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Creep limitations of current polycrystalline ceramic fibers
Composites Science and Technology 51 (1994) 213-222
CREEP LIMITATIONS OF C U R R E N T P O L Y C R Y S T A L L I N E C E R A M I C FIBERS James A. DiCarlo NASA Lewis Research Center, Materials Division, 21000 Brookpark Road, Cleveland, Ohio 44135, USA (Received 31 July 1992; accepted 20 December 1992)
typically related to the matrix (e.g. oxidation, plasticity) and occur at temperatures lower than -800°C where current ceramic fibers first begin to show significant creep-related stiffness and strength degradation. 1 For ceramic-matrix composites, however, this loss in fiber structural properties can occur at the higher use temperatures projected for a variety of high-performance structural applications. Indeed, in some cases, current ceramic fibers show stiffness and strength losses at temperatures well below those where similar effects occur in potential ceramic matrices. At the present time, the structural implications of these effects have not been fully comprehended, either experimentally or theoretically. Clearly, critical questions arise as to whether for any application above -800°C, the creep and creeprelated effects displayed by current ceramic fibers can be tolerated and, if not, what minimum properties should be displayed. The objective of this paper is to present an overview of the issues, property status, and potential for the use of creep-prone polycrystalline ceramic fibers in high-temperature structural ceramic composites. This is accomplished first by establishing some fiber creep requirements based on obtaining and maintaining high mechanical and environmental performance of the composites. With this background, the creep properties for a variety of currently available SiC and A1203 polycrystalline fibers are then reviewed. Focus is placed on these two compositions because of the general technical need for long-term operation in oxidizing environments. Finally, by comparing the measured creep behavior against the fiber creep requirements, conclusions are drawn concerning optimum performance to be achieved in a polycrystalline SiC fiber and potential microstructural approaches that can be implemented in current fibers to achieve this performance.
Abstract The objective of this paper is to present an overview of the issues, property status, and potential for use of creep-prone polycrystalline ceramic fibers in thermostructural ceramic composites. Issues arise because the fine-grained microstructures of high-strength fibers can result in creep-related property changes, often at temperatures as low as 800°C. The underlying mechanism is grain boundary sliding controlled by grain size and grain boundary character, and thus by the fiber processing method. With the assumption of upper and lower limit creep requirements, the creep properties of a variety of current SiC and Al203 polycrystalline fibers are reviewed and discussed. Property evaluation is based on the results of a simple bend stress relaxation test which allows predictive creep equations to be developed for each fiber type describing the effects of time, temperature, and applied stress. It is shown that sintered SiC fibers with grain sizes below lO00nm appear to offer the best performance in terms of strength and creep resistance. However, even these fibers may not be capable of long-term service above 1400°C. Keywords: fibers, creep, stress relaxation, silicon carbide, composites INTRODUCTION For a thermostructural composite, it is desirable that the reinforcing fibers display high stiffness (modulus) and high strength relative to the matrix, and maintain these properties to as high a temperature as possible. Among other benefits, this allows the composite to operate in the elastic range to high stresses over a wide temperature range. For polymer- and metalmatrix composites reinforced by currently available ceramic fibers, this condition is generally achieved up to the maximum use temperatures of these composites. This is the case because use-limiting effects are
CREEP REQUIREMENTS To put into perspective the creep requirements of polycrystalline ceramic fibers in ceramic composites, it
Composites Science and Technology 0266-3538/94/$07.00 © 1994 Elsevier Science Limited. 213
James A. DiCarlo
214
Constant Stress, ~ro
III
w e.
/i
c., -
_
7.7-o Time
Constant Strain,
eo
or 0 ---- EF. 0
Time
Fig. 1. Typical curves for (a) creep and (b) stress relaxation of polycrystalline ceramics. is important first to consider the general phenomena of creep and stress relaxation in polycrystalline monolithic ceramics and then to examine some simple concepts concerning expected behavior of fibers and matrices during composite creep. As shown in Fig. l(a), if a constant tensile stress of Oo is applied to a monolithic ceramic at constant temperature, a tensile strain • will develop in the material which consists of two parts:
c=•~+c~
(1)
Here, •e is the instantaneous elastic or timeindependent strain which is related to applied stress by Hooke's law; that is, •c = ao/E, where E is the Young's modulus of the monolithic. At time t = 0, a time-dependent creep strain •e also develops and grows at a rate that increases with stress and temperature. Typically the polycrystaUine ceramic shows three creep stages, I, II, and III, which differ both macroscopically and microscopically in their behavior. In the first or primary Stage I, the creep rate decreases with time until it levels off at a constant value during Stage II. The underlying mechanism
giving rise to this deformation is thermally activated grain boundary sliding in which the grains slide infinitesimal distances relative to each other due to atomic diffusion.2'3 The diffusion can take place within second phases in the boundaries, or within the host material by grain boundary diffusion (Coble creep) or lattice diffusion (Nabarro-Herring creep). -~Due to the large aggregate of grains, the infinitesimal sliding distances add up to yield a time-dependent macroscopic dimensional change of the material. Because of reduced diffusion distances, smaller grain sizes increase creep rates. For example, the steady-state creep rate in Stage II has been found both theoretically and experimentally to vary inversely with d y, where d is grain size and the power y is generally between 2 and 3 depending on the underlying diffusion mechanism. 3 In Stage I, the decreasing strain rate is explained by the buildup of internal elastic stresses due to grain to grain contact. 2 This is supported by the observance of time-dependent creep recovery when the applied stress is removed, Indeed, the amount of recoverable creep is a good measure of the amount of creep strain that actually occurred during Stage I. This type of time-dependent and recoverable creep is generally characterized as anelastic deformation. The anelastic creep strain can be as low as 50% of the elastic strain, due for example to the Zener effect,a or in some cases greater than 10 times the elastic strain. 4 In Stage II, the buildup of internal elastic stresses can no longer be maintained in the presence of the concurrent internal diffusion phenomena. This leads to additional grain sliding at a steady-state rate controlled by the underlying diffusion mechanism. Upon removal of the applied stress, the Stage II contribution to macroscopic creep will not recover, leaving a permanent material deformation. This type of time-dependent and non-recoverable creep is generally characterized as viscoelastic deformation. Finally, in Stage III, internal void or cavity damage at grain boundary junctions become large enough not only to contribute to an increased creep rate but to internal flaw growth, resulting in rapid creep rupture of the ceramic. Another property of polycrystalline ceramics controlled by the same creep machanism is that of stress relaxation. As shown in Fig. l(b), if a constant strain Eo is applied at constant temperature, the internal stress within the material will decrease with time due to stress relief by grain boundary sliding. For convenience, one can introduce a stress relaxation ratio m which is defined as the internal stress at any time over the initial stress o0 = Eeo; that is:
re(t, T, e,,)= o/oo
(2)
Because the stress initially relaxes during the anelastic primary creep stage and the underlying mechanisms
Creep limitations of current polycrystalline ceramic fibers typically display a wide distribution in relaxation times, one can make the fairly accurate approximation2 that: m = %/e = (1 + ¢c/%) -1
(3)
where E is given by eqn (1); that is, the total creep strain that would have occurred in a creep test at the same temperature with an applied stress of or0. Thus, only a small creep strain ec equal to ee (= %) would yield a significant stress drop of m = 0.5. In addition, because creep strain is nearly proportional to stress for grain boundary sliding,2'3 it follows that the m ratio is effectively independent of eo but only dependent on time and temperature. Independence of applied strain in turn implies that m measured in a bend stress relaxation test would be independent of position in the bend specimen and also would be equivalent to the stress relaxation ratio m' measured in a pure tensile test. Finally, by eqn (3), it can be seen that the m ratio varies simply from a value of unity for purely elastic materials to near zero for creep-prone materials with creep strain ratios, E~/ee, greater than 10. As will be pointed out, all of these properties often make the m ratio a more convenient parameter than creep strain to characterize the creep behavior of polycrystalline fibers.
Creep mismatch problem With this background, we can now turn our attention to ceramic composites which consist of at least two polycrystalline monolithic ceramics, the fiber and matrix, and possibly more--for example, if polycrystalline coatings are employed on the fiber and matrix. Figure 2 displays a schematic of a desirable stress-strain curve for a structural fiber-reinforced ceramic composite. For high performance of this material, it is important that the matrix fracture strain eum be high so that applied strains during service can be kept at reasonable levels without causing cracks through the matrix thickness. These cracks are
~Ura 81rain
Fig. 2. Schematic of desirable stress-stress curve for fiber-reinforced structural ceramic composites.
215
generally undesirable because they reduce composite life by allowing aggressive environments (e.g. oxygen) into the fiber-matrix interface region and by contributing to local fiber-matrix interracial abrasion under fatigue conditions. It is also important that the fiber fracture strain euf be as high as possible to avoid catastrophic composite fracture when the matrix first fractures and to allow the composite some strain capability by multiple matrix cracking. By simple stress analysis of a unidirectional composite, it can be shown that this multiple cracking condition can be achieved when:
Euf/Eum>
1+ X
(4)
X = Era Vra/EfVf
(5)
Here,
where Era, Ef and Vm, Vf are the Young's moduli and volume fractions of the matrix and fiber, respectively. Since the ;C parameter is typically fixed by conditions prior to composite fabrication, it follows that it is important that neither %m nor the ratio 6 u f / f u ra degrade during composite processing or use. One way in which degradation of these properties can occur is by the development of residual internal strains e g on the matrix (cf. Fig. 2). A well known and often observable example of this situation is under conditions where the coefficients of thermal expansion, ~m and c~f, of the matrix and fiber differ. After processing, upon cooling the composite across a temperature range AT, simple analysis for a unidirectional composite predicts an axial residual matrix strain given by: Era S m ( ( ~ ' m a c f ) A T / ( 1 + X)
(6)
R is Here, AT is assumed positive, and the sign of Em positive for tensile strain and negative for compressive strain. If a sufficiently high residual tensile strain R> develops during composite cooldown so that 6m eum, the matrix will crack before reaching room temperature. On the other hand, if high compressive 6Rm develops during cooldown, and if the opposing tensile strain on the fibers does not fracture them, the composites may display undesirable brittle behavior upon matrix failure. That is, by eqn (4), the new fiber to matrix fracture strain ratio may no longer be greater than (1 + X). In a manner similar to the thermal expansion mismatch problem described above, a creep mismatch between polycrystalline ceramic fibers and matrices can also lead to undesirable residual strains on the matrix. For example, assuming a constant tensile or compressive applied stress on the composite and isostrain conditions for the constituents, the load on the more creep-prone constituent will decrease with time while the load on the less creep-prone constituent increases with time. This time-dependent
James A. DiCarlo
216
load shift will cause an increasing elastic deformation of the less creep-prone constituent which will augment any intrinsic creep deformation it experiences. The net result is an observable creep deformation of the entire composite which is primarily controlled by stress loss or relaxation in the more creep-prone constituent. Detrimental aspects of this behavior arise, for example, when under tensile loading the fibers shed their load to the matrix causing it to reach its fracture strain during composite service. If fracture can be avoided, the matrix can still experience an undesirable tensile ERm after stress removal. Similar behavior is expected for a composite compression situation in which a more creep-prone matrix sheds its R load to the fibers, and then develops a tensile e,,1 after stress removal. To illustrate this creep mismatch problem, let us assume a unidirectional composite, a creep-prone fiber, an elastic matrix, and g > 1, that is, E m V m > EfVf. Under this situation, if the composite is subjected to an elastic tensile or compressive strain E0 at t = 0 , the fiber effectively experiences a constant strain as it stress relaxes with relaxation ratio mr(t, T). The matrix then deforms elastically beyond c0 and the composite displays a time-dependent creep strain ec(comp). Upon load removal and temperature R decrease, the matrix develops a residual strain em equal to ec(comp). Simple stress analysis shows that R Em ~
(1 - m,)e,,/z.
(7)
where signs of the strains are positive and negative for tensile and compressive conditions, respectively. Thus matrix residual strain increases with applied composite strain and fiber creep (stress relaxation), but decreases with relative load carried by the matrix (i.e. Z). It is important to realize that unlike the thermal mismatch ERm which can be avoided by choosing trm = trf, it is almost impossible to avoid the creep mismatch Era. R This is the case because fiber and matrix creep behavior must necessarily be different due to differences in microstructures (grain size, grain boundary phases, etc.) and to likely differences in internal stress states. Indeed, even for a fiber and matrix of the same chemical composition, the fiber should have a smaller grain size (and thus be more creep-prone) if it is to possess a fracture strain greater than that of the matrix.
Fiber creep goals To establish the creep property goals for a polycrystalline ceramic fiber in a ceramic-matrix composite, one must understand the dimensional requirements and the full range of stress, time, and temperature conditions the composite will experience in a particular application. Clearly this is beyond the scope of this paper not only because there is an
innumerable list of possible applications for ceramic composites, but also because for any given application it is probably impossible to understand all the conditions the composite will experience externally or the fiber will experience internally. For this reason the approach taken here is to develop fiber creep property goals based on a generic application in which the composite stress and temperature are held constant on the composite throughout all high-temperature service. Regarding goals for composite dimensional stability, it appears that guidelines for many structural components require a creep strain limit of 1%. If it can be assumed that fiber creep strain E,, controls composite creep, a fiber creep limit of 1% might seem appropriate for our generic application. However, this assumption may be unrealistic since it implies that the matrix either does not carry any load or creeps at the same rate as the fiber. Since the first case is structurally undesirable because it means a completely cracked, debonded or compliant matrix and the second case, as previously discussed, is highly unlikely, it would seem that a 1% creep strain limit for the fiber is optimistic and represents an upper extreme. More appropriately, a better situation (other than a completely elastic fiber) would be one that kept the matrix structurally intact while allowing some creep deformation in the fiber. A lower extreme condition would be that of an effectively elastic matrix and a stress relaxing fiber. Then the creep limit on the fiber would be that which avoided the development of a R high residual tensile strain em on the matrix after R long-term composite creep. Assuming Em results in only a small change in cure (for example, e,.R = 0 " I Eum ) and that high composite applied service strains are desirable (that is. e0--~ Corn), it follows from eqn (7) that the total fiber stress relaxation allowed in a simple unidirectional composite can be estimated by m = 1 -0.1 Z For typical ceramic composites, VmlVf= 1 to 1.25 and Em/Ef = 0"2 to 2 so that g m , x - 5. It follows then that m = 0 . 5 might be considered a more realistic lower 'creep' goal for the reinforcing fiber. Thus, for this study, creep goals for current and future polycrystalline ceramic fibers will consist of two extremes: an optimistic upper extreme goal of ec = 1% in which the matrix effectively plays no structural role, and perhaps a more realistic lower extreme goal of m = 0.5 in which the matrix remains elastic without cracking during structural use. It is assumed that these goals must be met within the service life of our generic application in which the ceramic composite is held at constant temperature and constant applied stress.
217
Creep limitations o f current polycrystaUine ceramic fibers
PROPERTY STATUS Experimental To present a broad perspective of the creep property status and issues for current fibers, this study has selected a variety of polycrystalline ceramic fibers which not only have potential for ceramic composite reinforcement but also display a wide range of microstructures. The base chemical compositions are SiC and A I 2 0 3 because these compositions have demonstrated long-time stability in oxygen (at least in bulk form up to 1400°C), and because most anticipated applications of high technical payoff will require operation in agressive oxidative environments. Key details (and references) concerning the seven selected fiber types are given in Table 1. Prime among these are the fabrication processes which almost necessarily give rise to microstructural factors controlling creep by grain boundary sliding, that is, grain size, grain boundary morphology, and second phases that may exist in the grain boundaries. As indicated, some of these processes are only newly developed, but most have been existent for sufficient duration to result in the production at one time or another of commercially available fibers. Generally, the developmental processes were recently initiated to address strength retention and creep-related problems of current commercial fibers. To evaluate and compare the creep performance for each fiber type, the bend stress relaxation (BSR) test shown in Fig. 3 was developed.13 For this test, a short length of fiber is placed in a pure bending mode of radius Ro and then thermally treated for a specific time t at a constant temperature T, typically in an inert environment. During the test, the local applied strains within the fiber are held constant, increasing linearly from the fiber centroid plane to the fiber surface where they reach a maximum of %* = D / 2 R o , where D is fiber diameter. For strong, small diameter fibers, bending modes of different applied strains can
Initial loop at room temperature
Unrestrained, relaxed fiber
. - - ~ treat ~
ft', tl I m = ~lcr 0 = 1 - RO/FIa - Graphite block
L•••ite
Ra - - - - ~ ~
~T,t
RO= 4 o r a m m
Fig. 3. Bend stress relaxation test for creep evaluation of polycrystalline ceramic fibers.
be achieved by simply tying the fibers into small loops of different radii. For weak or large diameter fibers, graphite jigs can be used to provide the pure bending constraint (cf. Fig. 3). After treatment, the fiber loop is broken or the jig pieces separated, revealing a fiber with a creepinduced curve of constant radius Ra. For all fiber types tested in this study, the radius ratio Ro/Ra showed a strong dependence on T, a weak dependence on t, and effectively little or no dependence on E*. This strain independence was to be expected based on grain boundary sliding as the controlling creep mechanism. Assuming after treatment that the local creep strains within the fiber were directly proportional to the local applied strains and inversely proportional to R~, it follows from eqn (2) that the radius ratio could then be simply used to measure the bend stress relaxation ratio m; that is: (8)
m = 1 - Ro/Ra
Table 1. Polycrystalline fiber specimens Fabrication processes
Type
Source
Primary phase
Other phases
/3-SIC
fl-SiC fl-SiC re-SiC
C, Si (C substrate) C, 12 Wt%O C, 0-5 Wt%O B, C
Ai203
Chemical vapor deposition
SCS-6
Polymer spinpyrolysis Polymer spinsinter (+ powder) Slurry spin
Nicalon (Hi-NiC)" (DC)" (Carb)"
Textron specialty materials Nippon Carbon Nippon Carbon Dow Corning Carborundum
FP PRD-166
DuPont DuPont
aDevelopmental fibers.
fl-SiC
A1203
Approx. grain size (nm)
Ref. nos
100
5, 6
B4C
3 4 3O 2000
7 8 9 10
SiO2 ZrO2, SiO2
300 400
11 12
218
James A. DiCarlo
In comparison to a traditional tensile creep test, the BSR test offers many advantages from both a practical and basic point of view. T M Clearly on the practical side, one can simply and quickly ascertain the onset and relative creep of various fibers without the need for long fibers, mechanical loading devices, large controlled-environment chambers, or chemically compatible sensors. On the basic side, as will be demonstrated and as predicted by eqn (3), BSR m data can also be used to yield fairly accurate predictions of the creep strain ec and stress relaxation ratio m' that will occur in a polycrystalline fiber when tested in a pure tensile mode under the same test conditions. That is, by eqn (3):
(cc)l(ee)
= (1)/(m') - 1 = (1)/(m) - i
(9)
Here, ee = e0 is the applied tensile elastic strain and e~/ee is defined as the normalized tensile creep strain ratio. Results mid discussion BSR results for the six most characterized fiber types are shown in Fig. 4 as a function of reciprocal absolute temperature. Test conditions were a treatment time of l h and average applied strains of about 0.2%. Because of good reproducibility (Am = +0-02), only three to eight specimens were used for each data point. Examining the Fig. 4 results, one can see that even the simple 1-h BSR test can shed much light on the creep performance of current fibers. For example, for each fiber type, one can discern its creep onset temperature and its creep resistance relative to the other fiber types. This fiber comparison can be done Fiber
F-O~ %
0.3
0 0
N~.alon
v a
D(~ Comw~ O~lum
N
PRD-166 FP
ml!
Temperature, °C
1800 1600
4.5
5.0
.3 .3 .2 .1 .1
5.5
1400
(LO
1200 1100 1000
6.5--7.0
7.5
8.0
900
8.5 9.0X10-4
Reciprocal tempecature, K -1
Fig. 4. One-hour bend stress relaxation results for a variety of polycrystalline SiC and A1203 fibers.
in terms of stress relaxation at a certain temperature or in terms of temperature capability to achieve a certain degree of stress relaxation. Thus, for example, in regard to temperature capability, at all stress relaxation levels the sintered SiC fibers rank highest and the A1203 fibers rank lowest. The somewhat better creep resistance of the PRD fiber relative to the FP fiber can be explained by the ZrO2 additions which inhibit grain boundary motion. 12 For the SiC-based SCS-6 and Nicalon fiber types, creep onset begins as low as the A1203 fibers. For the SCS-6 fiber, free silicon remaining during the chemical vapor deposition process appears a likely source for enhanced diffusion at the grain boundaries. 6 For the polymerderived Nicalon, excess oxygen from the curing stage and a microcrystalline character from low-temperature pyrolysis are likely contributors to its poor creep resistance. 7's As for the sintered SiC fibers, it would appear that the absence of silicon and oxygen at the boundaries does much to improve their temperature capability. As will be discussed, the difference between the sintered or-SiC and r-SiC results are probably related more to grain size than to phase. Although the data of Fig. 4 may be useful for estimating short-time fiber performance, many applications will require service times of 1000h and beyond. To understand better the behavior for these and other applications, the BSR measurements were extended to 100-h treatments and to applied strain levels up to 1%. These results were then empirically modeled to calculate accurately performance within the envelope of test conditions and to predict performance beyond this envelope. When treatment times were increased an order of magnitude, for example, to 10 and 100 h, the characteristic S-shaped curves of Fig. 4 shifted a specific A(1/T) to the right (lower temperatures) depending on the fiber type. This behavior is characteristic of a thermally activated process controlled by an activation energy Q that is effectively constant throughout the main part of the S-curve. For one order of magnitude time change: Q = 2.303R/A(1/T)
(10)
where R is the gas constant (8.314 J/mole). Calculated Q values for the various fiber types are listed in the last column of Table 2. Relatively, these values increase with increasing creep resistance. Absolutely, the Q values should be related to the diffusion mechanisms controlling grain boundary sliding. Although measured primarily in Stage I, the Q values for the sintered or-SiC fiber and the AI203 fibers agree closely with the activation energies measured in Stage II for bulk 0r-SiC and bulk A1203, respectively. 15"~6 Regarding modeling of the BSR data, previous tensile creep studies on SCS-6 fibers6'17 showed that in the creep region of interest (ec> 1%), the performance of these fibers could be accurately described by
219
Creep limitations o f current polycrystalline ceramic fibers Table 2. Best-fit parameters for creep and stress relaxation
Fiber
SCS-6 (BSR) SCS-6 (tensile) Nicalon Dow Coming Carborundum PRD-166 FP
Approximate fiber modulus E (GPa)
Aoa
400 400 200 400 400 370 370
2.8 x 105 1.0 x 10"~ 2.2 x 10s 69 x 105 0-83 x 105 7 x 10l° 49 x 10"'
n
B
Correlation
(K)
factoff
pb
Controlling creep energy,a
Q
(kJ/mol) 1-1 1.1e 1.2 1-5 1.2 1.35 1-5
0.37 0-37 0-40 0.45 0.31 0-53 0-58
25 800 + 950 25 80(Y 24 200 + 1 400 35 133 + 2 450 31 600 -t- 1 800 41 000 + 2 120 41 150:1:2 460
0.991
580 + 38
0-980 0-962 0.971 0-964 0.959
500 + 65 670 + 90 830 + 145 640 + 70 590 + 76
"Applies for stress o in gigapascals, time t in hours, and temperature T in Kelvin. bCalculated from p = RB/Q. CFor linear regression to determine B parameter. dCalculated from A(1/T). eAssumed from bend results. Average Stress cra : a0E/2 = 0.5 GPa
the following semi-empirical creep equation: ~
ec = Aoo~[t e x p ( - Q / R T ) ] p = Aoo~t p e x p ( - B / T )
-
1"~
(11)
Here, B = p Q / R and Ao, n and p are empirically determined creep parameters. In a more recent study, TM it was assumed that this equation could be applied to any polycrystalline ceramic fiber in the stress relaxation range m > 0-1 or in the normalized creep strain range co/co < 10. Equation (9) was then employed to convert the measured m data to normalized tensile creep strains which were then used with eqn (11) to calculate best-fit empirical parameters for the SCS-6, Nicalon, FP, and P R D fibers. The results are given in Table 2 together with new additional parameters for the sintered fibers. For use of these parameters with eqn (11), strain is in absolute units, stress in GPa, time in hours, and T in Kelvin. In that same study, ~4 the accuracy of the m conversion (eqn (9)) and the empirical model (eqn (11)) were demonstrated by plotting available tensile creep data and converted m data as a function of absolute reciprocal temperature. Figure 5 represents such a plot for the Nicalon fiber. It can be seen that: (1) at constant test times straight lines can be fit to all data, implying a constant B parameter; (2) the A ( 1 / T ) shift with each decade of time is constant, implying constant Q and p = Q / B R parameters; and (3) close agreement exists between the converted BSR m data and the tensile creep data of Jia et al.lS The same conclusions could be made for the P R D and FP fibers. 14 For the SCS-6 fiber, although the n, p and Q values agreed, the absolute values of E c / e e for the converted bend results were consistently higher than those for the actual tensile data. 17 This is the basis for the different A0 values listed in Table 2 for the SCS-6
" 1
--
o
0 I-I
•
10 100 1}Ten
z
.1
I 5.0
5.5
I
I
I
I
I
6.0 6.5 7.0 7.5 8.0 Reciprocal temperature, K -1
I 8.5
I 9.0x10"4
Fig. 5. Temperature-time dependence of normalized tensile creep strain for Nicalon fibers. Open points predicted from bend stress relaxation data; closed points are actual tensile data of Jia et al. ,8 fiber. This can be explained by the fact that, unlike the other fibers, the SCS-6 fiber is not homogeneous. That is, it contains not only a carbon substrate core but an outer SiC sheath which creeps or stress relaxes more than the inner sheath. Thus stress relaxation in bending should be greater than that in tension. In summary, the current status of six fiber types can be described by eqns (11) and (9) and the appropriate creep parameters listed in Table 2. These descriptions should be fairly accurate for the conditions: 0 - 9 > m > 0 . 1 or 0 - 1 < e c / e e < 10. Since these conditions fall within or near the upper and lower creep goals previously discussed, this methodology will be used in the next section to provide estimates of the potential for each fiber type in our generic composite application. PERFORMANCE
POTENTIAL
Given their current microstructures, the creep performance of the polycrystalline fibers can now be
220
James A. DiCarlo Assumed conditions: ¢o = 100 MPa; t = 1000 hr L FP r / / / / / . / , / / / / J
PRD-166 ~ / / / / ' / / / / / / / 1 Nicalon
I
~ / ' ~ ~ "
scs-e
Creep limits
I
J
1771 m _>0.5 I I =-.¢ _< 1%
I
//////////////A
C=~oru.dum
////////////////A I
>2GPaS
I
Dow C o ~ n g
I
Strength ~1_._j C A R B ~ ' ~ f 0015OO --
////////////J
0
Assumed conditions: ~o = 400 MPa t = l hour 1600 -- m=0.5forstressrelaxatJon, or~c/~o=l forcreep
I
~.1400 _
I I I
I
I
I
I
/
I
200 400 600 800 10001200140016001800 Temperature, °C
_
Fig. 6. Estimated fiber upper use temperatures under the assumed service conditions and creep limits. compared against the two creep limit goals of ec = 1% and m =0.5. Since creep depends on the time, temperature, and applied stress conditions, one must choose two of the three conditions in order to determine the third. For example, one might assume a particular operating temperature and stress in order to determine maximum service life, or assume a certain service life and temperature and calculate the maximum operating creep stress. However, since ceramic composites are being examined for their potential to operate at high temperature, the approach taken here will be to assume a common service life of 1000 h and a realistic applied stress of 100 MPa, and then estimate maximum or upper fiber use temperatures using the Table 2 parameters. The results are shown in Fig. 6 with the hatched bars for the lower limit of m =0.5 and the open bars for ec= 1%. Examining Fig. 6, one finds for both limits a similar ranking as indicated by the simple 1-h BSR data of Fig. 4. That is, in terms of temperature capability, the A1203 fibers rank lowest and the sintered SiC fibers highest. Also it can be seen that some temperature gain can be achieved in the unlikely situation in which a composite upper creep limit of 1% yields a fiber limit of the same value. It should be realized that these temperature gains quickly diminish when the applied stress is taken above the selected value of 100 MPa. This, however, will not be the case for the lower limit temperatures since rn is effectively independent of stress. Thus, the m = 0.5 temperatures probably represent the most conservative estimates for the six fiber types. Assuming this, one concludes that although the sintered o~-SiC fiber currently represents the best fiber for most structural ceramic composite applications, its potential appears to be limited to temperatures less than 1400°C. Given the above observations, the question arises whether improvements can be made in the other SiC fiber fabrication processes in order to achieve creep resistances comparable or better than that of the a~-SiC fiber. To shed light on this question, its of
S / ~
1100
"-- Sintered
DC ..,,~//////~ "
~-s~c
; th ko :
" / / t ( dr°.'-
• Hi-NiC ONiC
I 10
0 SCS I 102 Grain size d, nm
103
........I 104
Fig. 7. Grain size effects on SiC fiber upper use temperatures under the assumed conditions. Hatched area represents estimate of grain-size conditions for optimum creep resistance (solid line) and for strength greater than 2 GPa (dashed line), interest to examine Fig. 7 which plots upper use temperatures versus grain size for the SiC fibers. The assumed conditions of 1 h and m = 0.5 allow the 1-h BSR measurements to be used for the four SiC fibers plus the Hi-Nicalon fiber for which limited BSR data currently exist. In Fig. 7 one can see a general trend in which increased grain size results in improved creep resistance. As previously discussed, this is to be expected because increased grain size should increase diffusion distances, thereby slowing down grain boundary sliding and macroscopic creep. The solid line is drawn assuming that the tr-SiC fiber data point represents the ultimate in creep behavior and that grain size d can be introduced into the semi-empirical eqn (11) by the relation: A,, = A
~ / d 0"4 ~
A1/d p
(12)
The choice of an inverse d p dependence with p = 0-4 is based on the average p values observed for all the SiC fibers and the fact that grain boundary sliding in the early creep stages ( m > 0 . 4 ) is predicted theoretically2 and observed experimentally~9 to give a creep strain which is a function of the parameter q), where: (a = ( t / d ) e x p ( Q / R T )
(13)
It is interesting to note in Fig. 7 that the sintered fl-SiC fiber data points fall very near the solid line and the data points for all other SiC fibers fall below the line. This suggests that for the assumed conditions the solid line may indeed be the optimum performance to be expected from SiC polycrystalline fibers with equiaxed grains.
Creep limitations of current polycrystalline ceramic fibers
Improvement approaches Assuming then that the solid line of Fig. 7 represents near optimum performance for polycrystalline fibers, what direction should each fiber fabrication process take in order to produce a more creep resistant product? Clearly, the sintered fl-SiC fiber (DC) can move up line by increasing grain size. However, there is a limit because increasing grain size will eventually reduce fiber strength when the largest grain or flaws of this grain size become strength controlling. The vertical dashed line indicates the approximate grain size, 1000 nm, when SiC fiber strength is about 2 GPa, a value often quoted as a minimum fiber strength requirement for structural ceramic composites. This estimate is simply made using Grittith's Law and a SiC fracture toughness of 3 M P a m ~2. Since creep is controlled by the average grain size and strength by the largest, the possibility exists that the sintered fl-SiC has already reached its creep limit, unless perhaps its grain-size distribution can be narrowed. Thus, the hatched area in Fig. 7 represents a general grain-size envelope that fiber manufacturers should seek for optimum creep and strength performance. Even for the best performing sintered tr-SiC fiber, current grain size appears too large to meet fiber strength requirements. Although creep performance will suffer somewhat, the fabrication processes for this fiber should move in a direction which reduce average d to below 1000 nm while still maintaining a narrow distribution. For the chemical vapor deposition and polymerprecursor fibers, improvements appear to lie first in removing low-viscosity grain boundary phases, such as, free silicon and Si-O complexes, and perhaps amorphous carbon. Both directions appear possible by changes in processing conditions. Indeed, by using radiation curing, the Hi-Nicalon fiber has significantly reduced oxygen content with respect to the Nicalon fiber, s The payoff in better creep performance can be clearly seen in Fig. 7. If all oxygen can be removed, behavior similar to the polymer-derived and sintered fl-SiC fiber should be expected. Figure 7 suggests that as oxygen is removed, grain size increases in a beneficial manner. What processing improvements are needed to further increase grain size have yet to be determined. Two final points concerning process improvements need to be mentioned. First, to evaluate whether any process change has led to improved creep behavior, almost immediate information can be obtained by using the simple BSR test and the Fig. 7 solid line. That is, a few 1-h tests at different temperatures and at an approximate surface strain of Eo = 0.2% could determine whether any relative improvements in the m ratio have been made. The results can also yield an estimate of the upper use temperature for m = 0.5 which with grain-size information could then be
221
compared with the Fig. 7 line to determine whether optimum performance has been achieved. A second important point is that optimum creep performance in this study was derived empirically using primarily SiC fibers with equiaxed grains. Thus, as new fibers are developed, the concept of optimum performance may have to be redefined. Based on the grain boundary sliding mechanism, it may be possible to introduce different boundary conditions or second phases which inhibit sliding more than those currently present in the sintered SiC fibers. Likewise, grain shapes or aspect ratios can be changed with the long axis parallel to the fiber axis. By increasing diffusion distances, this approach has been successfully employed to improve the creep resistance of tungsten filaments while still maintaining strength properties. Indeed this method is currently being attempted to improve the creep resistance of polycrystalline AI203 fibers. 2° If the grain aspect ratio is taken to the extreme, one would then have single crystal fibers which should provide the optimum in creep performance, perhaps increasing upper use temperature capability by as much as 200°C. ~ However, although a commercial processing approach has been developed for continuous single crystal AI203 fibers, 2~ no method currently exists for continuous single crystal SiC fibers.
CONCLUDING REMARKS This study has attempted to demonstrate the current situation for using creep-prone polycrystalline ceramic fibers in structural ceramic composites. In so doing, some important issues have been exposed which will require further in-depth studies. One centers on the need to develop accurate fiber creep requirements from the performance requirements of the composite component. From a macroscopic point of view, the service conditions that the component experiences can be a complex variation of time- and spaciallydependent temperatures and stresses. How these influence fibers on the microscopic level adds a further degree of complexity, especially in light of the creep mismatch problem and the fact that fiber architectures can be very complex. An added problem is the lack of numerical or analytical codes to model and analyze time-dependent constituent behavior. In light of these complexities, it would appear that the best approach at the present time is to identify, as was done in this study, the fiber composition, fabrication process, and microstructure which gives the best creep performance and then to actually evaluate these fibers within the composite under the intended use conditions. The performance equations derived here may be of some use in setting approximate limits on the service conditions, but these certainly cannot replace
222
James A. DiCarlo
observation of performance under actual operating conditions. Another issue requiring further study is that there appears to be an optimum creep performance for equiaxed SiC fibers which, under the conditions of practical interest, is weakly dependent on grain size. Currently, all the tested SiC fibers fall at or below this optimum, suggesting possible directions for processing/microstructure improvements. Thus, opportunities apparently exist for obtaining more creep-resistant SiC fibers from a variety of fabrication processes. However, even if these performance gains can be achieved, upper use temperatures may be only about 1300°C for a 1000-h service life. This temperature will increase or decrease by about 60°C for an order of magnitude decrease or increase in service time, respectively. Thus, questions arise whether processing improvements in polycrystalline SiC fibers are indeed possible and whether these improvements are enough to significantly impact the use of these fibers in high-temperature structural ceramic composites. A negative answer to this latter question would then seem to imply the need to develop SiC fibers with directionally aligned grains or perhaps with single crystal microstructures.
ACKNOWLEDGMENTS The author wishes to acknowledge the invaluable professional and technical support of G. N. Morscher and T. C. Wagner.
REFERENCES 1. DiCarlo, J. A., High Temperature Structural Fibers-Status and Needs. N A S A TM, 105174, 1991. 2. Nowick, A. S. & Berry, B. S. Anelastic Relaxation in Crystalline Solids. Academic Press, New York, 1972. 3. Raj, R. & Ashby, M. F., On grain boundary sliding and diffusional creep. Met. Trans., 2 (1971) 1113-27. 4. Todd, R. I. & Hazzledine, P. M., Large anelastic strains at constant volume in superplastic tin-lead eutectic alloy. Scripta Met., 27 (1992) 127-32.
5. Ning, X. J. & Pirouz, P., The microstructure of SCS-6 SiC filter. J. Mater. Res., 6 (1991) 2234-48. 6. DiCarlo, J. A., Creep of chemically vapour deposited SiC fibers. J. Mater. Sci., 2 (1986) 217-24. 7. Yajima, S., Okamura, K., Hayashi, J. & Omori, M., Synthesis of continuous SiC fibers with high tensile strength. J. Am. Ceram. Soc., 59(7-8) (1976) 324-7. 8. Takeda, M. et al., Thermal stability of the low oxygen silicon carbide fibers derived From polycarbosilane. Ceram. Eng. Sci. Proc. 13 (1992) 209-17, 9. Lipowitz, J., Rabe, J. A. & Zank, G. A., Polycrystalline SiC fibers from organosilicon polymers. Ceramic Eng. Sci. Proc., 12 (1991) 1819-31. 10. Frechette, F., Dover, B., Venkateswaran, V. & Kim, J., High temperature continuous sintered SiC fiber for composite applications. Ceram. Eng. Sci. Proc., 12 (1991) 992-1006. i1. Stacey, M. H., Developments in continuous aluminabased fibres. J. British Cermic Trans., 87 (1988) 168-72. 12. Romine, J. C. New high-temperature ceramic fiber. Ceramic Eng. Sci. Proc., 8(7-8) (1987) 755-65. 13. Morscher, G. N. & DiCario, J. A., A simple test for thermomechanical evaluation of ceramic fibers. J. Am. Ceram. Soc., 75(1) (1992) 136-40. 14. DiCarlo, J. A. & Morscher, G. N., Creep and stress relaxation modeling of polycrystalline ceramic fibers. ASME, A D 22, A M D 122 (1991) 15-22. 15. Sargent, P. M. & Ashby, M. F., Deformationmechanism maps for silicon carbide. Scripta Met., 17 (1983) 951-7. 16. Cannon, R. W. & Langdon, T. G., Review: Creep of ceramics. Part 1. Mechanical characteristics. J. Mater. Sci., 18 (1983) 1-50. 17. Morscher, G. N., DiCarlo, J. A. & Wagner, T., Fiber creep evaluation by stress relaxation measurements. Ceram. Eng. Sci. Proc., 12(7-8) (1991) 1032-8. 18. Jia, N., Bodet, R. & Tressler, R. E., Creep and creep rupture behavior of Si-C-O and S i - N - C - O based continuous fibers. Paper presented at 93rd Annual Meeting, American Ceramic Society, 28 April-2 May 1991, Cincinnati, OH. 19. Ke, T. S., Experimental evidence of the viscous behavior of grain boundaries in metals. Phys. Reo. 17 (1947) 533-46. 20. Tressler, R. E. Penn State University, personal communication (1990). 21. Hurley, G. F. & Labelle, H. E., Elevated temperature strengthening of melt-grown sapphire by alloying. In Applied Polymer, Symposium No. 29, New and Specialty Fibers, ed. J. Economy. John Wiley, New York, 1976, pp. 131-49.
CREEP LIMITATIONS OF C U R R E N T P O L Y C R Y S T A L L I N E C E R A M I C FIBERS James A. DiCarlo NASA Lewis Research Center, Materials Division, 21000 Brookpark Road, Cleveland, Ohio 44135, USA (Received 31 July 1992; accepted 20 December 1992)
typically related to the matrix (e.g. oxidation, plasticity) and occur at temperatures lower than -800°C where current ceramic fibers first begin to show significant creep-related stiffness and strength degradation. 1 For ceramic-matrix composites, however, this loss in fiber structural properties can occur at the higher use temperatures projected for a variety of high-performance structural applications. Indeed, in some cases, current ceramic fibers show stiffness and strength losses at temperatures well below those where similar effects occur in potential ceramic matrices. At the present time, the structural implications of these effects have not been fully comprehended, either experimentally or theoretically. Clearly, critical questions arise as to whether for any application above -800°C, the creep and creeprelated effects displayed by current ceramic fibers can be tolerated and, if not, what minimum properties should be displayed. The objective of this paper is to present an overview of the issues, property status, and potential for the use of creep-prone polycrystalline ceramic fibers in high-temperature structural ceramic composites. This is accomplished first by establishing some fiber creep requirements based on obtaining and maintaining high mechanical and environmental performance of the composites. With this background, the creep properties for a variety of currently available SiC and A1203 polycrystalline fibers are then reviewed. Focus is placed on these two compositions because of the general technical need for long-term operation in oxidizing environments. Finally, by comparing the measured creep behavior against the fiber creep requirements, conclusions are drawn concerning optimum performance to be achieved in a polycrystalline SiC fiber and potential microstructural approaches that can be implemented in current fibers to achieve this performance.
Abstract The objective of this paper is to present an overview of the issues, property status, and potential for use of creep-prone polycrystalline ceramic fibers in thermostructural ceramic composites. Issues arise because the fine-grained microstructures of high-strength fibers can result in creep-related property changes, often at temperatures as low as 800°C. The underlying mechanism is grain boundary sliding controlled by grain size and grain boundary character, and thus by the fiber processing method. With the assumption of upper and lower limit creep requirements, the creep properties of a variety of current SiC and Al203 polycrystalline fibers are reviewed and discussed. Property evaluation is based on the results of a simple bend stress relaxation test which allows predictive creep equations to be developed for each fiber type describing the effects of time, temperature, and applied stress. It is shown that sintered SiC fibers with grain sizes below lO00nm appear to offer the best performance in terms of strength and creep resistance. However, even these fibers may not be capable of long-term service above 1400°C. Keywords: fibers, creep, stress relaxation, silicon carbide, composites INTRODUCTION For a thermostructural composite, it is desirable that the reinforcing fibers display high stiffness (modulus) and high strength relative to the matrix, and maintain these properties to as high a temperature as possible. Among other benefits, this allows the composite to operate in the elastic range to high stresses over a wide temperature range. For polymer- and metalmatrix composites reinforced by currently available ceramic fibers, this condition is generally achieved up to the maximum use temperatures of these composites. This is the case because use-limiting effects are
CREEP REQUIREMENTS To put into perspective the creep requirements of polycrystalline ceramic fibers in ceramic composites, it
Composites Science and Technology 0266-3538/94/$07.00 © 1994 Elsevier Science Limited. 213
James A. DiCarlo
214
Constant Stress, ~ro
III
w e.
/i
c., -
_
7.7-o Time
Constant Strain,
eo
or 0 ---- EF. 0
Time
Fig. 1. Typical curves for (a) creep and (b) stress relaxation of polycrystalline ceramics. is important first to consider the general phenomena of creep and stress relaxation in polycrystalline monolithic ceramics and then to examine some simple concepts concerning expected behavior of fibers and matrices during composite creep. As shown in Fig. l(a), if a constant tensile stress of Oo is applied to a monolithic ceramic at constant temperature, a tensile strain • will develop in the material which consists of two parts:
c=•~+c~
(1)
Here, •e is the instantaneous elastic or timeindependent strain which is related to applied stress by Hooke's law; that is, •c = ao/E, where E is the Young's modulus of the monolithic. At time t = 0, a time-dependent creep strain •e also develops and grows at a rate that increases with stress and temperature. Typically the polycrystaUine ceramic shows three creep stages, I, II, and III, which differ both macroscopically and microscopically in their behavior. In the first or primary Stage I, the creep rate decreases with time until it levels off at a constant value during Stage II. The underlying mechanism
giving rise to this deformation is thermally activated grain boundary sliding in which the grains slide infinitesimal distances relative to each other due to atomic diffusion.2'3 The diffusion can take place within second phases in the boundaries, or within the host material by grain boundary diffusion (Coble creep) or lattice diffusion (Nabarro-Herring creep). -~Due to the large aggregate of grains, the infinitesimal sliding distances add up to yield a time-dependent macroscopic dimensional change of the material. Because of reduced diffusion distances, smaller grain sizes increase creep rates. For example, the steady-state creep rate in Stage II has been found both theoretically and experimentally to vary inversely with d y, where d is grain size and the power y is generally between 2 and 3 depending on the underlying diffusion mechanism. 3 In Stage I, the decreasing strain rate is explained by the buildup of internal elastic stresses due to grain to grain contact. 2 This is supported by the observance of time-dependent creep recovery when the applied stress is removed, Indeed, the amount of recoverable creep is a good measure of the amount of creep strain that actually occurred during Stage I. This type of time-dependent and recoverable creep is generally characterized as anelastic deformation. The anelastic creep strain can be as low as 50% of the elastic strain, due for example to the Zener effect,a or in some cases greater than 10 times the elastic strain. 4 In Stage II, the buildup of internal elastic stresses can no longer be maintained in the presence of the concurrent internal diffusion phenomena. This leads to additional grain sliding at a steady-state rate controlled by the underlying diffusion mechanism. Upon removal of the applied stress, the Stage II contribution to macroscopic creep will not recover, leaving a permanent material deformation. This type of time-dependent and non-recoverable creep is generally characterized as viscoelastic deformation. Finally, in Stage III, internal void or cavity damage at grain boundary junctions become large enough not only to contribute to an increased creep rate but to internal flaw growth, resulting in rapid creep rupture of the ceramic. Another property of polycrystalline ceramics controlled by the same creep machanism is that of stress relaxation. As shown in Fig. l(b), if a constant strain Eo is applied at constant temperature, the internal stress within the material will decrease with time due to stress relief by grain boundary sliding. For convenience, one can introduce a stress relaxation ratio m which is defined as the internal stress at any time over the initial stress o0 = Eeo; that is:
re(t, T, e,,)= o/oo
(2)
Because the stress initially relaxes during the anelastic primary creep stage and the underlying mechanisms
Creep limitations of current polycrystalline ceramic fibers typically display a wide distribution in relaxation times, one can make the fairly accurate approximation2 that: m = %/e = (1 + ¢c/%) -1
(3)
where E is given by eqn (1); that is, the total creep strain that would have occurred in a creep test at the same temperature with an applied stress of or0. Thus, only a small creep strain ec equal to ee (= %) would yield a significant stress drop of m = 0.5. In addition, because creep strain is nearly proportional to stress for grain boundary sliding,2'3 it follows that the m ratio is effectively independent of eo but only dependent on time and temperature. Independence of applied strain in turn implies that m measured in a bend stress relaxation test would be independent of position in the bend specimen and also would be equivalent to the stress relaxation ratio m' measured in a pure tensile test. Finally, by eqn (3), it can be seen that the m ratio varies simply from a value of unity for purely elastic materials to near zero for creep-prone materials with creep strain ratios, E~/ee, greater than 10. As will be pointed out, all of these properties often make the m ratio a more convenient parameter than creep strain to characterize the creep behavior of polycrystalline fibers.
Creep mismatch problem With this background, we can now turn our attention to ceramic composites which consist of at least two polycrystalline monolithic ceramics, the fiber and matrix, and possibly more--for example, if polycrystalline coatings are employed on the fiber and matrix. Figure 2 displays a schematic of a desirable stress-strain curve for a structural fiber-reinforced ceramic composite. For high performance of this material, it is important that the matrix fracture strain eum be high so that applied strains during service can be kept at reasonable levels without causing cracks through the matrix thickness. These cracks are
~Ura 81rain
Fig. 2. Schematic of desirable stress-stress curve for fiber-reinforced structural ceramic composites.
215
generally undesirable because they reduce composite life by allowing aggressive environments (e.g. oxygen) into the fiber-matrix interface region and by contributing to local fiber-matrix interracial abrasion under fatigue conditions. It is also important that the fiber fracture strain euf be as high as possible to avoid catastrophic composite fracture when the matrix first fractures and to allow the composite some strain capability by multiple matrix cracking. By simple stress analysis of a unidirectional composite, it can be shown that this multiple cracking condition can be achieved when:
Euf/Eum>
1+ X
(4)
X = Era Vra/EfVf
(5)
Here,
where Era, Ef and Vm, Vf are the Young's moduli and volume fractions of the matrix and fiber, respectively. Since the ;C parameter is typically fixed by conditions prior to composite fabrication, it follows that it is important that neither %m nor the ratio 6 u f / f u ra degrade during composite processing or use. One way in which degradation of these properties can occur is by the development of residual internal strains e g on the matrix (cf. Fig. 2). A well known and often observable example of this situation is under conditions where the coefficients of thermal expansion, ~m and c~f, of the matrix and fiber differ. After processing, upon cooling the composite across a temperature range AT, simple analysis for a unidirectional composite predicts an axial residual matrix strain given by: Era S m ( ( ~ ' m a c f ) A T / ( 1 + X)
(6)
R is Here, AT is assumed positive, and the sign of Em positive for tensile strain and negative for compressive strain. If a sufficiently high residual tensile strain R> develops during composite cooldown so that 6m eum, the matrix will crack before reaching room temperature. On the other hand, if high compressive 6Rm develops during cooldown, and if the opposing tensile strain on the fibers does not fracture them, the composites may display undesirable brittle behavior upon matrix failure. That is, by eqn (4), the new fiber to matrix fracture strain ratio may no longer be greater than (1 + X). In a manner similar to the thermal expansion mismatch problem described above, a creep mismatch between polycrystalline ceramic fibers and matrices can also lead to undesirable residual strains on the matrix. For example, assuming a constant tensile or compressive applied stress on the composite and isostrain conditions for the constituents, the load on the more creep-prone constituent will decrease with time while the load on the less creep-prone constituent increases with time. This time-dependent
James A. DiCarlo
216
load shift will cause an increasing elastic deformation of the less creep-prone constituent which will augment any intrinsic creep deformation it experiences. The net result is an observable creep deformation of the entire composite which is primarily controlled by stress loss or relaxation in the more creep-prone constituent. Detrimental aspects of this behavior arise, for example, when under tensile loading the fibers shed their load to the matrix causing it to reach its fracture strain during composite service. If fracture can be avoided, the matrix can still experience an undesirable tensile ERm after stress removal. Similar behavior is expected for a composite compression situation in which a more creep-prone matrix sheds its R load to the fibers, and then develops a tensile e,,1 after stress removal. To illustrate this creep mismatch problem, let us assume a unidirectional composite, a creep-prone fiber, an elastic matrix, and g > 1, that is, E m V m > EfVf. Under this situation, if the composite is subjected to an elastic tensile or compressive strain E0 at t = 0 , the fiber effectively experiences a constant strain as it stress relaxes with relaxation ratio mr(t, T). The matrix then deforms elastically beyond c0 and the composite displays a time-dependent creep strain ec(comp). Upon load removal and temperature R decrease, the matrix develops a residual strain em equal to ec(comp). Simple stress analysis shows that R Em ~
(1 - m,)e,,/z.
(7)
where signs of the strains are positive and negative for tensile and compressive conditions, respectively. Thus matrix residual strain increases with applied composite strain and fiber creep (stress relaxation), but decreases with relative load carried by the matrix (i.e. Z). It is important to realize that unlike the thermal mismatch ERm which can be avoided by choosing trm = trf, it is almost impossible to avoid the creep mismatch Era. R This is the case because fiber and matrix creep behavior must necessarily be different due to differences in microstructures (grain size, grain boundary phases, etc.) and to likely differences in internal stress states. Indeed, even for a fiber and matrix of the same chemical composition, the fiber should have a smaller grain size (and thus be more creep-prone) if it is to possess a fracture strain greater than that of the matrix.
Fiber creep goals To establish the creep property goals for a polycrystalline ceramic fiber in a ceramic-matrix composite, one must understand the dimensional requirements and the full range of stress, time, and temperature conditions the composite will experience in a particular application. Clearly this is beyond the scope of this paper not only because there is an
innumerable list of possible applications for ceramic composites, but also because for any given application it is probably impossible to understand all the conditions the composite will experience externally or the fiber will experience internally. For this reason the approach taken here is to develop fiber creep property goals based on a generic application in which the composite stress and temperature are held constant on the composite throughout all high-temperature service. Regarding goals for composite dimensional stability, it appears that guidelines for many structural components require a creep strain limit of 1%. If it can be assumed that fiber creep strain E,, controls composite creep, a fiber creep limit of 1% might seem appropriate for our generic application. However, this assumption may be unrealistic since it implies that the matrix either does not carry any load or creeps at the same rate as the fiber. Since the first case is structurally undesirable because it means a completely cracked, debonded or compliant matrix and the second case, as previously discussed, is highly unlikely, it would seem that a 1% creep strain limit for the fiber is optimistic and represents an upper extreme. More appropriately, a better situation (other than a completely elastic fiber) would be one that kept the matrix structurally intact while allowing some creep deformation in the fiber. A lower extreme condition would be that of an effectively elastic matrix and a stress relaxing fiber. Then the creep limit on the fiber would be that which avoided the development of a R high residual tensile strain em on the matrix after R long-term composite creep. Assuming Em results in only a small change in cure (for example, e,.R = 0 " I Eum ) and that high composite applied service strains are desirable (that is. e0--~ Corn), it follows from eqn (7) that the total fiber stress relaxation allowed in a simple unidirectional composite can be estimated by m = 1 -0.1 Z For typical ceramic composites, VmlVf= 1 to 1.25 and Em/Ef = 0"2 to 2 so that g m , x - 5. It follows then that m = 0 . 5 might be considered a more realistic lower 'creep' goal for the reinforcing fiber. Thus, for this study, creep goals for current and future polycrystalline ceramic fibers will consist of two extremes: an optimistic upper extreme goal of ec = 1% in which the matrix effectively plays no structural role, and perhaps a more realistic lower extreme goal of m = 0.5 in which the matrix remains elastic without cracking during structural use. It is assumed that these goals must be met within the service life of our generic application in which the ceramic composite is held at constant temperature and constant applied stress.
217
Creep limitations o f current polycrystaUine ceramic fibers
PROPERTY STATUS Experimental To present a broad perspective of the creep property status and issues for current fibers, this study has selected a variety of polycrystalline ceramic fibers which not only have potential for ceramic composite reinforcement but also display a wide range of microstructures. The base chemical compositions are SiC and A I 2 0 3 because these compositions have demonstrated long-time stability in oxygen (at least in bulk form up to 1400°C), and because most anticipated applications of high technical payoff will require operation in agressive oxidative environments. Key details (and references) concerning the seven selected fiber types are given in Table 1. Prime among these are the fabrication processes which almost necessarily give rise to microstructural factors controlling creep by grain boundary sliding, that is, grain size, grain boundary morphology, and second phases that may exist in the grain boundaries. As indicated, some of these processes are only newly developed, but most have been existent for sufficient duration to result in the production at one time or another of commercially available fibers. Generally, the developmental processes were recently initiated to address strength retention and creep-related problems of current commercial fibers. To evaluate and compare the creep performance for each fiber type, the bend stress relaxation (BSR) test shown in Fig. 3 was developed.13 For this test, a short length of fiber is placed in a pure bending mode of radius Ro and then thermally treated for a specific time t at a constant temperature T, typically in an inert environment. During the test, the local applied strains within the fiber are held constant, increasing linearly from the fiber centroid plane to the fiber surface where they reach a maximum of %* = D / 2 R o , where D is fiber diameter. For strong, small diameter fibers, bending modes of different applied strains can
Initial loop at room temperature
Unrestrained, relaxed fiber
. - - ~ treat ~
ft', tl I m = ~lcr 0 = 1 - RO/FIa - Graphite block
L•••ite
Ra - - - - ~ ~
~T,t
RO= 4 o r a m m
Fig. 3. Bend stress relaxation test for creep evaluation of polycrystalline ceramic fibers.
be achieved by simply tying the fibers into small loops of different radii. For weak or large diameter fibers, graphite jigs can be used to provide the pure bending constraint (cf. Fig. 3). After treatment, the fiber loop is broken or the jig pieces separated, revealing a fiber with a creepinduced curve of constant radius Ra. For all fiber types tested in this study, the radius ratio Ro/Ra showed a strong dependence on T, a weak dependence on t, and effectively little or no dependence on E*. This strain independence was to be expected based on grain boundary sliding as the controlling creep mechanism. Assuming after treatment that the local creep strains within the fiber were directly proportional to the local applied strains and inversely proportional to R~, it follows from eqn (2) that the radius ratio could then be simply used to measure the bend stress relaxation ratio m; that is: (8)
m = 1 - Ro/Ra
Table 1. Polycrystalline fiber specimens Fabrication processes
Type
Source
Primary phase
Other phases
/3-SIC
fl-SiC fl-SiC re-SiC
C, Si (C substrate) C, 12 Wt%O C, 0-5 Wt%O B, C
Ai203
Chemical vapor deposition
SCS-6
Polymer spinpyrolysis Polymer spinsinter (+ powder) Slurry spin
Nicalon (Hi-NiC)" (DC)" (Carb)"
Textron specialty materials Nippon Carbon Nippon Carbon Dow Corning Carborundum
FP PRD-166
DuPont DuPont
aDevelopmental fibers.
fl-SiC
A1203
Approx. grain size (nm)
Ref. nos
100
5, 6
B4C
3 4 3O 2000
7 8 9 10
SiO2 ZrO2, SiO2
300 400
11 12
218
James A. DiCarlo
In comparison to a traditional tensile creep test, the BSR test offers many advantages from both a practical and basic point of view. T M Clearly on the practical side, one can simply and quickly ascertain the onset and relative creep of various fibers without the need for long fibers, mechanical loading devices, large controlled-environment chambers, or chemically compatible sensors. On the basic side, as will be demonstrated and as predicted by eqn (3), BSR m data can also be used to yield fairly accurate predictions of the creep strain ec and stress relaxation ratio m' that will occur in a polycrystalline fiber when tested in a pure tensile mode under the same test conditions. That is, by eqn (3):
(cc)l(ee)
= (1)/(m') - 1 = (1)/(m) - i
(9)
Here, ee = e0 is the applied tensile elastic strain and e~/ee is defined as the normalized tensile creep strain ratio. Results mid discussion BSR results for the six most characterized fiber types are shown in Fig. 4 as a function of reciprocal absolute temperature. Test conditions were a treatment time of l h and average applied strains of about 0.2%. Because of good reproducibility (Am = +0-02), only three to eight specimens were used for each data point. Examining the Fig. 4 results, one can see that even the simple 1-h BSR test can shed much light on the creep performance of current fibers. For example, for each fiber type, one can discern its creep onset temperature and its creep resistance relative to the other fiber types. This fiber comparison can be done Fiber
F-O~ %
0.3
0 0
N~.alon
v a
D(~ Comw~ O~lum
N
PRD-166 FP
ml!
Temperature, °C
1800 1600
4.5
5.0
.3 .3 .2 .1 .1
5.5
1400
(LO
1200 1100 1000
6.5--7.0
7.5
8.0
900
8.5 9.0X10-4
Reciprocal tempecature, K -1
Fig. 4. One-hour bend stress relaxation results for a variety of polycrystalline SiC and A1203 fibers.
in terms of stress relaxation at a certain temperature or in terms of temperature capability to achieve a certain degree of stress relaxation. Thus, for example, in regard to temperature capability, at all stress relaxation levels the sintered SiC fibers rank highest and the A1203 fibers rank lowest. The somewhat better creep resistance of the PRD fiber relative to the FP fiber can be explained by the ZrO2 additions which inhibit grain boundary motion. 12 For the SiC-based SCS-6 and Nicalon fiber types, creep onset begins as low as the A1203 fibers. For the SCS-6 fiber, free silicon remaining during the chemical vapor deposition process appears a likely source for enhanced diffusion at the grain boundaries. 6 For the polymerderived Nicalon, excess oxygen from the curing stage and a microcrystalline character from low-temperature pyrolysis are likely contributors to its poor creep resistance. 7's As for the sintered SiC fibers, it would appear that the absence of silicon and oxygen at the boundaries does much to improve their temperature capability. As will be discussed, the difference between the sintered or-SiC and r-SiC results are probably related more to grain size than to phase. Although the data of Fig. 4 may be useful for estimating short-time fiber performance, many applications will require service times of 1000h and beyond. To understand better the behavior for these and other applications, the BSR measurements were extended to 100-h treatments and to applied strain levels up to 1%. These results were then empirically modeled to calculate accurately performance within the envelope of test conditions and to predict performance beyond this envelope. When treatment times were increased an order of magnitude, for example, to 10 and 100 h, the characteristic S-shaped curves of Fig. 4 shifted a specific A(1/T) to the right (lower temperatures) depending on the fiber type. This behavior is characteristic of a thermally activated process controlled by an activation energy Q that is effectively constant throughout the main part of the S-curve. For one order of magnitude time change: Q = 2.303R/A(1/T)
(10)
where R is the gas constant (8.314 J/mole). Calculated Q values for the various fiber types are listed in the last column of Table 2. Relatively, these values increase with increasing creep resistance. Absolutely, the Q values should be related to the diffusion mechanisms controlling grain boundary sliding. Although measured primarily in Stage I, the Q values for the sintered or-SiC fiber and the AI203 fibers agree closely with the activation energies measured in Stage II for bulk 0r-SiC and bulk A1203, respectively. 15"~6 Regarding modeling of the BSR data, previous tensile creep studies on SCS-6 fibers6'17 showed that in the creep region of interest (ec> 1%), the performance of these fibers could be accurately described by
219
Creep limitations o f current polycrystalline ceramic fibers Table 2. Best-fit parameters for creep and stress relaxation
Fiber
SCS-6 (BSR) SCS-6 (tensile) Nicalon Dow Coming Carborundum PRD-166 FP
Approximate fiber modulus E (GPa)
Aoa
400 400 200 400 400 370 370
2.8 x 105 1.0 x 10"~ 2.2 x 10s 69 x 105 0-83 x 105 7 x 10l° 49 x 10"'
n
B
Correlation
(K)
factoff
pb
Controlling creep energy,a
Q
(kJ/mol) 1-1 1.1e 1.2 1-5 1.2 1.35 1-5
0.37 0-37 0-40 0.45 0.31 0-53 0-58
25 800 + 950 25 80(Y 24 200 + 1 400 35 133 + 2 450 31 600 -t- 1 800 41 000 + 2 120 41 150:1:2 460
0.991
580 + 38
0-980 0-962 0.971 0-964 0.959
500 + 65 670 + 90 830 + 145 640 + 70 590 + 76
"Applies for stress o in gigapascals, time t in hours, and temperature T in Kelvin. bCalculated from p = RB/Q. CFor linear regression to determine B parameter. dCalculated from A(1/T). eAssumed from bend results. Average Stress cra : a0E/2 = 0.5 GPa
the following semi-empirical creep equation: ~
ec = Aoo~[t e x p ( - Q / R T ) ] p = Aoo~t p e x p ( - B / T )
-
1"~
(11)
Here, B = p Q / R and Ao, n and p are empirically determined creep parameters. In a more recent study, TM it was assumed that this equation could be applied to any polycrystalline ceramic fiber in the stress relaxation range m > 0-1 or in the normalized creep strain range co/co < 10. Equation (9) was then employed to convert the measured m data to normalized tensile creep strains which were then used with eqn (11) to calculate best-fit empirical parameters for the SCS-6, Nicalon, FP, and P R D fibers. The results are given in Table 2 together with new additional parameters for the sintered fibers. For use of these parameters with eqn (11), strain is in absolute units, stress in GPa, time in hours, and T in Kelvin. In that same study, ~4 the accuracy of the m conversion (eqn (9)) and the empirical model (eqn (11)) were demonstrated by plotting available tensile creep data and converted m data as a function of absolute reciprocal temperature. Figure 5 represents such a plot for the Nicalon fiber. It can be seen that: (1) at constant test times straight lines can be fit to all data, implying a constant B parameter; (2) the A ( 1 / T ) shift with each decade of time is constant, implying constant Q and p = Q / B R parameters; and (3) close agreement exists between the converted BSR m data and the tensile creep data of Jia et al.lS The same conclusions could be made for the P R D and FP fibers. 14 For the SCS-6 fiber, although the n, p and Q values agreed, the absolute values of E c / e e for the converted bend results were consistently higher than those for the actual tensile data. 17 This is the basis for the different A0 values listed in Table 2 for the SCS-6
" 1
--
o
0 I-I
•
10 100 1}Ten
z
.1
I 5.0
5.5
I
I
I
I
I
6.0 6.5 7.0 7.5 8.0 Reciprocal temperature, K -1
I 8.5
I 9.0x10"4
Fig. 5. Temperature-time dependence of normalized tensile creep strain for Nicalon fibers. Open points predicted from bend stress relaxation data; closed points are actual tensile data of Jia et al. ,8 fiber. This can be explained by the fact that, unlike the other fibers, the SCS-6 fiber is not homogeneous. That is, it contains not only a carbon substrate core but an outer SiC sheath which creeps or stress relaxes more than the inner sheath. Thus stress relaxation in bending should be greater than that in tension. In summary, the current status of six fiber types can be described by eqns (11) and (9) and the appropriate creep parameters listed in Table 2. These descriptions should be fairly accurate for the conditions: 0 - 9 > m > 0 . 1 or 0 - 1 < e c / e e < 10. Since these conditions fall within or near the upper and lower creep goals previously discussed, this methodology will be used in the next section to provide estimates of the potential for each fiber type in our generic composite application. PERFORMANCE
POTENTIAL
Given their current microstructures, the creep performance of the polycrystalline fibers can now be
220
James A. DiCarlo Assumed conditions: ¢o = 100 MPa; t = 1000 hr L FP r / / / / / . / , / / / / J
PRD-166 ~ / / / / ' / / / / / / / 1 Nicalon
I
~ / ' ~ ~ "
scs-e
Creep limits
I
J
1771 m _>0.5 I I =-.¢ _< 1%
I
//////////////A
C=~oru.dum
////////////////A I
>2GPaS
I
Dow C o ~ n g
I
Strength ~1_._j C A R B ~ ' ~ f 0015OO --
////////////J
0
Assumed conditions: ~o = 400 MPa t = l hour 1600 -- m=0.5forstressrelaxatJon, or~c/~o=l forcreep
I
~.1400 _
I I I
I
I
I
I
/
I
200 400 600 800 10001200140016001800 Temperature, °C
_
Fig. 6. Estimated fiber upper use temperatures under the assumed service conditions and creep limits. compared against the two creep limit goals of ec = 1% and m =0.5. Since creep depends on the time, temperature, and applied stress conditions, one must choose two of the three conditions in order to determine the third. For example, one might assume a particular operating temperature and stress in order to determine maximum service life, or assume a certain service life and temperature and calculate the maximum operating creep stress. However, since ceramic composites are being examined for their potential to operate at high temperature, the approach taken here will be to assume a common service life of 1000 h and a realistic applied stress of 100 MPa, and then estimate maximum or upper fiber use temperatures using the Table 2 parameters. The results are shown in Fig. 6 with the hatched bars for the lower limit of m =0.5 and the open bars for ec= 1%. Examining Fig. 6, one finds for both limits a similar ranking as indicated by the simple 1-h BSR data of Fig. 4. That is, in terms of temperature capability, the A1203 fibers rank lowest and the sintered SiC fibers highest. Also it can be seen that some temperature gain can be achieved in the unlikely situation in which a composite upper creep limit of 1% yields a fiber limit of the same value. It should be realized that these temperature gains quickly diminish when the applied stress is taken above the selected value of 100 MPa. This, however, will not be the case for the lower limit temperatures since rn is effectively independent of stress. Thus, the m = 0.5 temperatures probably represent the most conservative estimates for the six fiber types. Assuming this, one concludes that although the sintered o~-SiC fiber currently represents the best fiber for most structural ceramic composite applications, its potential appears to be limited to temperatures less than 1400°C. Given the above observations, the question arises whether improvements can be made in the other SiC fiber fabrication processes in order to achieve creep resistances comparable or better than that of the a~-SiC fiber. To shed light on this question, its of
S / ~
1100
"-- Sintered
DC ..,,~//////~ "
~-s~c
; th ko :
" / / t ( dr°.'-
• Hi-NiC ONiC
I 10
0 SCS I 102 Grain size d, nm
103
........I 104
Fig. 7. Grain size effects on SiC fiber upper use temperatures under the assumed conditions. Hatched area represents estimate of grain-size conditions for optimum creep resistance (solid line) and for strength greater than 2 GPa (dashed line), interest to examine Fig. 7 which plots upper use temperatures versus grain size for the SiC fibers. The assumed conditions of 1 h and m = 0.5 allow the 1-h BSR measurements to be used for the four SiC fibers plus the Hi-Nicalon fiber for which limited BSR data currently exist. In Fig. 7 one can see a general trend in which increased grain size results in improved creep resistance. As previously discussed, this is to be expected because increased grain size should increase diffusion distances, thereby slowing down grain boundary sliding and macroscopic creep. The solid line is drawn assuming that the tr-SiC fiber data point represents the ultimate in creep behavior and that grain size d can be introduced into the semi-empirical eqn (11) by the relation: A,, = A
~ / d 0"4 ~
A1/d p
(12)
The choice of an inverse d p dependence with p = 0-4 is based on the average p values observed for all the SiC fibers and the fact that grain boundary sliding in the early creep stages ( m > 0 . 4 ) is predicted theoretically2 and observed experimentally~9 to give a creep strain which is a function of the parameter q), where: (a = ( t / d ) e x p ( Q / R T )
(13)
It is interesting to note in Fig. 7 that the sintered fl-SiC fiber data points fall very near the solid line and the data points for all other SiC fibers fall below the line. This suggests that for the assumed conditions the solid line may indeed be the optimum performance to be expected from SiC polycrystalline fibers with equiaxed grains.
Creep limitations of current polycrystalline ceramic fibers
Improvement approaches Assuming then that the solid line of Fig. 7 represents near optimum performance for polycrystalline fibers, what direction should each fiber fabrication process take in order to produce a more creep resistant product? Clearly, the sintered fl-SiC fiber (DC) can move up line by increasing grain size. However, there is a limit because increasing grain size will eventually reduce fiber strength when the largest grain or flaws of this grain size become strength controlling. The vertical dashed line indicates the approximate grain size, 1000 nm, when SiC fiber strength is about 2 GPa, a value often quoted as a minimum fiber strength requirement for structural ceramic composites. This estimate is simply made using Grittith's Law and a SiC fracture toughness of 3 M P a m ~2. Since creep is controlled by the average grain size and strength by the largest, the possibility exists that the sintered fl-SiC has already reached its creep limit, unless perhaps its grain-size distribution can be narrowed. Thus, the hatched area in Fig. 7 represents a general grain-size envelope that fiber manufacturers should seek for optimum creep and strength performance. Even for the best performing sintered tr-SiC fiber, current grain size appears too large to meet fiber strength requirements. Although creep performance will suffer somewhat, the fabrication processes for this fiber should move in a direction which reduce average d to below 1000 nm while still maintaining a narrow distribution. For the chemical vapor deposition and polymerprecursor fibers, improvements appear to lie first in removing low-viscosity grain boundary phases, such as, free silicon and Si-O complexes, and perhaps amorphous carbon. Both directions appear possible by changes in processing conditions. Indeed, by using radiation curing, the Hi-Nicalon fiber has significantly reduced oxygen content with respect to the Nicalon fiber, s The payoff in better creep performance can be clearly seen in Fig. 7. If all oxygen can be removed, behavior similar to the polymer-derived and sintered fl-SiC fiber should be expected. Figure 7 suggests that as oxygen is removed, grain size increases in a beneficial manner. What processing improvements are needed to further increase grain size have yet to be determined. Two final points concerning process improvements need to be mentioned. First, to evaluate whether any process change has led to improved creep behavior, almost immediate information can be obtained by using the simple BSR test and the Fig. 7 solid line. That is, a few 1-h tests at different temperatures and at an approximate surface strain of Eo = 0.2% could determine whether any relative improvements in the m ratio have been made. The results can also yield an estimate of the upper use temperature for m = 0.5 which with grain-size information could then be
221
compared with the Fig. 7 line to determine whether optimum performance has been achieved. A second important point is that optimum creep performance in this study was derived empirically using primarily SiC fibers with equiaxed grains. Thus, as new fibers are developed, the concept of optimum performance may have to be redefined. Based on the grain boundary sliding mechanism, it may be possible to introduce different boundary conditions or second phases which inhibit sliding more than those currently present in the sintered SiC fibers. Likewise, grain shapes or aspect ratios can be changed with the long axis parallel to the fiber axis. By increasing diffusion distances, this approach has been successfully employed to improve the creep resistance of tungsten filaments while still maintaining strength properties. Indeed this method is currently being attempted to improve the creep resistance of polycrystalline AI203 fibers. 2° If the grain aspect ratio is taken to the extreme, one would then have single crystal fibers which should provide the optimum in creep performance, perhaps increasing upper use temperature capability by as much as 200°C. ~ However, although a commercial processing approach has been developed for continuous single crystal AI203 fibers, 2~ no method currently exists for continuous single crystal SiC fibers.
CONCLUDING REMARKS This study has attempted to demonstrate the current situation for using creep-prone polycrystalline ceramic fibers in structural ceramic composites. In so doing, some important issues have been exposed which will require further in-depth studies. One centers on the need to develop accurate fiber creep requirements from the performance requirements of the composite component. From a macroscopic point of view, the service conditions that the component experiences can be a complex variation of time- and spaciallydependent temperatures and stresses. How these influence fibers on the microscopic level adds a further degree of complexity, especially in light of the creep mismatch problem and the fact that fiber architectures can be very complex. An added problem is the lack of numerical or analytical codes to model and analyze time-dependent constituent behavior. In light of these complexities, it would appear that the best approach at the present time is to identify, as was done in this study, the fiber composition, fabrication process, and microstructure which gives the best creep performance and then to actually evaluate these fibers within the composite under the intended use conditions. The performance equations derived here may be of some use in setting approximate limits on the service conditions, but these certainly cannot replace
222
James A. DiCarlo
observation of performance under actual operating conditions. Another issue requiring further study is that there appears to be an optimum creep performance for equiaxed SiC fibers which, under the conditions of practical interest, is weakly dependent on grain size. Currently, all the tested SiC fibers fall at or below this optimum, suggesting possible directions for processing/microstructure improvements. Thus, opportunities apparently exist for obtaining more creep-resistant SiC fibers from a variety of fabrication processes. However, even if these performance gains can be achieved, upper use temperatures may be only about 1300°C for a 1000-h service life. This temperature will increase or decrease by about 60°C for an order of magnitude decrease or increase in service time, respectively. Thus, questions arise whether processing improvements in polycrystalline SiC fibers are indeed possible and whether these improvements are enough to significantly impact the use of these fibers in high-temperature structural ceramic composites. A negative answer to this latter question would then seem to imply the need to develop SiC fibers with directionally aligned grains or perhaps with single crystal microstructures.
ACKNOWLEDGMENTS The author wishes to acknowledge the invaluable professional and technical support of G. N. Morscher and T. C. Wagner.
REFERENCES 1. DiCarlo, J. A., High Temperature Structural Fibers-Status and Needs. N A S A TM, 105174, 1991. 2. Nowick, A. S. & Berry, B. S. Anelastic Relaxation in Crystalline Solids. Academic Press, New York, 1972. 3. Raj, R. & Ashby, M. F., On grain boundary sliding and diffusional creep. Met. Trans., 2 (1971) 1113-27. 4. Todd, R. I. & Hazzledine, P. M., Large anelastic strains at constant volume in superplastic tin-lead eutectic alloy. Scripta Met., 27 (1992) 127-32.
5. Ning, X. J. & Pirouz, P., The microstructure of SCS-6 SiC filter. J. Mater. Res., 6 (1991) 2234-48. 6. DiCarlo, J. A., Creep of chemically vapour deposited SiC fibers. J. Mater. Sci., 2 (1986) 217-24. 7. Yajima, S., Okamura, K., Hayashi, J. & Omori, M., Synthesis of continuous SiC fibers with high tensile strength. J. Am. Ceram. Soc., 59(7-8) (1976) 324-7. 8. Takeda, M. et al., Thermal stability of the low oxygen silicon carbide fibers derived From polycarbosilane. Ceram. Eng. Sci. Proc. 13 (1992) 209-17, 9. Lipowitz, J., Rabe, J. A. & Zank, G. A., Polycrystalline SiC fibers from organosilicon polymers. Ceramic Eng. Sci. Proc., 12 (1991) 1819-31. 10. Frechette, F., Dover, B., Venkateswaran, V. & Kim, J., High temperature continuous sintered SiC fiber for composite applications. Ceram. Eng. Sci. Proc., 12 (1991) 992-1006. i1. Stacey, M. H., Developments in continuous aluminabased fibres. J. British Cermic Trans., 87 (1988) 168-72. 12. Romine, J. C. New high-temperature ceramic fiber. Ceramic Eng. Sci. Proc., 8(7-8) (1987) 755-65. 13. Morscher, G. N. & DiCario, J. A., A simple test for thermomechanical evaluation of ceramic fibers. J. Am. Ceram. Soc., 75(1) (1992) 136-40. 14. DiCarlo, J. A. & Morscher, G. N., Creep and stress relaxation modeling of polycrystalline ceramic fibers. ASME, A D 22, A M D 122 (1991) 15-22. 15. Sargent, P. M. & Ashby, M. F., Deformationmechanism maps for silicon carbide. Scripta Met., 17 (1983) 951-7. 16. Cannon, R. W. & Langdon, T. G., Review: Creep of ceramics. Part 1. Mechanical characteristics. J. Mater. Sci., 18 (1983) 1-50. 17. Morscher, G. N., DiCarlo, J. A. & Wagner, T., Fiber creep evaluation by stress relaxation measurements. Ceram. Eng. Sci. Proc., 12(7-8) (1991) 1032-8. 18. Jia, N., Bodet, R. & Tressler, R. E., Creep and creep rupture behavior of Si-C-O and S i - N - C - O based continuous fibers. Paper presented at 93rd Annual Meeting, American Ceramic Society, 28 April-2 May 1991, Cincinnati, OH. 19. Ke, T. S., Experimental evidence of the viscous behavior of grain boundaries in metals. Phys. Reo. 17 (1947) 533-46. 20. Tressler, R. E. Penn State University, personal communication (1990). 21. Hurley, G. F. & Labelle, H. E., Elevated temperature strengthening of melt-grown sapphire by alloying. In Applied Polymer, Symposium No. 29, New and Specialty Fibers, ed. J. Economy. John Wiley, New York, 1976, pp. 131-49.