by further fiber failure.23 The corresponding stress- displacement curve is linear to the peak load and failure is catastrophic.

A preexisting crack without bridging fibers (e.g., a notch cut by a saw) grows in the matrix initially without breaking fibers. A bridging zone develops behind the advancing crackfront, resulting in increasing closure tractions as the crack grows. Consequently, the applied stress intensity factor needed for continued crack growth increases, so that crackgrowth is dictated by an increasing crack resistance curve (R-curve), as depicted in Figure 12. In general, calculation of the rising part of the R-curveand the amount of crack extension needed to achieve steady state requires numerical solution of an integral equation to obtain the crack opening displacements.23 In composites with small bridging zones, the steady- state toughness increment is of primary interest. However, in composites with large bridging zones, the entire R-curve must be specified, because the steady-state toughness may never be achieved by a stable crack(e.g., if the crack extension needed is larger than the specimen width).

TOUGHNESS - - - - - - - -

G

CRACK EXTENSION, A,

Figure 12. Aschernatic resistancecurve for crackextension in auniaxialcornposite.

A simple analytical solution for the steady-state toughness increment has been derived using the J-integral?021

" AS;< = 21;(u)du (2)


where uo is the crack opening at the end of the bridging zone. For the steady-state crack, uo is the displacement above which the bridgingforces are zero and AG, is given by the area beneath the p(u) curve, as depicted in Figure 11 b. Therefore, if the p(u) relation is specified, then AGc can be evaluated without having to determine the crack opening displacements.

Transition in Failure Mechanism and Optimization of Properties The resultsof the previous twosectionsallow somegeneralconclusions

to be drawn concerning the dependence of steady-state toughness and steady-state matrix cracking stress on microstructural properties. If we begin with a composite that fails by the noncatastrophic, multiple matrix cracking mechanism, then, as discussed previously, any change in the nature of the bridging ligaments that stiffens the increasing portion of the p(u) relation causes the steady-state matrix cracking stress to increase. However, if uc exceeds the peak in the p(u) curve, a transition in failure mechanism must occur and the steady-state toughness is given by the area beneath the p(u) curve. Then, further increase in stiffness of the increasing portion of the p(u) relation would usually lead to a decrease in toughness, provided that the peakvalue of p(u) remainsconstant. Therefore, the optimum properties(i.e., maximumu,orAGJ occur inthevicinityof the transition between the two failure mechanisms.

Residual Stress Residual microstructural stresses arise generally from thermal contrac-

tion during cooling from an elevated processing temperature.The residual stresses before cracking are of opposite sign in the reinforcing ligaments and matrix, and the average residual stress normal to a potential crack plane that spans many microstructural units is zero. Therefore, in the absence of a bridging zone, the microstructural residual stresses have no effect on the steady-state fracture toughness. Moreover, the fracture mechanics analysis of bridging, expressed in terms of the p(u) relation, is unaffected by the presence of residual microstructural stresses. However, the residual stresses influence the p(u) relation and thereby the magni- tudes of the matrix cracking stress and the fracture toughness.24

The influence of the residual stress on the p(u) relation is dependent upon the mechanisms of interfacialsliding andfiberfailure.An offset in the origin always occurs by an amount

0 0 = - q E / E m (3)

where q is the axial residual stress in the matrix and E and E,,, are the Youngs moduli of the composite and the matrix. The remainder of the p(u) curve is simply translated by a, for some composites, but in general, the shape of the p(u) curve may be modified as well.

For composites in which residual stress translates the p(u) relation uni- formly by a,, it is readily deduced from Figure 1 1 a that the matrix cracking stress a, is either increased (q compressive) ordecreased(q tensile) byu,. However, the sign and magnitude of the steady-state toughness change

STRESSlDISPLACEMENT RUPTURE LAW CONDITION

Linear Stress

Displacement

Frictional Surface With Pullout Roughness Stress

Coulomb Friction Stress

RESIDUAL STRESS IN FIBER

COMPRESSION TENSION

Decrease Decrease

Increase Decrease

---------- Negligible

I-------- Decrease

E(GPa)

A 1 2 0 3 400

Si3N4 300

Sic 400

Amorphous interface 70

f (Jrn-?) a (IC*)

20-30 7 x 10-6

60-80 3 x 10-6

15-20 4 x 10-6

-- 4-8


the normal interfacial stress. In principle, with a statistical distribution of fiber strengt hs, some f i ber fail u re occurs ahead oft he crack tip as well as in the crack wake. Analysis of fiber failure ahead of the crack has not been attempted, partly because the problem is complex and because of a perception that fiber failures close to the crack plane that cause pullout are most likely to occur in the crackwake. Such behavior has indeed been observed in glass reinforced plastics.27 However, there is no direct evidence that fiber failure ahead of the matrix crack can be neglected in ceramic matrix composites.

rn =O

.-. ---,m=~ ------____

'- - - - - - - - - - - - - - - - I I 1 0 , I

1 2 3 4 5

rn = 5

0 - 0

(UIV)

Figure 13. Nondimensional stress, crackopening curves for bridging plus pullout for various values of the statistical shape parameter, m.

Nevertheless, it is insightful to examine solutions for the p(u) relation based on wake failure in composites having debonded interfaces subject tosmal1,constant sliding resistance,z,and negligibleresidualstress.28The analysis involves calculation of the distribution of fiber failure sites as a function of applied stress, a, and hence, the reduction in stress due to fiber failure. The fiber strengths are defined by a Wiebull distribution with shape and scaling parameters m and S,. The results summarized in Figure 13 indicate that the initial, rising portion of the p(u) curve is dominated by intact fibers, the peak is dominated by multiple fiberfailures, analogous to bundle failure, and the tail is governed by pullout. Theinitial rising portion of the curve is closelyapproximated by the limiting solution (m = m) for all m:14

[ 4Tf2EfE2 2['*2 J /2 (5) P =

R E ~ , ( I - f)


where R is the fiber radius, f the volume fraction of fibers and 6, E,,,, and E are the Youngs moduli of the fibers, matrix, and composite. However, the tail of the curve is more sensitive to m: as m decreases, corresponding to a broader distribution of fiber strengths, more fibers fail further from the matrix crack, causing the extent of pullout to increase.

Correlation of calculated and experimentally measured pull-out lengths on the fracture surfaces of broken test pieces provides a route for measuring the statistical parameters and interfacial sliding resistanceJ1.29 The calculated cumulative probability that the pull-out length will be I h is plotted in Figure 14. The results indicate that the pull-out lengths tend to increase as m decreases, as expected. Preliminary estimates of the eff ects of residual strain suggest that the pull-out length usuallydecreases as the residual strain increases, when the residual stress at the interface is compressive. However, specific trends are sensitive to m, as well as to the friction coefficient k Measurements in the LAS/SiC system are discussed below.

-

c

0.0 0.1 0.2 0.3 0.4 0.5 Non-Dimensional Pull-Out Length (2 T h /ax)

Figure 14. The cumulative pull-out distribution for several values of the shape parameter, m.

Matrix Cracking For composites in which the interfacial residual stress is tensile or zero

and thesliding resistancecan be represented byauniquestress,z, the Nu) relation is given by Equation (5) and the steady-state matrix cracking stress evaluated from Equation (1) is12,14

- - ac - ( 6 f ~ ~ ~ F ~ j-.- (6) E [ I - f ] E E , R E m


where G,, is the fracture energy of the unreinforced matrix and q is the axial residual stress in the matrix. Experiments conducted on a number of ceramic matrix composites are consistent with Equation (6). When the interface is subject to residual compression, r depends on the local applied stress and the solution for a, is more complex. However, to first order, r may be simply replaced by pq.

When the applied stress exceeds a,, multiple matrix cracking is ex- pected12.14 and 0bserved.12~13 The saturation crack spacing D is between one and two times the distance overwhich the applied stress in the matrix builds up from zero at the crack surface to the value for an uncracked composite. For unbonded fibers, this is the distance over which sliding occurs at the interface. In this case, the range of crack spacings is given by1 2

(7)

Experimental observations13 have again confirmed this feature of matrix cracking.

The most crucial aspects of the above interpretation of steady-state cracking and of behavior prediction concern determination of rand q for actual composite systems. Both are difficult to measure. Two basic approaches have been used to measure the sliding resistance z: indentation3031 and measurement of crack opening hysteresis.13 Both approaches are readily applicable when G,, and r are small. The former method is most insightful when used with a nanoindenter system, whereupon rcan be obtained on single fibers either from a push through force on thin sections or from the hysteresis in the loading/unloading cycle on thick sections (Figure 15a). This method has the obvious disadvantage that the fiber is in axial compression so that the interface is also compressed during the test, with attendant changes in r. However,thiseffect has been shown to be negligible for ceramic composites systems having very small sliding stresses (r I 1 OMPa). Calculation of rfrom direct measurement of crack opening hysteresis during load cycling (Figure 15b) avoids thiscomplication because the fibers are subjected to axial tensile loading. However, this approach also has several drawbacks. Measurements are obtained only after matrix cracking and so correspond to a range of crack opening displacements beyond those that dictate formation of the matrix cracks. Furthermore, interpretation of the results is complex when appreciable fiber failure accompanies matrix cracking. Consequently, other approaches applicable to composites having larger 5 are being investigated. One of these is discussed in the following section.

Ultimate Strength Following multiple matrix cracking, the axial stress in each of the fibers

varies from the maximum, equal to o,/f, between the crack surfaces to a minimum, la , E&, halfway between adjacent matrix cracks. The probability and location of fiber failure subject to weakest link statistics in such a stress field can be readily derived. However, calculation of the maximum

Oc R / Z T c D c GcR/f 7


6 a) Indentation Testing

U

P

U Governed t

by Residual Stress b) Tensile Testing

Figure 15. (a)Aload/unload cyclefor nanoindentation ofafiber.(b)Crackopening hysteresis for a composite with intact fibers revealing the trends in both sliding resistance and residual stress.

loadsupported bythecomposite(i.e., a bundleofsuchfibers) requiresthat the stress redistribution caused by the fractured fibers be modelled. Such an analysis has not been attempted. Nevertheless, a lower bound for the maximum load can be derived by simply allowing failed fibers to have no load-bearing ability. Then, a modified bundle failure analysis yields the following expression for the ultimate strength:

Mechanical Behavior of Ceramic Matrix Composites

[ l - ( l - r D / R S ) ]

21

(m + 1)[1- (1 - T D / R S ~ ) ~ ] (8)

-1

[ A m+l 1 A cu = fS exp - where

( R $ / T D ) ~ + = (A,/2RRL)(RSJ.rD)m[l- (1 - L D / & ) ~ ]

with L being the gauge length. In the one composite system for which analysis of the ultimate strength has been performed (LAS/Sic),l129 Equa- tion (8) agrees quite well with measured values.

The ultimate strength anticipated from the above argument is expected to be influenced by the residual stress. Specifically, in systems for which the fiber is subject to residual compression, the axial compression should suppress fiber failure and elevate the ultimate strength to a level exceeding that predicted by Equation 8.32 This effect may be estimated by regarding the matrix as clamping onto the fiber and, thus, simply superposing the residual stress onto S.

Resistance Curves When mode I failure is dominated by propagation of a single matrix

crack accompanied by fiber failure and pullout, the mechanical properties are characterized by a resistance curve. The entire R-curve has been evaluated for fibers with asingle-valued strength, S (i.e., m = a), and small sliding resistance, I: Although this analysis does not account for pullout of broken fibers (f i bers must f ail between the crack su rfaces for m = - ), some useful trendswith microstructural propertiesare evident.The steady-state toughness increase obtained from Equations (2) and (5) is

(9) S 3 R f ( l - f)E; 6 . r E f E , AGC = 2

and the amount of crack extension needed to achieve steady state is

(1 0) *~ - [ GimrR2(1 - f) E, T~ P E:E 5 r

It is noteworthy that both AGc and Ac increase with the ratio R/r, whereas the steady-state matrix cracking stress decreases (Equation 6). Further- more, a simple relationship exists between the steady-state toughening and the matrixcrackingstress(Figure 16).33These resultsindicate thatfor reasonable values of fS/oc( 4 - a,

a c. - 2 2 -

O 1 IT - 2 3

OF

I I I I


bridging fibers and the failed fibers that experience Pullout. Knowledge of the magnitude of the statistical shape Parameter, m, for the fibers within the composite is therefore a prerequisite to Optimizing the shear Properties of the interface for high toughness.

The shape of the rising R-curve is expected to be sensitive to m; in general, the amount of crackgrowth needed to approach steady state must increase with decreasing m. However, the actual slope of the resistance curve has not yet been evaluated, because numerical methodsare needed to determine the upper limit of Equation (2), as dictated by the crack opening at the end of the bridging zone. Further research concerning this phenomenon is a major priority.

Property Transition Nonlinear macroscopic mechanical behavior in tension is most desirable

for structural purposes. Therefore, analysis of the transition between this and the linear response is important. The transition is dependent upon the nature of preexisting defects in the composite, in particularthe length of un- bridged crack However, a useful lower bound, which applies to preexisting defects that are fully bridged, is given by the requirement that the steady-state matrix cracking stress be smaller than the ultimate strength in order to obtain the noncatastrophic failure mode. For m = -, this condition isgiven by setting u, in Equation (4) equal to the fiber strength S.

Equations (6) and (9) for steady-state matrix cracking and asymptotic toughening (at m = -) allow the general trends for property optimization, to be quantified. The variation in strength of a composite with the ratior/R is shown schematically in Figure 17. In the region of noncatastrophic response, the matrix cracking stress increases with the parameter r/R, whereas the ultimate strength is not affected (for small m the ultimate strength is a weakly decreasing function of r). At a critical value of r/R, where a, exceeds Sf, the transition to linear response occurs. With further increase in d R , the fracture toughness decreases (and the strength also decreases if thepreexistingflawsremain unchanged).Therefore,optimum values of r / R exist near the transition point.

The behavior illustrated in Figure 17 has been investigated systematically by heat treating the LAS/SiCcomposite under conditions where the interfacial graphite layer disappears and is replaced by SiO, (Figures 4 to 6), resulting in an increased frictional sliding resistance. The variation in tensile mechanical properties with heat treatment time are shown in Figure 18.The influence of heat treatment on interfacial propertiescan be inferred from measurements of fiber pull-out lengths in broken test pieces and comparison with the calculations of Figure 14.29 The results reveal that the pull-out distribution gradually changes as the gap caused by C removal is filled with Si02 (Figure 19). In particular, the median pullout length decreases and the proportion of fibers that actually pull out exhibit distributions consistent with the predictions of the weakest fiber failure analysis, such that r increases by about an order of magnitude when a partial SiO, layer replaces C. This change in rand the accompanying dramatic change in the mechanical properties of the composite are consistent with the response depicted in Figure 17.


0n I c I 1

1 PEAK STRESS

Sf

0

CRACKING I

I INONCATASTROPHIC I FAILURE1

TIR

Figure 17. Trends in composite properties with r / R for large m.

*p - *D v

I- Deviation from

100 -

0 2 4 20 24

EXPOSURE TIME (h)

Figure 18. Effects of heat treatment on the tensile stress-strain behavior of LAS/ Sic composites.


03 0 Calculation

02

0 1

c o 0 $ 0 4 - L - VI fj 03 r : 0 2 3 cr

U. E 0 1

0

0 5

0 4

03

0 2

0 1

o

I Crack Fronl

' 5 ' 10 ' 15 ' 20 ' 2 5 ' 30' 35' 4 0 ' 45 ' 50 ' 55 ' 60' 65 ' 70 ' 75 ' 100 Pull Out. h ( p m )

Figure 19. Histograms indicating trends in pull-out length with heat treatment.

Residual Stress Large mismatches in thermal expansion coefficient between fiber and

matrixare clearly undesirable. In particular, relatively large matrixcontrac- tion, am/af>>l, causes premature, or even spontaneous, matrix cracking (Equation 4). Such behavior is not necessarilystructurally detrimental, but concerns regarding thermal fatigue, the ingress of environmental fluids, etc., have discouraged the development of materials having these characteristics. Conversely, relatively small matrix contractions, af/am>>l, ther- mally debond the fiber from the matrix. When sufficiently extensive, the resultant radial separations negate the influence of the fibers. Consequent- ly,values of & / a m close to unityare required. Indeed, mode I axial properties subject to an interface that easily debonds and slides freely along the debond involve an optimum residual stress, with amaximum matrixcracking stress, when the interfacial stress is compressive, given by34

(1 2) 1 / 2 O, /E = (213) [ f w m c / ~ ~ , ~ l whereA= 1-(1 -E/~)/2.When~/am>>1,suchthattheinterfacialstressis


tensile, asperities on the debond surface may provide a discrete sliding stress, z, that depends on suchfeaturesas the asperity amplitude. For such cases, the optimum residual strain has not been determined.

The fracture resistance is also influenced by the residual stress. However, the sign and magnitude of the change in toughness induced by residual stress depends on the mechanisms of interfacesliding and fiberfailure, as summarized in Table 2.24 Subject to adequate debonding, the salient result for ceramics reinforced with brittle fibers is that AG,, is unaffected when the interfacial stress is tensile and the interface is characterized bya unique z, whereas AG,, usually decreases with increasing compressive interfacial residual stress because the pull-out lengths decrease, as apparent when z is equated to pq.

Residual stresses in composites are difficult to measure. Even when the composite is fully elastic, so that no interface debonding or sliding occur on cooling, the residual stresses at the surfaces are complex. Consequently, methods such as X-ray diffraction that probe thin surface layers are difficult to interpret. Neutron diffraction, which typically averages over a much largervolume of material, is usually moresatisfactory. Measurement difficulties are exacerbated when debonding andsliding occuron cooling. These processes initiate preferentially at the surface and spread into the body along the interface, thereby alleviating the residual stress over the debond/slip length. For small z and Gic, these lengths are large (many multiples of the fiber diameter).l3 Consequently, a valid measure of the residual stress can only be obtained using processes that penetrate well into the material. One independent approach for measuring q that has merit in some cases involves use of the same crack-opening measurements described in Figure 15b. Specifically, the residual axial stress in the matrix is related directly to the stress at which crack closure occurs.13 The method is, however, restricted to materialsforwhich matrixcracking is not accompanied by extensive fiber failure.

MIXED MODE FAILURE

Mode II Failure Mechanisms Flexural tests performed on uniaxial composites reveal that a shear

damage mechanism exists (Figures 1,20)35 and that such damage often initiates at quite low shear stresses, e.g., 20MPa in LAS/SiC. The damage consistsof en echelon matrix microcracks inclined at about 7~/4 to the fiber axis (Figure 21). With further loading, the microcracks coalesce, causing matrix material to be ejected and resulting in the formation of a discrete mode II crack The crack is defined by the planar zone of ejected matrix. The crackalso has a microcrack damage zonesimilar to that present upon crack initiation.

The microcracks that govern mode I I failure are presumably caused by stress concentrations in the matrix and form normal to the local principal tensile stress, but then deflect parallel to the mode I I plane and coalesce. An adequate model that incorporates the above features has not been


developed. Consequently, the underlying phenomena are briefly noted without elaboration. The stress concentrations in the matrix have magni- tudegoverned by the elastic properties, the fiberspacing and the interface strength. The growth and coalescence of the microcracks are influenced by the matrix toughness GmC The shear strength seemingly decreases as the mode I toughness increases.

Figure 20. Delamination cracking in notched flexure tests.

\ ,

_ _ * _ _ - .-_-e ------ -- - - - 1 Figure 21. Matrix microcracks that precede mode II failure in tough composites.


Delamination Cracking Delamination isacommon damage mode in the presenceof notches13,25

(Figures 1 and 20). Delamination cracks nucleate near the notch baseand extend stably. Analysis of such data is based on the solutions used for mixed mode interface cracking in beamsp modified to take account of the elastic anisotropy. The fracture resistance is found to increase with crack extension. The existence of a resistance curve is attributed to intact fibers within thecrackthat resist thedisplacement of thecracksurfacesand thus shield the crack tip in a manner analogous to fiber bridging in mode 1. However, the fracture energies are typically of the same order as the fracture energy of the matrix, e.g., -20Jm-2 for LAS matrix composite.

INTERFACIAL DEBONDING AND SLIDING

Mechanics of Interfacial Cracks The results and discussion of the preceding sections point to several

problems involving debonding along interfaces that are central to determining mechanical properties of composites. Such debonding occurs both at the tipof amatrixcrackand in thecrackwake(Figure2). It typically involves two materials with different elastic constants and mixed shear and tensile loading. Furthermore, since the interface can have lower fracture resistance than either the matrix orfibers, the debond crack can continue to extend along the interface under mixed mode conditions rather than seeking a plane normal to the principal tensile stress. Therefore, it is necessarytodirect attention to thedependence of fractureresistance, Gi,, on the mix of shear and tensile loading.

The mechanics of cracks at bimaterial interfaces was derived in a series of studies in the 196037-41 and has received renewed interest and elucida- tion recently.42-45 An additional complexity in the fracture mechanics arises from the fact that the shear and tensile components of stress and displacement along thecrack plane aheadof, and behind, thecracktipare not decoupled as they are in linear elastic fracture mechanicsfor homogeneous materials, i.e., a tensile (mode I) remote loading generally results in both tensile and shear stresses and displacements near the crack tip. Moreover, the mix depends on the crack length as well as the mismatch in elastic constants and the position relative to the crack tip. Despite this complication, it is possible to specify the crack tip field in terms of a position-independent stress intensity factor, K, which contains all of the information concerning applied loads and the geometrical configuration. However, to accomplish this, the stresses and displacements are expressed in complex notation, with opening and shear components as the real and imaginaryparts. In thisscheme, the ratioof opening toshearcrack tip displacements(u and v respectively, Figure 22) isdescribed bya phase angle cp = tan-l(v/u).

Because of the interdependence of the opening and shearing compo- nentsof the remoteloading and thecracktipdisplacements,cpdiffersfrom


y . t O E l , v 1 !

X U I . - . - . - . - _ -Je- - - - - - - - --t Interface

+$ _ _ _

Crack Tip Interfacial Crack

0 Ep.V2 4 =tan- (v/u)

Figure 22. The displacement of the surface of a crack at a bimaterial interface indicating the shear and opening displacement that accompany most external loading conditions.

the phase angle of K(Y = tan-l(K,/K,)) by an amount that depends on the mismatch of elastic constants and position:

Q = v + E l n r + tan-% (1 3)

where

I - b E = iwid

and G, (1 - 2 ~ ~ ) -G2 (1 - 2 ~ , )

2 [G, (1- v.J +G2(1 -vJ] b =

where b is one of Dundurs parameters,46 with G the shear modulus, v the Poissons ratio, and r the distance from the crack tip. Because of the In r term,whichdescribesaslowosciIlation in theratiovh withr,thevalueofY is dependent upon the choice of length units. However, this does not present a difficulty provided a consistent choice is maintained.

In most practical examples, the parameter E is small, often less than 0.01.43 Consequently, several schemes for ignoring the effect of E in Equation (13) have been proposed, so that Y represents the relative proportions of mode I I and mode I in the crack tip field.41M However, even in this case, the proportion of mode I to mode II in the crack tip field differs from that in the applied remote field.

The strain energy release rate can be calculated in terms of the crack surface d is p lacem en ts?2

(1 4) x (1 + 4E3 (u + v) = 8r [(1-vl)/G,+(1-v,)/G,]


Alternatively, G can be expressed in terms of the modulus of the stress intensity factor, IKP=K12+K,,2, in a form similar to that for homogeneous maferials?1

(1 5) ClI t, steady state obtains for long debonds and the solutions given in Figure 24 are directly applicable. Fort < t,, the debond crack is subject to normal compression and resultant friction. In this case, G diminishes with increase in debond length, d, representative of stable crack growth:

(1 7) Fd (1 - f) (1 - vF) G/EfR(AaA132 a F2/4+F/2 - R [(I - f) (1 - Zv) + 1 + f] where

F = ( t -q)/E,AaAT


andq istheaxial residualstressin thematrix,asgoverned byAaAT,f,andX. In this instance, G is strictly mode II and debonding should thus be predicted by equating G to G, at Y = d2. Such predictions have not been attempted. However, it is insightful to note that, for weak interfaces (G, K --- X - -X + C X - / X I - - 7 b

(40)

(41)

(42)

(43)

2 2 I 3 J ( x ) = b x - bx + -x 3

2 1 Crn I - c,= (Ef - ci + Z C m ) , ,

in which K, = ( (b L3 )2

G, - = (+ + 5 ) ( + ) q s + (3e - - 2 7 ) - ; [ W k c r 1 , cm - TC T* ] Ts '"j + [ -+-- i k ( z )2] x [ (+kcr)2ls - - 0 ; c r t + (a) 7

a& is the solution of the equation Gs + KxGcr = c,&

Ern E/ Cf ,u

cm 7s

Cf a

(44) + -2(1+ 1 p r n ) ( A ) 2 B ( b ) -s,(~)Ic

bErn

This completes the planar model analysis.


Simplified Planar Model In the preceding treatment the shear and transverse normal stresses in

the matrix are explicitly incorporated into the model. A special case of the above model isalsodeveloped in which the transverse normal stress in the matrix is set to zero so that the shear stress in the matrix is constant for a given coordinate (z). We referto thissimplified model as planar model I and to the model that accounts for the matrix transverse normal stress as planar model 11.

The planar model I is developed by following the procedure given in this chapter's three previous sections. The matrix stress om, in the free body diagrams shown in Figure 4 is assumed to be zero. This leads to

ol,x = 0 and 7,&,~> = 7 m x z ( z > ( 45 )

The equilibrium equation of the fiber only need to be considered to solve for the axial stress in the fiber and the shear stress in the matrix (or the interface), that is, the first line of the Equation (1 8). The axial stress in the matrix is then obtained by the rule of mixtures relation. The stresses in the cracked region and the critical cracking conditions for the three interface conditions considered earlier are obtained as described earlier. The debonding length analysis will be same for planar models I and ll because the lateral stress that inducesdebonding isthecracktiptransversestressand does not depend on the shear lag stress field.

RESULTS AND DISCUSSION

In thissection, we present some typical resultsforaSiC/LAScomposite, using the current planar model. Where possible we will show the results of the current model along side those obtained using the cylindrical BHE model. The material and geometric properties used in obtaining the results are: E, = 8 5 GPa, Ef=200GPa,p,,,= .25,p,= .25,a = 8 pm, K, = 2 0 MPadm, G, = 44 N/m. We present the results for the three interface conditions separately.

Perfectly Bonded Case

Here weassume that thecomposite hasaverystrong interfacesuch that matrixcracking takes placewithout any interface damage. Figure Gshows the distributions of various normalized stress components in the matrix and the axial stress distribution in the fiber. The normalization being carried out is as follows: - OF - 0; a/ =

0,"

Ir - a L - om Om, =

a,"

2

v)

e - 2 - I- v)

n

0 W

-4 w N -J -6- Q z E -8-

-

4 - IOo

- I I I < 0 7 -

< - PERFECTLY BONDED CASE, ~ ~ ~ 0 . 5 -

MATRIX z FIBER

MATRIX

MATRIX CRACK j B cr,i 1 64 I 1 I 16 32 48 -

-


PERFECTLY BONDED CASE, c, =0.5, u= IGPo

LT -

MATRIX

-

-

I 60 80

COORDINATE z (,urn)

Figure 7. Comparison of cylindrical and planar model stress distributions.

12 - I I I 1 PERFECTLY BONDED CASE

h z IO- PLANAR I - WITHOUT TRANSVERSE STRESS - v (3 PLANAR2 - WITH TRANSVERSE STRESS

2 6 = -

8 - - W rx -

-

-

1.0

VOLUME FRACTION, cf

Figure 8. Matrix critical cracking stress-perfectly bonded case.

Micromechanical Models for Brittle Matrix Composites 59

Slipping Case Forfrictionally constrained fibers the predicted matrixcracking stress is

presented in Figure 9. For the interface frictional strength (z,) a value of 2 MPa and for the frictional coefficient (K,) a value of 0.2 have been used in the present calculations. As in the perfectly bonded case, the current models predict a trend similar to the cylindrical model and lower critical stresses.

1.2- I I I I

SLIPPING CASE, r = 2MPa h z 1.0- PLANAR I - WITHOUT TRANSVERSE STRESS - 0 PLANAR 2 - WITH TRANSVERSE STRESS 0 0.8 - cn W cr h 0.6- - -I 3 0.4- -

Y * EXPERIMENTAL C43 -

E g 0.2- -

OO 0.2 0.4 0.6 0.8 1.0 VOLUME FRACTION, c f

Figure 9. Matrix critical cracking stress-slipping case.

Initially Bonded-Debonded Case Finally, we present results in which the initially bonded fiber-matrix

interface debonds as the matrix crack approaches it. The results for this case are interesting in that the planar model predicts a unique critical stress for a given fiber volume fraction, Figure 10, and a unique debond length for a given critical debonding energy release rate (Gd), Figure 1 1. In Figure 11 the debond length is same for the two planar models. In both cases the predicted values of the planar models seem to agree with corresponding lower branches of the double-valued functions of the cylindrical model over a significant range of the abscissa.

CONCLUSIONS

The Nicalon-LAS unidirectional composite system is one of the few which has been characterized completely, aside from the bonded-debonding mode, in accordance with the present models.These input values have


been given earlier here. For a fiber volume ratio of 0.5 and the estimated frictional resistance of 2 MPa determined from push-out and indentation tests, the matrixcracking stresswascalculated to be 265 MPa bythe BHE model neglecting initial stresses for the large slip case. The experimental observation for the onset of nonlinearity in a direct tension test was 290 f 20 MPa. Despite this apparently reasonable agreement between theory and experiment, we cannot conclude that predictive modeling of the so- called steady state cracking mode in BMC has been validated. Neither the planar nor cylindrical models described in this chapter give reasonable results for the limiting cases corresponding to zero fiber volume ratio and perfect bonding. Both models produce a zero cracking strength in the former case ratherthan the monolithicmatrix result. Forthecaseof perfect bonding both models predict extremely high cracking strength, especially the BHE model.

For example, the BHE results for the aforementioned Nicalon-LAS compositewith perfect bonding is a matrixcracking strength of 1625 MPa, neglecting initial stresses. While we cannot prove this result is incorrect, the prediction lacks credibility. With increasing fiber volume ratio, the model prediction becomes even more incredulous as shown in Figure 8. Thustheapparentagreementatasingledatapoint may bemorefortuitous than representative of a rational understanding of the failure process.

To amplify the point just made in the previous paragraph, experimental observations by Kim4 fora Nicalon-LAS unidirectional composite indicate approximately the same point of initiation of nonlinearity as above. How- ever, the corresponding damage mode consisted of a rather diffuse distribution of matrix and fiber transverse cracks, as well as longitudinal splitting of the matrix. In fact, numerous other experiments on Nicalon fibers reinforced by other glass-ceramic matrices failed to produce the steady state cracking mode at first damage. Unfortunately, Nicalon fiber- glass ceramic matrix composites are far from ideal in the sense that the contain many irregularities in fiber size, spacing, and orientation. Therefore, a convincing demonstration that the concept of steady state cracking in BMC corresponds to an event of practical significance in realistic composites remains to be seen. It seems evident that under certain conditions of stiff and strong fibers reinforced by a weak and extremely brittle matrix this failure mode will prevail as shown in earlier work by Aveston, Cooper, and Kelly.* However, we cannot claim that the phenomenon is well under- stood at this time.

With regard to the comparison between cylindrical and planar models, we have observed considerable differences between the two in the examples treated here, although in many cases the trends are qualitatively similar. The most striking departure is demonstrated in Figures 10 and 1 1. Hereadouble-valued result isgiven by BHE,while the planar model yields a monotonic dependence. It should be noted that the lower branch of the BHEcurvein Figure1OcorrespondstotheupperbranchinFigurell.Thus the larger predicted matrix cracking stress leads to the smaller debond length. From this observation we can conclude that both states are not physically possible within asingle specimen.The authors3 favor the upper

Micromechanical Models for Brittle Matrix Composites 61

IO I I I I Gd/G, = 0.1 BONDED DEBONDING CASE, h B s? 8 - -

3 v)

IT 6 - - I- cn

-

IT 0 2 - -

1.0

VOLUME FRACTION, cf

Figure 10. Matrix critical cracking stress-bonded-debonding case.

branch of Figure 11. Forthe higherfibervolumefractions,at leastforG,,/G, = 0.1, reasonable agreement between the two model predictionsof matrix cracking stress is observed along this branch. However, at lower volume fractionsthisagreement is totally lacking. Furthermore, the debond length predictionsof Figure 1 1 differdrastically,especiallywith the upper branch of the BHE curve. It is not clear at this time whether the double-valued dependence noted in Figure 1 0 and 1 1 for the BHE model (but not for the planar model) representspotentialfailurestatesorsimplyanartifact ofthe modeling assumptions.

The most favorable agreement between the planar and cylindrical models iswhen the transverse stress is neglected, and the results of planar model II differ significantly from those of BHE. Since a, is neglected in the cylindrical model, this trend may not be meaningful except to indicate that planar vs cylindrical correlation might be achieved by varying the magni- tudeof thetransversestressin theplanarmodel.Thus, byconsidering axas a parameter, we may be able to bring the two models intocloser agreement. This may be of interest from an empirical viewpoint, but it is not satisfying academically or realistically since the effect of transverse stress on the bond strength has been ignored.

In general, we must admit that the model comparison has been somewhat disappointing with regard to application of planar modeling concepts to represent micromechanical relations. While the BHE cylindrical model has definite shortcomings, one cannot reconcile that using a cruder approximate representative of the geometry can lead to a more valid prediction. Thus it appears that application of planar models in composite micro- mechanics must be employed with caution with the appreciation that they have limited fundamental value. Their use in trend studies and semi- empirical calculations may be helpful when guided by appropriate experiments so long as one does not extrapolate too far from the experimental conditions. As mentioned earlier, only through systematic modeling and

4

e I' 3- - i i

I- (3 z W -I 2 -

n z w D g

1 I I BONDED DEBONDING CASE, ct = 0.5

-

-

I -

PLANAR

I I - '

Fiber_Reinforced_Ceramic_Composites/Fiber Reinforced Ceramic Composites/12333_03.pdf3

Properties of Ceramic Fibers from Organosilicon Polymers

Neal R. Langley, Gary E. LeGrow and Jonathan Lipowitz

SYNTHETIC METHODS FOR PRECERAMIC POLYMERS AND CERAMIC FIBERS

Introduction The use of preceramic polymers, which are processed as ordinary

polymers to the desired shape and then pyrolyzed to ceramic parts, can overcome the difficulty of high temperature shaping of ceramics.

The organometallic polymer route to ceramic materials has several advantages, including shaping and forming of parts, high purity, and the abilityto produce metastable compositions. Organometallic polymers can be prepared in high purity. The purity of the polymer is determined by the purityof thestarting chemicalsand theenvironment ofthesynthesis, both its atmosphere and its material of construction. Generally, the purer the starting materials, the purer and more consistent the polymer. Organo- metallic polymers of awide rangeof compositions can be produced which can correspondingly produce a wide range of ceramic materials, also with a wide range of properties.

Practical preparation of ceramic fibers via the organometallic polymer route requires several important polymer properties and process characteristics. The polymer must be tractable with a softening point above ambient temperature. The rheological characteristics of the polymer should permit softening of the polymer, extrusion, and draw down to form small- diameter fibers without any degradation of the polymer's physical or chemical properties. Once the polymer is in the form of a fiber, it must be able to be rendered infusible while in the solid state so that on heating above the softening point of the original polymer, the fiber will retain its shape. Further heating usually causes significant weight loss until a

63


temperature of approximately 800C is reached. Above 800"C, negligible weight loss occurs; however, densification occurs with isotropic shrinkage, finally producing maximum tensile strength and modulus at 11 00 to 1400C. At higher temperatures, thermal degradation typically begins.

Polymeric Precursors for Silicon Carbide Ceramic Fibers The first demonstrated preparation of silicon carbide containing ceramic

fibers from a polymeric percursor was reported in 1975 by Yajima and co- workers.1 The process, shown in Equations 1 through 4, is based upon permethylpolysilanes, produced from dimethyldichlorosilanes by a Wurtz coupling reaction with sodium.2 Thermal decomposition of the permethylpolysilanes at 350"C, catalyzed by3 to 5 percent polyborodiphenylsiloxane, yields a polycarbosilane.3 The polycarbosilane is then fractionated to remove both higher and lower molecular weight fractions. The midrange fraction, with a number-average molecular weight of 1000 to 2000, is suitable for melt spinning into 10 to 20 pm diameter fibers. The fibers are then crosslinked by oxidation of Si-H bonds by heating in a i r ( l l 0 to 190"C), and finally pyrolyzed to 1 100 to 1200C in an inert atmosphere to produce blackceramic fibers.4 This technology has been commercialized by Nippon Carbon Co. (Japan). Silicon carbide containing ceramic fibers is supplied under the trade mark NICALON.

x MezSiCIz + 2x Na + (Me,Si),

(Me,Si), + heat (350%) + (MeHSiCHZ),

(MeHSiCH2)x + air -+ (Me0KSiCH2),

(MeOxSiCH2), + heat (12OO0C) -+ Si + Cceramic

+ 2x NaCl (1)

(2)

(3)

(4)

A chemical modification of polycarbosilane was produced by reaction with an alkoxy titanate.5 Based upon infrared spectral evidence, the titanium appears to be bonded through oxygen to silicon. Ceramic fibers were produced from this type of modified polycarbosilanewhich contain 6 to 15 atom% titanium relativetosilicon.Thesefibersalsocontainasmuch as 50 atom % oxygen relative to silicon. This technology has been commercialized by Ube Industries, Ltd. (Japan). The silicon-carbon-oxygen-titanium- containing ceramic fibers are supplied under the trade name of TYRAN NO FIBER.6

Another method of preparation of a polycarbosilane, by passing various simple organosilicon compounds such as Me,Si through a hot tube at 700C, was patented by Bayer AG in 1974.7 The polycarbosilane polymer produced by this method was dry spun, using methylene chloride as a solvent, into 10 to 20 pm diameter fibers which were crosslinked by

Properties of Ceramic Fibers from Organosilicon Polymers 65

heating inair(200C),andthen pyrolyzedtol200Cinan inertatmosphere to produce black ceramic fibers7 Although Verbeeks patent preceded Yajimas patent by one year, commercialization of this technology by Bayer AG has apparently not been attempted.

Branched polycarbosilanes have been prepared by a one-step process of reacting chloromethyl- or vinylsilanes with other methylchlorosilanes using potassium.8.9 These branched polycarbosilanes were converted to SIC containing ceramic materials by inert atmospheric pyrolysis10

Organopolysilane copolymers have been synthesized by alkali metal dechlorination of mixtures of diorganodichlorosilanes.11 Polysilastyrene is produced (Equation 5) from a 1/1 ratio of dimethyldichlorosilane and methylphenyldichlorosilane.1~ Polysilastyrene was drawn into fibers which were then crosslinked by UV radiation and then pyrolyzed in a vacuum to 11 00C producing Sic containing ceramic fibers.13

x Me2SiC12 + x MePhSiC12 + 4x Na -+ (5) (Me,SiMePhSi), + 4x NaCl

Branched poly(methylchlorosilanes) have been produced (Equation 6) by Bu,PBr catalyzed redistribution of a mixture of methylchlorodisilanes.14 These polymers were derivatized (Equation 7) with various organometallic reagentssuch asGrignard Reagents to replacethesi-Clfunctionalitywith Si-R groups. These Si, C, and H-containing polymers have been spun into fibers of 15 to 20 pm diameter, cured, and pyrolyzed to 1200% producing ceramic fibers of a wide range of composition from Sic with excess silicon through stoichiometric Sic to SIC with excess carbon.15

x MeC12SiSiC12Me + Bu4PBr cat. .+ (MeSi),Cly + x MeSiCI3

(MeSi),CI, + y RMgCl -+ (MeSi).R, + y MgC12

(6)

(7)

Polymeric Precursors for Silicon Nitride Ceramic Fibers Dichlorosilane has been reacted with ammonia (Equation 8) to produce

silazane polymers.16-18 Pyrolysis of these polymers to 1200C led to ceramic materials having a composition between Si,N4 and Silicon.19

H2SiC12 + xs NH3 + (H2SiNH), + NH4CI (8)

Polymer Precursors of Silicon, Carbon, and Nitrogen Ceramic Fibers Organotrichlorosilanes have been reacted with ammonia to produce

organosilsesquiazanes,20,21 as shown in Equation 9. Methylsilsesquiazane was dry spin into fibers and pyrolyzed to 1200C in nitrogen to produce an amorphous ceramic fiber of a composition between silicon carbide and silicon nifride.20

(9) x RSiCI3 + xs NH3 .+ [RSi(NH)3/21x -+ x NH4CI


Methyltrichlorosilane was reacted with methylamine (Equation 10) to produce methyltris(methylamino)silane. This compound, when heated to 520C, undergoes a thermal rearrangement producing a polycarbosilazane polymer.2,22 Verbeek also described the synthesis of an analogous polycarbosilazane from the reaction of dimethyldichlorosilane and methylamine? The silazane derived from methyltrichlorosilane was melt spun intofibersof 12 to 14pmdiameter,exposed tomoistureforcrosslink- ing, and then pyrolyzed to 1200C in nitrogen. The resultant ceramic fiber had a composition between silicon carbide and silicon nitride.

MeSiC13 + 6 MeNH, -+ MeSi(NHMe)3 + 3 MeNH,CI (1 0)

Direct ammonolysis of methyltrichlorosilane, trichlorosilane, or methylchlorodisilanes produces the corresponding silsesquiazanes; however, there is a tendency to form gels particularly in the case of trichlorosilane. Thistendencycan be partlyalleviated by reacting thetrichlorosilaneswith hexamethyldisilazane. In this manner Si-CI/Si-N redistribution (Equations 1 1-1 2) produces Me3SiCI and incorporates carbon in the form of Me3SiNH- substituents into the silsesquiazane polymers.23-25 These polymers have been melt spun into fibers, cured by exposure to HSiCI,, and then pyrolyzed to 1000 to 1 200C to produceceramic fiber scomposed primarilyof silicon and nitrogen, but containing 2 to 10 percent carbon and ~2 percent oxyge n.26

RSiC13 + (Me3Si),NH + RSi(NHSiMe3)3 + Me3SiCI (1 1)

(Me3SiI2NH (1 2) x RSi (NHSiMe3I3 + [RSi(NH)3/21x(NHSiMe3), + -

Direct ammonolysis of methyldichlorosilane produced (Equation 13) a mixture of cyclic silazanes. Treatment of these cyclics with potassium hybride (Equation 14) resulted in a dehydrogenation reaction and formation of a branched polysilazane.27 Pyrolysis of these polysilazanes to 1000C in nitrogen produced ceramic materials containing silicon, carbon, and nitrogen.

3 x - y 2

x MeHSiCI, + xs NH3 -+ [MeHSi(NH)I x + NH4CI

[MeHSi(NH)], + KH -+ [MeHSi(NH)IY (MeSiN), + Hz

(1 3)

(1 4)

The reaction of methylamine with dichlorosilanezs and organodichloro- silanes29 has produced (Equation 15) N-methylpolysilazanes. Ruthenium carbonyl has been employed as a catalyst for a dehydrocoupling poly- merization (Equation 16) of these polymers which leads to an increased ceramic yield.30 Fibers have been obtained from some of these polymers which werecuredwith ammoniaand then pyrolyzed toform ceramicfibers composed of silicon, carbon, and nitrogen.29

x RHSiCI, + xs MeNH, -+ [RHSi(NH)I, + MeNH,CI (1 5)


Aphilosophyused inourcharacterizationworkhas been toutilizealarge variety of techniques providing a wide range of compositional and structural information. Those techniques which prove useful are retained and extensively utilized. This approach leads to a comprehensive picture of structure and composition since the complimentary nature of various techniques provides a rather complete and detailed picture of compositional and structural features. SEM, microprobe, and analytical TEM techniques will not be discussed since they will be presented in another chapter of this volu me.31

Though many ceramic polymer precursors have been described in the literature, relativelyfew have been spun intofibers,cured, and pyrolyzedto form ceramic fibers. None have previously been thoroughlycharacterized with respect tocomposition andstructure. In general, X-raydiffractograms have been run and in some cases elemental analyses were obtained. If crystallinity was observed, crystallite size was sometimes determined by the Scherrer line-broadening technique. Characterization rarely proceeded beyond this stage, except for NICALON fibers.34-36,61,64

Verbeek et al.20,32 were the first to report preparation and properties of ceramic fibers prepared by pyrolysis of polymeric precursors, which were polysilazanes in theirwork Fibers wereamorphous toX-rayswhen pyrolyzed to 1200C and their compositions were reported as homogeneous mixtures of the species Si3N4, Sic, and C. The compositions given were presumably based on elemental analysis, though analysis methods were not described. Oxygen was probably also present since fibers were air cured.At higher pyrolysistemperatures(1600 to2000C),p-SiC andp-Si3N4 were identified by x-ray diffraction (XRD).

The earliest reports of Yajima et al.133 describe a microcrystalline p-Sic fiber of -10 nm grain size based on XRD. High tensile strength and modulus (up to 600 GPa) suggest a microcrystalline SIC fiber as does invariant strength with temperature up to 1400C. The cure method is not described. A later publication3 describes air cure and vacuum pyrolysis. The structure is claimed to be a uniform cohesion of fine p-Sic particles, the crystal structure having no orientation. The fiber after air cure gave much lower modulus values than described in earlier publications. A later publication34shows that -10 wt%oxygen is present in theceramicsi-C-0 fiber as a result of air cure. It also shows that tensile strength begins to decrease above 1200C ascrystallization of p-Sic proceeds. It was theorized that oxygen is concentrated mainly at and near the fiber surface. High resolution, lattice-imaging TEM showed that regions of crystalline Sic, 0.25 nm (1 11) plane spacing, and graphitic-like C, 0.34 nm (002) plane spacing, are present in fibers heated to 1300 in vacuum.34 Some crystallization and thermal decomposition is expected in vacuum at 1300C with loss of CO and Si0 and growth of p-Sic crystallites(seeThermal Stability section).

Others3536 haveattempted tocharacterize thestructureof theseSi-C-0 fibers prior to thermal stability studies. One difficulty in characterization has been that changes in processing methods as fibers proceeded from laboratorysynthesis to commercial production (NICALON fibers, Nippon Carbon Co.) appears to have caused changes in composition and structure.35


The elemental composition found by Andersson and Warren35 was similar to that reported for standard grade NICALON" (SGN), consisting of about 15 wt%oxygen.37Theoxygen(andSiandC) isdistributed uniformlyacross the fiber cross section by electron microprobe analysis.36 Fibers are essen- tially amorphous to a pyrolysis temperature of about 1200C. Nanometer sizecrystallitesofp-Sicaredetected byX-ray line broadeningat aboutthis temperature. Crystallite size increases with increasing pyrolysis temperature.

Composition In ourearlyworkon characterization of ceramicsderivedfrom polymers,

it was recognized that accurate methods of total elemental analysis are required to obtain compositional information because many polymer- derived ceramic materials are amorphous (or partially amorphous) to X-rays and nonstoichiometric in composition. Analysis for C, N, and H using methods developed for organic or organosilicon materials often give low total analyses because refractory silicon carbide or nitride is present. Recommended procedures37 follow.

Silicon-Sodium peroxide fusion and analysis by atomic ad- sorption spectrometry38 or by an inductively coupled argon plasma. Trace metallic elements may also be determined by this method. Carbon, hydrogen, nitrogen-A high temperature combustion method is used (Carlo Erba Elemental Analyzer Model 1 106). Oxygen-A high temperature carbothermic reduction method is used (Leco Corp. Oxygen Determinator Model R0316). Chlorine-Silver nitrate titration of an aqueous solution of material from the sodium peroxide fusion.

These methods give good accuracy for high purity silicon nitride and silicon carbide standards and generally give analyses totalling 95 to 102 wt% for polymer-derived ceramic fibers. Elemental analyses of some powdered ceramic fibers are shown in Table 1.

Table 1 : Typical Elemental Analyses in W t %

Ceramic F iber S i C N 0 Tota la

M P S 6 8 . 3 27.1 0 0 . 9 9 6 . 3 CG NICALON 58. 4c 31 .2 0 . 1 10.0 99 .7 SG NICALON 56.2' 3 0 . 4 0 . 1 14.2 100.9

Ube Tyranno 49 .3 27.9 0 . 3 2 0 . 2 98. gb

MPDZ-PhVi 4 7 . 0 28.9 14 .5 7 . 5 97 .9 HP Z 59 .4 10.1 28 .9 3 . 1 101.5

a Also A. u s u a l l y 0.1 wt%; C 1 . u s u a l l y O.lwt% (not i n SGN or CGN).

b Inc ludes 1 . 2 w t X T i . c C G , Ceramic Grade. SG, Standard Grade.


Fibers we have characterized include an old standard grade of commercial Si-C-0 NICALON" Fiber(Nippon CarbonCo., Ltd.) containing-l5wt% oxygen, referred to as SGN, and a newer ceramic grade NICALON" fiber containing -1 0 wt% oxygen, referred to as CGN.39-41 A Si-C-N-0 ceramic fiber derived from methylpolydisilylazane (MPDZ) polymer, referred to as MPDZ fiber,W5 and a Si-N-C ceramic fiber derived from hydridopolysilazane (HPZ) polymer, referred to as HPZ fiber,24,26 were characterized. A Si-C-0 fiber derived from methylpolysilane polymer, referred to as MPS fiber,14,42,43 was also characterized. The latter three fibers were developed at Dow Corning under a government (DARPA) sponsored contract "Advanced Ceramics Based on Polymer Processing" (Air Force Contract F33615-83-

Nominal rule-of-mixture compositions can be calculated by assigning all oxygen tosilicon as SiO,; then all nitrogen to silicon as Si,N4; then carbon to silicon as Sic. Excess carbon or silicon, if any, is considered to be in the elemental standard state. These rules are obtained using increasing Gibbs free energies of formation per equivalent of silicon for the species SiO,, Si3N4, and SiC.44 Composition for the Tyranno Si-C-0-Ti fiber (Ube Industries, Ltd.)W was not calculated. Compositions are shown in Table 2 as mole % or equivalent % for excess C or Si and as wt% in parentheses.

It is apparent from X-ray diffraction that actual compositions are not mixtures of these thermodynamically stable crystalline species(Figure 2). XRD of Tyranno fiber also shows a primarily amorphous structure at pyrolysis temperatures up to 1300"C.s

The XRD data in Figure 2 is shown for ceramic fibers pyrolyzed to less than 1400C. At higher temperatures, many of the fibers undergo crystallization with simultaneous loss of gases such as CO, SiO, and N,. These phenomenawill be described in asubsequent section on thermal stability. The Si-C-0 ceramic fibers show some evidence for microcrystallinity. Broad peakspresentat28=36"areassignedtothe[lll] reflectionofp-Sic. Crystallinity can also be detected by darkfield TEM and lattice imaging TEM techniques.3134 Additional broad peaks at 28 = -60" and 70" can be assigned to [221] and [311] reflections ofp-Sic.

C-5006).

Table 2: Typical Rule of Mixtures Compositions in Mol % (Wt %)

MPDZ-PhVi H P Z SGN CGN MPS

Si02 8.0 (14.3) 6.8 (5.8) 15.1 (26.8) 10.9 (19.1) 1.1 (1.8)

Si3N4 8.9 (37.1) 35.5 (71.3) 0 0 0

SIC 22.9 (27.2) 32.3 (18.5) 52.1 (61.6) 60.6 (70.9) 85.8 (94.0)

C 60.2 (21.4) 25.5 (4.4) 32.8 (11.6) 28.5 (10.0) 0

si 0 0 0 0 13.1 (4.3)

\ Intensity - - s G L L _ d *

MPDZ - PhVi ----- W+IUh w

HPZ

'--*-A

I I I I


Chemical Bonding Infrared spectrawere obtained on ceramic fibers as potassium bromide

disks (Figure 3). This technique eliminates all absorption bands due to nonceramic components. Note that a strong band covering the region from 700 to 1200 cm-1 is present. This band is due to SiOSi bonds at 1 180 cm-1, SiNSi bands at 950 cm-1, and SiCSi at 830 cm-1.47 Veryweak bands at -2950 cm-1 are assigned to traces of C-H bonds remaining in the ceramic. Sharp crystalline bands for Sic, Si,N,, and Si02 which appear below 800 cm-1 are not present. InfraredspectraforTyrannoceramicfiber are similar to CGN and SGN spectra.48

Transmission

Wave Number (cm-')

Figure 3. Infrared spectra of powdered ceramic fibers.

Solid state, magic angle sample spinning nuclear magnetic resonance spectroscopy (MAS NMR) has proven to be a powerful technique for determining details of the chemical bonding present in these fibers.37,49,50 Strong evidence is found for the presence of a glassy phase in all these fibers, which consists of mixed bonding of carbon and oxygen atoms to silicon in NICALON" and MPS fibers, Si(C,O),, and of mixed bonding of C, N, and 0 atoms to silicon in MPDZ-PhVi and HPZ fibers, Si(C,N,O),.

Figures 4 to 8 show9 Si MAS NMR spectra of the ceramic fibers. In the Si-C-0 fibers, there are five possible tetrahedral structures for the silicon atoms if only Si-C and Si-0 bonds are present as suggested by the infrared spectra. The tetrahedral structures are Sic4, SiC,O, SiC202, SiCO,, and SiO,.AIl appear to be present in the NICALON" fibers. Oxygenated tetra-

I

intensity

I

SIC


-46

Intensity

50 0 -50 -100 -150 -200 PPM Chemical Shilt, 6

Figure 8. 29Si MAS NMR spectrum of HPZ ceramic fiber.

Table 4: Model Compounds for 29Si NMR Chemical Shifts of Si(C,0l4 Ceramicsa

Si Tetrahedral Structure Compound 6 .PPm

sio4 Si (OS iMe 3) 4 -105 Si (OMe) 4 70

s io4 Si02(1) -114.0 Si03C MeSi(OSiNe3) 3 - 65.0

- 41.5 3 sio3c Me S i (OMe) sio2c2 Me2Si(OSiMe3) 2 - 22.1

Me2Si(OMe)2 - 2.5 siocg Me 3 S iOMe + 17.2 sioc3 (Me3SiI20 + 6.7

b s ic4 68-Sic -13.9, - 20.2, -24.5 b $-sic - 18.3

Me4Si 0

a ref. 53

b ref. 51, 52


hedral structures are more prevalent in SGN than in CGN, reflecting the higher oxygen concentration in SGN. Chemical shifts are assigned based on known shifts for the various silicon carbide p0lytypes51~52 and for organosilicon compounds with mixed, Si(C,O),, tetrahedral structures.53 Model compounds used are shown in Table 4. Note that the most intense signal at6=-11 to-15ppmfromSiCOfibersoccursclosetothechemical shift of pSiC (a=-18.3 ppm) as predicted by rule-of-mixtures calculation showing Sic to be the dominant species and by X-ray diffraction which shows the beta (3C) structure to be predominant. The 2QSi MAS NMR spectrum for MPS fiber is rich in Sic, structure but contains no SiCO, or SiO, structure. This occurs because MPS fiber contains only 6 wt% oxygen.

It is important tospin thesample rapidly(>4 KHz) to minimize interference from spinning side bands and to allow sufficient time between signal acquisitions(21 minute) to permit slow spin-lattice relaxation(T,) tooccur. At least 200 signals were acquired to reduce the signal to noise ratio(S/N) to acceptable values.

Each possible tetrahedral structure about silicon is broadened by the well known distribution in bond angles and bond lengths in amorphous structures54 as well as by second nearest neighbor effects. There are 15 possible tetrahedral structures (SiC,NbO, with a+b+c=4) in the Si-C-N-0 ceramicsystem represented by MPDZ-PhVi and HPZ.A broad, unresolved envelope covering this chemical shift region is thus expected and observed. Thespectrum for MPDZ-PhVi isenrichedin tetrahedralspeciesin the high- field (SiO,, SiCO,) and low-field (SiC30) chemical shift regions because of its high oxygen and carbon content. Chemical shifts of the oxygenated species cover a broader chemical shift range than do the Si(C,N), model compoundsshown inTable 5.Chemical shiftsof compoundswithC, N,and 0 bonded to silicon are not available.

Table 5: Chemical Shifts of Model Si(C,N), Compoundsa

Tetrahedra l S truc ture Compound 6. PPm

+ 5 .9 SiC3N Me3SiNMe 2 SiC2N2 M+=2Sf(NMe2)2 - 1 .7

SiN4 Si (Me2l4 -28.1

SIN4 B-Si3N4 -48.7

sicN3 MeSi(NMe2)3 -17.5

a-Si N b 3 4 -46.8. -48.9 SIN4

b

b amorphous-Si N - 4 6 . 4 3 4 S i N 4

a ref . 53

b ref. 55

100 -

Intensity (normalized)

50 -

0


Table 6: Chemical Species Distribution by Si-2P XPS (Mol %, Powdered Sample)

Fiber si-0 si-x* si-c

CGN 7 31 62 SGN 12 39 50 MFDZ-PhVi 7 50 43 mz + 3 38 5 1

* Si-X i s mixed spec ie s . + 7 m o l % S I - S i spec ies or S i with bond vacancy is a l so present.

Nanostructure

Nanometer-scale crystallites of p-Sic (nanocrystallinity) weredescribed in Si-C-0 based ceramics in an earliersection on composition.TEM datawill be presentedelsewhere in thisvolume for NICALON fibers.31 Good agreement is obtained between crystallite sizes calculated by X-ray line broadening and by dark-field TEM measurements.

Another nanoscale heterogeneity in these ceramics appears to be excess carbon. We describe this elemental carbon which is calculated by the rule-of-mixtures method as excesscarbon rather than freecarbon, the more usual term, because it is not readily available or free for oxidation. By using a conventional oxidative analysis method for measuring free carbon in silicon carbidep only 2 percent of the calculated elemental carbon can be found. This method corrects for oxidation of Sic during oxidation of elemental carbon by simultaneous measurement of weight loss and CO, evolved. If extremely fine grinding and oxidation for 6 hours rather than 15 minutes as specified in the test method is employed, up to two-thirds of the expected carbon can be obtained. The elemental carbon thus appears to be present in extremely finely divided form, perhaps surrounded by amorphous material, such that grinding will not preferentially fracture the sample at large carbon grains as in conventional ceramics containing free carbon. Thus the carbon is oxidized with difficulty.

Raman spectroscopyisa useful method toobtain additional information about elemental carbon. Raman spectra of ceramic fibers are shown in Figure 10. The Raman signals were too weakfor relative intensity measurements. Fibers were finely powdered to give a representative bulk analysis. The onlyfeatures observed are two broad bands at 1575 to 161 Ocm-1 and at 1350cm-.The bandsare notwellseparated in HPZfibers.These bands are very similar to those observed for pyrolytic carbons (PyC) which have been prepared at similar temperatures to these ceramic fibers (1 100- 1400C). Such materials are pyrolysis or combustion products of pitch and various organic compou nds.60-63

The 1575 cm-1 band is assigned to the graphitic(po1yaromatic) structure. It shifts toward higher frequency with extremely small crystallite size60 or with a greateramount of disorderedcarbon.60-63 Vidano and Fischback63 claim that agraphitic band occurs at 1580cm-1 and adisorderedcarbon


f *\ I \

Absorption

-----/-

I 1800 1600 1400 1200 1000 800 600 400

Wave Number (CW)

Figure 10. Ramanspectraof powdered ceramicfibers. Reprinted by permissionof the American Ceramic Society.36

band at 1620 cm-1. The 1350 cm-1 band is due to unorganized carbon and increases with decreasing crystallite size. From the intensity ratio of the twopeaks,an approximatecrystallitesize(in nm) in thegraphiticaaxis (in the polyaromatic plane) can be calculated as follows:60

La = 4.5 (115,5/11350)

Approximate crystallite domain sizes calculated in this way are % 4 nm for all the ceramics examined. The c-axis dimensions (0002 plane) must be small, no more than a few polyaromatic layers in extent, because of the inability to determine the graphite structures by XRD (20=%25 for the 0002 reflection). A turbostratic structure of the polyaromatic sheets, as is found in pyrolytic carbons pyrolyzed at these temperatures, could also serve to weaken the (0002) reflection.

Unfortunately, the Raman intensity of these graphitic bands is so low duetoself-absorption that it has not been possible to use their intensityas a measure of the amount of excess carbon present.

Sasaki et al.64 have recently observed similar Raman spectra for NICA- LON fiber pyrolyzed to 1400C. The graphiticstructure is less developed at lower pyrolysis temperatures. Above 1 400C, a Raman spectrum of crystalline silicon carbide begins to develop as gases such as CO and Si0 are lost and larger crystallites of silicon carbide develop. These phenomena are discussed in detail in the section on thermal stability.


Yajima et al.34 have detected the 0002 graphite spacing (0.34 nm) by lattice imaging TEM of NICALON" fibers that have been heated at 1300C in vacuum. Crystallization of large p-Sic crystallites under these conditions, along with loss of CO and SiO, may promote separation of largergraphitic crystallites between the Sic grains.

Porosity Apparent densities (p) of ceramic fibers have been measured by a

gradient column method.65 Low and high density solvents used were perchloroethylene and bromoform respectively, which cover a density range of 1.9 to 2.7 g/cm3. Theoretical densities (pt) were calculated from rule-of-mixtures compositions using a volume additivity rule.

p i = density of component i Pt = ?PiVi

I Vi = volume fraction of component i Component densities (g/cm3) used were?7 SiOJI), 2.2; Si3N4(l), 3.0;

SiC(I), 3.0; p-Sic, 3.2; and Pyc, 2.2 (I = liquid, or amorphous). Gradient column densities and calculated pore volume fractions are

shown in Table 7. Densitiesfor the Si-C-0 fibers, which contain p-Sic nanocrystallinity, were calculated using amorphous and crystalline densities are required, using the wt% crystallinities shown in Table 3. Note that the calculated porosities are substantial. However, various methods used to detect porosity66 indicate that little open porosity is present. For example, mercury porosimetry measurements show little open porosity. Helium pycnometry measurements give densities only a few percent higher than gradient column densities. Surface area measurements by mercury porosimetry and by the BET nitrogen or krypton methods give values close to calculated for fibers of the diameter used. Microscopy measurements (SEM, TEM) show little visible porosity. However, preliminaryx-rayscatter- ing measurements show that considerable porosity is present in the nanometer size range.67 Treatment of this scattering data by standard techniquessuch as the Guinier method and the Porod method66showthat the nanometer-scale closed porosity is present in all these ceramicfibers.'

Table 7: Density and Porosity of Ceramic Fibers

Density: p Porosity! P

MPS 2.62 0.10 CG N I C A L O N 2.55 0.16 SG NILAmV 2 s 2 0,_16 HPZ 2.32 0.20 MPDZ-PhVi 2.18 0.18

a gradient column method

b p = 1 - (P/P,) 7 i"i p t = theoretical rule-of-mixtures density; p t =


This nanoporosity is believed to arise during the pyrolysis process, in which hundreds of volumes of pyrolysisgases are evolved per unit volume of cured polymer undergoing pyrolysis.37 It is possible that open porosity developsduring the pyrolysis process to permit escape of pyrolysisgases. This open porosity may subsequently close after pyrolysis gases have evolved(above800"C). In thetemperature range800to 1200"C,densityof the ceramic fiber continues to increase as does tensile strength and modulus. Open pores may undergo closure to the globular nanopores found byx-ray scattering asa result of aviscousflow processof theglassy phase.

The closed nanoporosity present in these polymer-derived ceramic fibers is being characterized more extensively.

Discussion of Amorphous Structure Amorphous structures appear to be the major phase present in these

covalent ceramics prepared by pyrolysis of polymeric precursors if pyrolysis temperatures are not sufficiently high to provide the activation energy required for crystallization. These amorphous covalent structures, which approach random bonding tosilicon, appearto be thecontinuous phase in the ceramic fibers. They can be described as a glassy silicon oxycarbide phase for the Si-C-0 system of fibers and a glassysilicon oxynitrocarbide phase for the Si-C-N-0 fibers. These types of structures are believed to be thermodynamically metastable. The thermodynamically stable structure is a mixture of crystalline species (e.g., rule-of-mixtures composition). Amor- phous compositionscontaining this much carbon or nitrogen have not been reported previously since they are difficult to prepare by methods tradition- ally used to prepare carbon and nitrogen-containing silica glasses.6970 Silicon oxycarbonitride glasses have not been previously reported. Such materials have been previously proposed and referred to as "ceramic alloys" by Baney.15

Based purely on statistics, polymer-derived ceramics with C, N, and 0 bonding to silicon approaching a random distribution would be expected to be amorphous at moderate pyrolysis temperatures unless bulk elemental composition is close to that of a known crystalline species in the Si-C-N-0 system (e.g., SiO,, Sic, Si3N,, or Si,N,O). This appears to be the case since the Si-C-0 ceramics, but not the Si-C-N-0 ceramics, contain p-Sic nanocrystallinity which increases in size and wt% at a fixed pyrolysis temperature as the composition more closely approaches Sic. Another crystalline phase, nanocrystalline pyrolytic carbon, appears to be present in all these compositions which contain excess carbon.

It is not unexpected that pyrolysis of an amorphous polymer in which mixed bonding to silicon is present (e.g., methyl groups on silicon atoms which are also bonded to 0 and/or NH groups) will lead to an amorphous ceramic which retains mixed bonding to silicon. Crystallization does not occur unless the composition is close to silicon carbide or nitride, at least for pyrolysis temperatures less than 1400C. Above 1 400C, N,, CO, and Si0 may evolve from the ceramic with crystallization (thermal decomposition). This is thecase also for Si-N-0 ceramics derived from polymers71 and


for other polymer-derived ceramics.72 The Si-N-0 ceramic fibers were prepared from polycarbosilanefibers which had been pyrolyzed in ammonia, a process which removes all carbon from the material and substitutes nitrogen on silicon.

THERMAL STABILITY

Introduction Thermal stability is a broad term used to describechangesof properties,

structure, morphology, and composition as functions of the temperature history and aging atmosphere. It is necessary to retain good mechanical properties of ceramic fibers at temperatures required to fabricate them into ceramic matrix composites, approximately 1200 to 1400C. This section is concerned first with changes in the tensile strength of Si-C-N-0, Si-C-0, and Si-C ceramic fibers, then with the parameters of structure and composition which control these properties. Emphasis is on irreversible changes (measure

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Fiber Reinforced Ceramic Composites