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SiC composite
MATERIALS SCIENCE & ENGINEERING
ELSEVIER
Materials Scienceand Engineering A230 (1997) 171- 182
Tension-tension
fatigue of a cross-woven C/Sic composite Mingde Wang, Campbell
Department
of M&vkds
Science
md Etlgiueering,
Unicersit~~
Laird *
of Penng~lrn~~in,
Philadelphia,
PA 191046272,
C;S’A
Received2 November1996;receivedin revisedform 12November1996
Abstract The tension-tension fatigue behavior of a crosswoven C/Sic compositewas studiedin termsof damagemodesand damage development.The fatigue stressversuslife diagram(S-N) curve and an endurancelimit of 320-340 MPa (about SO’% tensileUTS) for IO6cycleswereobtained for the C/Sic composite.Different fatigue behaviorswerefound for samplesthat failed during fatigue and for samplesthat survived 10” cycles. Sevenfatigue damagemodeswere observed,the developmentof which were usedto explain the different fatigue behaviors.For the fatigue-failed samples,the degreesof damageof the sevenmodesincreasedwith increaseof cycles,leadingto an increasein elapsedstrain and a decreasein compositemodulus.For fatigue-survivedsamples,the developmentof all the damagemodesexcept for fiber breaking causedan initial increaseof elapsedstrain and decreaseof compositemodulus,but at high cycles,fiber bundle realignmentand straighteningin thesesamplesled to partial recovery of the modulusand cessationof the damagedevelopment.0 1997Elsevier ScienceS.A. Kqwords:
C/Sic composite; Fatigue;Fiberbundle
1. Introduction
but also to predict the cyclic stress-strain loops [lo? 111, and to estimate interFacial friction by C. Cho et al. [12].
Fiber reinforced ceramic matrix composites (CMCs) have been under extensive development as potential high temperature materials because of their possible high strength and toughness at both room and high temperatures. Since the applications of these composites are likely to include jet engines and heat engines, the cyclic fatigue behavior of these composites has been the focus of a number of studies, especially in composite systems with unidirectional or cross-ply fiber arrangements such as Sic fiber (Nicalon)/lithium-aluminum-silicate (LAS) composite [l-7]. In those unidirectional or cross-ply ceramic composites, there seems to exist a definite fatigue limit, which approximately coincides with the proportional limit of the 0” plies in monotonic loading. In those composite systems, the most studied fatigue damage mode has been the same as in monotonic loading, i.e., matrix cracking. Because of the simple fiber arrangement in the unidirectionalcross-ply composites, it has been possible not only to relate matrix cracking during fatigue to the reduction of composite stiffness through mechanical modeling [8,9],
The fatigue failure mechanisms in the fiber reinforced CMCs are associated with fiber-matrix interface degra-
*Correspondingauthor. Tel.: + 1 215 8986664;fax: + 1 215 5732128; e-mail:[email protected] 0921-5093/97/$17.00 0 1997ElsevierScienceS.A. All rightsrcsrrved. HISO921-5093(97)00018-X
dation and degradation of fiber strength. A recent overview paper by A.G. Evans et al. [13] deals with interface degradation mechanism, giving a detailed
analysis and modeling of the fatigue behavior of unidirectional and cross-ply CMCs. In ceramic composites with fiber reinforcements in the form of textile preforms such as cross woven C/Sic composites, cyclic damage is a more complex process. For example; in an exploratory study of a cross woven C/Sic composite by Wang et al. [14], ply delamination and matrix wear between plies and fiber bundles were observed and found to be important in fatigue failure
of this composite, in addition to matrix cracking. These damage modes were also found to be more gradual and localized during fatigue than in monotonic loading. Such fatigue damage modes were identified in a study on a similar type of cross woven C/Sic composite by Shuler et al. [15]: and the mechanisms for these damage modes were proposed to result from non-uniform stress-strain distributions produced by the cross weaving of the fiber bundles in the composite. As in unidirectional CMCs, the modulus reduction was related
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largely to the matrix cracking in that study, and other fatigue damage modes were observed from the fracture surfaces. Damage development under cyclic loading was demonstrated as it developed. In a recent study on a slightly different cross woven C/Sic composite by Wang et al. 1161, six damage modes were identified during monotonic tension due to the complex composite microstructures, and they were found to be responsible for the mostly nonlinear stress-strain behavior of the composite, which is shown in Fig. 1. As an extension to the monotonic tensile behavior, the purpose of the present study is to investigate how the complex damage develops under cyclic loading conditions and the effect of the cyclic damage on the fatigue behavior of a typical textile ceramic composite, In the present study, the fatigue stress versus life relation was first obtained for an unbalanced cross woven C/Sic composite. The fatigue behavior of the composite, represented by using elapsed strain and modulus, was found to be different for fatigue stresses above the fatigue limit and at the fatigue limit. Fatigue damage and damage evolution were examined and used to explain the differences in the fatigue behavior at stresses above and below the fatigue limit.
2. Materials and experimental procedures
The experimental composite used in this study is an unbalanced cross woven C/Sic composite supplied by DuPont. The microstructure of the composite was characterized in detail previously [16]. The interested reader is referred to that paper for details of the material. Here? only those microstructural features relevant to the present study are briefly summarized as follows.
Fatigued
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Fig. 1. The stress--strain curve of a typical virgin C/Sic composite. together with the stress-strain curve of a fatigued survived sample. The off-set strain in the strain axis for the latter is equal to the residual strain accumulated through fatigue.
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Fig. 2. Longitudinal cross section of C&C composite, showing coating layers on both sample surfaces and interbundle voids (A and B). Bundles X and Y were laid up mostly through contacts between transverse bundles while Y and Z have extensive degrees of longitudinal contacts. Notice that the travel of bundle C ends at D, indicating bundle misalignment with respect to the polished plane. E and F indicate bundle sections with differeni matrix distributions.
The typical microstructure of the C/Sic composite can be seen in a longitudinal cross section: as shown in Fig. 2. Several features can be seen from this figure. First, the undulations of the longitudinal bundles in different plies are not in synchrony with each other (e.g., the first and second bundles from the top surface in Fig. 1). a situation termed ‘random phase lay up’, Second, as a consequence of the random phase lay up, the contacts between the plies in the composite are highly irregular, i.e., some plies contact through transverse bundles iX and Y in Fig. l)? and other plies are laid up through various degrees of contact between longitudinal bundles (Y and Z in Fig. 2). Third, this random phase lay up also caused a non-uniform distribution and size variation in the interbundle voids (e.g., A and 3 in Fig. 2-white areas). Most of these voids are oriented in the longitudinal direction (e.g., A in Fig. 2)> and their sharp tips can act as sites for local ply delamination under cyclic stresses. Lastly, a less obvious feature in the composite is the misalignment of some plies (and bundles in those plies) with respect to the polished plane of the sample (the loading direction in the fatigue tests was parallel to this plane), as evidenced in Fig. 2 by the discontinuation of the travel of the bundle marked C from right to left at point D. In fact, all the plies have a certain degree of misalignment due to the random phase lay-up of the fabrics. The longitudinal bundles in the composite consist of three typical kinds of sections with different matrix distributions, as described in detail 1161. The first section, which exists between crossover of the bundle and is not in contact of the adjunct plies (e.g., Section E in
Fig. 2), has a continuous matrix distribution, and is thus termed ‘matrix-continuous’ section. The second cross section, which is often positioned in the crossover of the fabric and is also pressed by the adjunct ply (e.g., Section F in Fig. 2), has a discontinuous matrix distribution and thus, is termed ‘matrix-discontinuous’ section. In this section, the fibers are densely packed (i.e., mostly fiber-fiber contact) and the matrix exists as isolated ‘matrix pockets’, filling the interstitials of the dense fiber pack. The third section is a ‘mixed section’ consisting of the above two, i.e., one part of the section has continuous matrix distribution and the other has discontinuous matrix distribution. Because of the different matrix distributions in a fiber bundle with the same number of 3000 fibers, the cross section area of the ‘matrix-continuous’ section is larger than that of the ‘matrix-discontinuous’ section (e.g., see Sections E and F in Fig. 2). Therefore, the present composite possesses a high degree of variability in microstructure at both composite level and fiber bundle level, which can lead to complex damage modes such as ply delamination. Because of the complex microstructures of the composite, the volume fractions of the fiber, matrix and void can only be approximated (see Ph.D. Thesis by M. Wang [17]) and are about 53, 34 and 13%, respectively. In addition to its complex microstructure, the composite is coated with thick Sic coatings on both surfaces of the composite plate (Fig. 2) and thin SIC coatings over fiber bundle sections which are open to the interbundle voids (e.g., Section E in Fig. 2). Matrix cracks were found in these coatings and in the transverse fiber bundles in the as-received composite [16,17]. Since these cracks resulted from the CVI processing, they are termed processing-induced cracks. The crack populations can be quantified as ‘surface coating crack density’ and ‘bundle coating crack density’. The measurements of such densities, which were also described in [17], are represented by numbers of cracks per unit length (no. mn- ‘). These crack densities were obtained both for as-received samples and for samples fatigued at all the stress levels conducted in this study.
For cyclic fatigue tests, straight-sided composite specimens 65 mm long, 4 mm wide and 4.4 mm thick were used. Both ends of the specimen were tabbed using clear polyester resin for gripping. For tension-tension fatigue tests with R = 0, in which tension was applied in the ‘longitudinal’ direction, an Instron 1361 screwdriven universal testing machine was used, which was equipped with a MTS hydraulic grip system. Strains were measured using an Instron dynamic extensometer with a gage length of 10 mm. To avoid extensometer slippage during a test, the extensometer knife edges
were glued to the specimen using a 5 min epoxy. The load and uniaxial displacement during a fatigue test were recorded using a X- Y recorder and a data acquisition system running in an IBM PC. Pulsating tension fatigue tests were conducted in load control using a half sinusoidal wave at 1 Hz. Based on the monotonic strength of 405 MPa for the experimental composite (see Fig. l), the controlled load levels were chosen to be 380, 360, 340 and 320 MPa. A typical test was run continuously until the sample failed (i.e., fractured in the gage section, otherwise the result was considered invalid) or survived 10” cycles. The fatigue life of a failed sample was counted to the last cycle number at which the sample fractured. For the fatigue-survived samples, a tensile test was finally performed at a loading rate of about 5 MPa s- ’ (i.e., the same as used in monotonic tests) to determine the residual strength. During fatigue tests, individual stress-strain loops were recorded for the following cycles: l> 2, 4, lo> 20, 10, lo’, 2 x lo’, 4 x lo’, lo”, 2 x lo”, 4 x lo’, lo”, 2 x lo’, 4 x lOa, 105, 2 x lo”, 4 x 10’ and 10h. From these loops, the elapsed strains, which are defined as the accumulative peak strain at the above selected cycles, can be obtained. The composite modulus can also be obtained by linear fitting of the loading portion of the stress-strain loops except for the first cycle. The cycle intervals chosen were found to describe the fatigue behavior in a representative f’ashion. In addition, the strain versus time during fatigue was recorded continuously on a chart recorder. From this recording, the elapsed strain for the last cycle before fatigue failure was obtained for comparison among the fatigued samples and with the tensile failure strain. In order to understand the fatigue behavior of the composite, fatigue damage evolution was monitored at two stress levels, i.e., 360 and 320 MPa, which are above and at the fatigue (or endurance) limit respectively. Side-polished samples were examined and areas of interest were selected in the SEM before loading. Subsequently, the same areas of interest were observed (1) at cycles 1, 10 and 10’ for a sample fatigued at 360 MPa; (2) at cycles 1, lo”, 10” and lo6 for the sample fatigued at 320 MPa. The fracture surfaces of fatiguefailed samples were examined using SEM and the results for a typical fracture are also reported here.
3. Tension-tension composite
fatigue behavior of C/SC
The fatigue stress versus life diagram (S-N diagram), obtained from a total of 12 specimens, is shown in Fig.
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3. From this figure, it is seen that a fatigue limit’ (or more accurately endurance limit) for lo6 cycles exists at stresses 320-340 MPa for the subject C/Sic composite. For present purposes, we define this stress as the IO6 cycle endurance limit. Specifically, the two samples fatigued at 320 MPa plus one sample fatigued at 340 MPa survived lOh cycles. These three samples are referred to as ‘fatigue-survived’ samples. All the other samples, which failed before lo6 cycles, are referred to as ‘fatigue-failed’ samples. It is also seen from Fig. 3 that, in general, the average fatigue life decreases as the fatigue stress increases. However, and as usual, in the S-N diagram, there is some degree of scatter, which indicates the sensitivity of the fatigue life to the microstructure variability in the C/Sic composite samples.
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3.2, Cyclic stress-stmin behavior
In order to understand the fatigue behavior of the C/Sic composite in detail, it is necessary to study the hysteresis stress-strain loops of the fatigued samples. Fig. 4 shows the typical hysterisis stress’strain loops of a fatigued sample at selected cycles. From Fig. 4, it is seen that the loading portion of the first cycle is largely non-linear, similar in shape to that of a monotonic tensile test (Fig. 1). The first cycle shows an open loop with large residual (irreversible) strain when the sample was unloaded. In the subsequent cycles, the stress-strain
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LIFE (Cycles) Fig. 3. S-N curve (fatigue stress vs, life curve) for C/Sic composites fatigued at 1 Hz. The solid lines indicate the possible range of experimental data, and show the trend of life behavior. The point at 10° cycles represents the monotonic strength.
’ The solid lines in this figure were drawn to indicate the trend of the relation between fatigue stress and life, and are based on the scatter of the experimental data, as well as an assessment of fatigue scatter in many material systems. Clearly, -there are insufficient specimens to establish high cycle endurance behavior on a proper statistical basis.
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Fig. 4. Stress-strain loops at selected cycles for a fatigue stress of 360 MPa (sample F-4). This behavior is typical of the stress-strain loops for all the fatigue samples. Definition of elapsed strain is indicated in this figure.
loops appear closed. In contrast to the first loading, the typical loading curves of the subsequent cycles are mostly ‘linear’ except for the very beginning. It is also seen from Fig. 4 that fatigue loading produces cyclic creep* in the composite (which is indicated by rachetting of the loop peaks). In the meantime, the loops tend to tilt downwards as the cycle count increases. Therefore, it is useful to choose the elapsed strain to describe the cyclic creep, and the composite modulus to describe cyclic damage for the fatigued samples because the former represents the dimensional changes, and the latter is both an indication of damage as well as an important design parameter for practical engineering applications. In order to compare the changes of the modulus among the fatigue samples, the modulus was normalized with respect to the average tensile modulus of 137 GPa and then used in the modulus versus cycles plots (Fig. 5(a)-(d)). This approach was adopted because the tensile modulus is characteristic of the C/Sic composite [16,18], and because the limited linear part of the first fatigue loading curve did not contain enough data to obtain a reliable value of the modulus for the individual specimens. 3.3. Variations of elapsedstrai~z and modztlzts lrlith cycling
Fig. 5(a)-(d) show the elapsed strains and normalized composite moduli versus cycles for samples fatigued at 380, 3609 340 and 320 MPa, respectively. From Fig. 5(a)-(d), it can be seen that, for the fktiguefailed samples, in general, the elapsed strains increased continuously from cycle one, while the normalized ’ We employ this term as conventionally used for metal specimens
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Fig. 5. Elapsed strain and normalized modulus vs. cycles for fatigue samples at: (a) 380 MPa; (b) 360 MPa; (c) 340 MPa; and (d) 320 MPa.
modulus decreased steadily after a large drop at cycle one. For the fatigue-survived samples (two at 320 MPa in Fig. 5(d) and one at 340 MPa (F17) in Fig. 5(c)), the elapsed strain increased only modestly in the beginning, and then, stabilized at high cycles (lo4 for F17 and 340 MPa and lo5 at 320 MPa. Also note that in Fig. 4(c), the elapsed strain for samples F17 and F20 appear to decrease after 10” and 10’ cycles, and this is considered to be experimental error caused by possible slippage of the extensometer during fatigue). For these samples, the relationship between the normalized modulus and fatigue cycles is very interesting, as indicated for sample F17 in Fig. 5(c) and the two samples in Fig. 5(d), i.e., the modulus decreased rather modestly after the big drop in the first cycle, but subsequently the modulus recovered at high cycles ( > lo5 cycles).
It is interesting to note that the elapsed strains of all the fatigue-failed samples at one cycle before failure seem to approach a value of about 0.7%, regardless of fatigue stress levels and fatigue life (Fig. 5(a)-(c)). Since this critical elapsed strain of 0.7% for fatigue failure is comparable to the average failure strain of the composite under monotonic tensile loading (which is about 0.69%, see Fig. l), it is reasonable to suggest that the failure of the C/Sic composite is strain-controlled, i.e., no matter what the loading condition is, once the failure strain is reached, the composite will fail. On the other hand, for the fatigue-survived samples, since the elapsed failure strain of about 0.7% was not reached during fatigue, they survived the fatigue test for lo6 cycles. Therefore, the elapsed strain is a useful parameter which can be used to describe cyclic creep and to
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monitor whether the composite will fail under other testing conditions such as creep. Fig. 5(a)-(d) also show that there exists a relatively large data scatter in the change of both elapsed strain and normalized modulus with cycles at the same fatigue stress. Since all the fatigue tests were carefully conducted using the same sample size and test procedures, this large data scatter can only indicate that the fatigue behavior of CiSiC composite is sample specific, i.e., the fatigue behavior is very sensitive to microstructurally-induced differences that may exist from sample to sample. 3.4. Damage and fktwe
during cyclic iondhg of
C/Sic composites
In order to understand the different fatigue behaviors of the C/Sic composite described above. damage evolution and fracture during tension-tension fatigue were examined at two stress levels. The first stress level was 320 MPa, representing fatigue-survived samples, and the second, 360 MPa. representing fatigue-failed samples. In general, a total of seven fatigue damage modes were observed in both samples. These damage modes are: (1) matrix cracking; (2) transverse bundle cracking; (3) interfacial debonding; (4) fiber breaking; (5) ply delamination; (6) bundle splitting and (7) matrix wear. All the damage modes except matrix wear (item 7) were observed after the first fatigue cycle at both stress levels. Since the six damage modes produced in the first cycle are the same as those observed in monotonic tensile tests r16], they are not shown here in detail, instead they will be pointed out in the following discussion. After the first cycle, all the damage modes accumulate in a subtle manner: leading to a gradual increase in elapsed strain and decrease in modulus, as indicated in Fig. 5(a)-(d). Matrix cracking, observed during fatigue at the two stress levels, includes matrix cracks in the longitudinal bundles and in the composite surface coating. Since the longitudinal bundles are the major load carrying elements in the composite for the fatigue tests, matrix cracking behavior was examined mainly in these bundles. As described in Section 2.1, the longitudinal bundles are composed of ‘matrix-continuous’ sections. ‘matrix-discontinuous’ sections and ‘mixed’ sections, thus, matrix cracking in these typical sections was monitored during fatigue. In the matrix-continuous section, matrix cracks run across most of the bundle in a mostly continuous and planar manner as shown in Fig. 6. In the matrix-discontinuous section, matrix cracking occurs in the form of matrix damage zones, which consist of small separated matrix cracks inside the bundle, as shown in Fig. 7(a). The matrix cracking behavior in the ‘mixed’ section is a mixture of the above two mechanisms 1171. These types of matrix cracking formed in the first cycle, which are the same as those observed in the monotonic tensile test [16]. With in-
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crease of cycles, however, matrix cracking continues to occur mainly in terms of new matrix damage zones in the ‘matrix-discontinuous! sections and ‘mixed’ sections of the longitudinal bundles, as shown in Fig. 7(b). Matrix crack multiplication was found to play an important role in the decrease of the composite modulus in the tensile test of the present composite [16], and in the fatigue of ceramic composites including unidirectional Nicaloniglass ceramics [8,9], and across woven composite [15]. To determine whether a similar relationship exists in the fatigue of the present composite, the densities of bundle coating cracks and surface coating cracks were measured as a function of cycles for stress levels of 380. 360. 340 and 320 MPa. The results are plotted in Fig. 8, from which it is seen that the crack densities for the longitudinal bundle coating cracks and the sample surface coating cracks saturate in the first two cycles with the biggest increase in the first cycle. The results in Fig. 8 also confirm the observation that the development of matrix cracking after the first two cycles took place inside the longitudinal bundles, i.e., in terms of matrix damage zone formation (e.g., Fig. 7), since the formation of matrix damage zones is not reflected in the bundle coating crack densities. Fig. 8 also shows a large data scatter? indicating again that the fatigue behavior of the composite is sample specific. Longitudinal bundle splitting, which was seen in monotonic tensile loading to the high stress level of 385 MPa [16], was observed at both fatigue stress levels of 320 and 360 MPa. This type of damage was found to occur mostly in the crossover sections of the longitudinal bundle. A typical example is shown in Fig. 9(a) for a sample fatigued for 1000 cycles at 320 MPa, With cycling, the splitting was seen to grow (Fig. 9(a)). Also in Fig. 9(b), matrix wear (in terms of matrix fragmentation) is seen to accompany the longitudinal bundle splitting.
Fig. 6. Matrix cracking in a mostly ‘matrix continuous’ section of a longitudinal bundle. The matrix cracks are continuous and straight across the bundle. The loading direction is horizontal.
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The fracture surface of a typical fatigue-failed sample is shown in Fig. 12(a), which indicates that failure of the composite consisted of fractures of the longitudinal bundles. From Fig. 12(a), it is seen that most of the fractures of the longitudinal bundles appeared to occur at or near the crossover sections of the bundle (e.g., A and B) with few bundles fracturing between crossover sections (e.g., C). Examination of the longitudinal bundle fracture surfaces which occurred at or near the crossover sections of the bundle revealed a ragged fracture surface with limited fiber pull-outs, while those which occurred between the crossovers showed either fracture surfaces with long fiber pull-outs or flat fracture surfaces with little fiber pull-out. Since these types of bundle fracture surfaces are the same as those observed in monotonic tensile fracture [16], they are now shown in detail here. However, in the flat fracture surface of a longitudinal bundle (Fig. 1 l(b)), an interesting feature of the matrix wear, which was closely associated with fatigue loading, was observed. This appears as the rough area on the left of the fractograph, indicated by W. 3.5. Residuul strer?gth
(b) Fig. 7. Matrix cracking in a ‘matrix-discontinuous’ sections of a longitudinal bundle at 320 MPa: (a) after 1 cycle, two matrix damage zones formed (N); (b) after 1000 cycles, another matrix damage zone developed (M).
Ply delamination, which was seen in monotonic tensile loading to the high stress level of 385 MPa [16], was observed at both fatigue stress levels of 320 and 360 MPa. An example of this damage mode is shown in Fig. 10, which was taken after 1 cycle at 320 MPa. Ply delamination was also observed to grow with increase of fatigue cycles for both stress levels of 320 and 360 MPa. Fiber breaking was observed after the first cycle in both samples fatigued at 320 and 360 MPa. However, only the higher stressed sample caused continued fiber fracture with increase of cycles. Again, like bundle splitting, the continued fiber fracture was seen to occur mostly in the crossover section of the longitudinal bundles. A typical example is shown in Fig. 1l(a) and (b), in which continued fiber fractures were seen with increase of cycles, together with the growth of bundle splitting and matrix wear in the sample fatigued at the stress of 360 MPa.
The residual strength of the fatigue-survived samples (two at 320 and one at 340 MPa) is listed in Table 1, together with the accumulative failure strains (i.e., the sum of the residual strain at lo6 cycles and the subsequent tensile failure strain). The tensile stress-strain curve of a typical fatigue-survived sample is shown in Fig. 1. From Table 1 and Fig. 1, it is interesting to see that the residual strength of these fatigued samples increased by an amount ranging from about 10% (F9) to 20% (F15), and the proportional limit of the fatigue sample is higher than that of the virgin sample. Again, the residual strength is sample specific. However, if we 6
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Cycles Fig. 8. Crack densities vs. cycles for the four fatigue stress levels studied. The crack densities were measured on composite surface coating and longitudinal bundle coating. The horizontal arrows positioned at 100 cycles indicate that the samples did not fail at the maximum cycle plotted in this figure.
Fig. 9. Growth of longitudinal bundle splitting at 320 MPa: (a) after 10” cycle, longitudinal splitting appeared (S) and N indicates matrix cracking: (b) after 10” cycles, the splitting grew with the associated production of matrix fragments.
The present C/Sic composite exhibits seven fatigue damage modes at all the stress levels studied here. These damage modes include: (1) matrix cracking, (2) transverse bundle cracking, (3) mterf’acial debonding, (4) fiber breaking, (5) ply delamination, (6) longitudinal bundle splitting, and (7) matrix wear. During the first fatigue cycle, all these modes of damage except for matrix wear (item 7) were observed to take place. For example, the crack densities of the sample surface coating and the longitudinal bundle coating (Fig. 8), which indicate matrix crack multiplication, increased to a level close to saturation soon after the first cycle for all the fatigue stress levels. The damage and damage accumulation. therefore, are believed to have caused the largely nonlinear stress-strain behavior of the first cycle (Fig. 4) leading to a large reduction of the composite modulus (Fig. 5(a)-(d)), similar to behavior during monotonic tensile loading (Fig. 1). Subsequent cyclic loading of the composite caused gradual damage developments in the composite, leading to nearly closed stress-strain loops and gradual changes in modulus and elapsed strain (Fig. 4 and Fig. 5(a)-(d)). The damage developments observed in this study include increases of bundle splitting, ply delamination and matrix cracking. continued fiber breaking, and matrix wear. Of these damage developments, bundle splitting and ply delamination are the unique and important damage mechanisms in the present composite. which influence and are enhanced by other damage mechanisms. Longitudinal bundle splitting is a damage mode that was observed 1.0 increase during fatigue (Fig. 9(a) and (b), and Fig. IO). Longitudinal bundle splitting is a result (also an .indication) of the ‘permanent’ straightening of longitudinal bundles. The mechanism for this is closely related to the cross weave of the C/Sic com-
consider the accumulative failure strain for the fatigue survived samples as the ‘failure strain’. the values of this ‘failure strain’ are about the critical value of 0.7%. comparable to monotonic tensile failure strain and fatigue elapsed failure strain of the CjSiC composite (see Section 3.3), confirming that the fracture of the C/Sic composite is controlled by the total strain. 4. Discussion
The fatigue behaviors of the C/Sic composite reported in the previous section can be understood qualitatively in terms of damage and damage development occurring in the composite during cyclic loading. Through this understanding, the failure behaviors can be related to the complex microstructure of the C/Sic composite. Thus. we point out key microstructural features responsible for important fatigue properties, such as the fatigue limit.
Fig. 10, Ply deiam:nations (II) between longitudinal bundles in C/Sic composites after 1 cycle at 320 MPa.
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bundle shear strength, both of which make it easier for the bundle to split. Like bundle splitting, ply delamination was also seen after the first loading cycle in the composite (Fig. 9). Ply delamination can be caused by either the shear stress generated from unequal displacement between contacting bundles in neighboring plies, or tensile stress normal to the longitudinal bundle due to the straightening of undulated bundles. These two mechanisms can take place in the present cross woven C/Sic composite in which the bundle contacts are highly irregular (Fig. 2) so that both shear stress and normal stress can be generated in these irregularly contacted areas, and both stresses can also be increased by the interbundle voids which have sharp tips oriented in the bundle direction, as shown in Fig. 2. The occurrence of ply delamination and bundle splitting decreases the composite integrity, thus reducing its
Fig. Il. Continued fiber breaking during fatigue at 360 MPa: (a) after 1 cycle, fiber breaking (B), debonding (D) and new cracks (N) are formed; (b) 100 cycles, other fiber breaking (B), matrix wear (W). Bundle splitting growth is indicated by S in (a) and (b).
posite as follows: Under applied stress, the undulation of the longitudinal bundles causes complex stresses which include tension and shear inside the bundle, especially in bundle sections that crossover transverse bundles. In these bundle sections, the fibers most likely contact each other and the matrix is discontinuous, as described in Section 2. l., so that the shear strength is very low. Therefore, when the shear stress caused by bundle undulation under the applied tensile stress exceeds this low shear strength, fracture occurs in the fiber direction, causing the bundle section to split. The bundle splitting damage allows the undulated section to straighten ‘permanently’. With cycling, the stress state in the bundle can change due to damage such as ply delamination and matrix wear, resulting in the growth of the splitting. For example, bundle splitting can be enhanced by ply dalamination which relieves the constraints of the bundles, while matrix wear decreases
(4
@I Fig. 12. (a) Fracture surface of a sample fatigued at 360; (b) Flat fracture surface of a longitudinal bundle with little fiber pull-out looking down the fiber direction (b). Notice the matrix wear at the bundle fracture surface.
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Table 1 Residual strength of the fatigue survivedsamples Sample no.
Fatigue stress (MPa)
Residual strength(MPa)
Residual-tensile
Failure strain (‘X0)
Tensile test F-9 F-15 F-17
N.A. 320 320 340
405 466 491 453
1 l.l5 1.21 1.1’
0.70 0.67” 0.73” 0.69’
“Failure strain for those samples = the sum of the tensile Failure strain obtained in the final tensile test and the residual strain at 10’ cycles.
modulus and increasing its extension. However, their influences on the fatigue behavior of the present CjSiC composite lie in the associated structural adjustments of the longitudinal bundles. For example, the ‘permanent’ straightening of undulated fiber bundles, which is associated with bundle splitting, can redistribute the load amongst the fibers in the bundle. And ply delamination can relieve constraint in bundles that are misaligned with respect to the loading direction (Fig. 2), permitting them to realign locally to the stress axis, and thus allowing the bundle section to bear higher load. These structural adjustments and the associated stress redistribution are the main causes for the increase of matrix cracking and continued fiber breaking during fatigue of the cross woven C/Sic composite. Similar damage mechanisms. termed geometric stiffening, were proposed in Ref. [5] to explain the fatigue modulus recovery in fatigue survived samples of a cross woven composite. Matrix cracking was observed to increase after the first few cycles as the formation of new matrix damage zones developed inside the fiber bundles (Fig. 7). This continued matrix cracking is associated with bundle splitting and ply delamination. During cyclic loading, the structural adjustments caused by bundle splitting and ply delamination can subject the affected bundle sections to higher stress, as described above. These affected bundle sections are most likely in cross-over sections of longitudinal bundles (Fig. 10 and Fig. 11) where the matrix distribution is often discontinuous. The high stress in matrix discontinuous sections leads to more matrix damage zones during fatigue. The development of these matrix damage zones during fatigue contributed to the gradual decrease of the modulus and the increase of elapsed strain. The continued matrix cracking during tensile fatigue was also observed in many unidirectional-cross ply glass ceramic composites [4,8,9] as well as in a balanced cross woven C/Sic composite [15], although the mechanisms for these composites are different from those in the present composite. For example, stress corrosion of the matrix is proposed for glass ceramic composites [8]: while a high loading rate was found to slow down matrix cracking for a cross woven C/Sic composite [15].
Continued fiber fracture in the fatigue-failed samples is also related to bundle splitting and ply delamination, which subject fibers in the affected bundle sections to a higher stress due to the aforementioned structural adjustments. In addition. the affected bundle sections are often at the bundle crossover (‘weak’ sections) e.g., Fig. 11 in which both fiber fracture and bundle splitting can be seen. Thus. the local stress can be raised to higher levels than the average applied stress, because these sections often have smaller cross section areas and the stress distribution can be quite non-uniform, leading to continued fiber fracture during fatigue. The continued fiber fracture leads not only to modulus decrease and increase in elapsed strain (Fig. 5(d)--(d)), but also to reduction in composite residual strength. Therefore, the plausible fatigue failure mechanism is that, because of the continued fiber fracture, the composite residual strength is gradually reduced to the level of the applied stress, until finally the longitudinal bundle fails in the ‘weak’ bundle sections. Once one bundle fails, it triggers other bundles to fracture due to their having to carry increased load, leading to the composite fatigue fracture. This mechanism is supported by the observation of the fatigue fracture surl&e of the composite which consists of mostly bundle fractures near the crossover poinls, as indicated in Fig. 12(a). Continued fiber fracture can occur as a result of fiber-matrix interface degradation [13] and by attrition of bridging fibers at the fiber-matrix interface as reported 1191, and can be expected to occur here also. even at the low stress of 320 MPa. Such a mechanism of failure could possibly induce failure at stress levels below the fatigue limit reported here. but if operable, was too gradual to be observed. Based on the above discussions, the behavioral differences between the fatigue-Pdiled samples and fatiguesurvived samples can be explained as follows: For the fatigue-Failed samples, the development of all seven damage modes caused a continuous increase in elapsed strain and decrease in modulus till sample failure (Fig. 5(a)-6(c)). For the Fatigue-survived samples, on the other hand, the development of these damage modes except continued fiber fracture led to an initial modest increase in elapsed strain and decrease in modulus {Fig. 5(c) and (d)). For these samples, after a certain number
of cycles, the structural adjustments associated with the increase of ply delamination and bundle splitting become the dominant damage mechanisms for the composite. These structural adjustments, i.e., the straightening of fibers and the realignment of the fiber bundles, can redistribute the load amongst the fiber bundles more evently without breaking the fibers. Therefore, the composite experienced saturation of the elapsed strain and a modulus recovery at high cycle numbers. These mechanisms are also responsible for the increase in residual strength, compared to that of the monotonic tensile strength (Table 1 and Fig. l), because a higher stress can be taken by the realigned and straightened bundles in the final tensile test if the fibers carry it more evenly. A similar increase of the residual strength in a balanced cross-woven C/Sic composite after lo6 cycles was reported in [15]. In that study, the increase in residual strength was proposed to result from the relief of the cross concentration near fiber bundle crossovers by a process of matrix fracture and interfacial debonding. Given the complexity in the damage and fracture mechanisms observed here, such a stress relief mechanism could very well operate in the subject composite. The sample specific fatigue behavior, indicated by the relatively large data scatter in the fatigue test results (Fig. 3 and Fig. 5(a)-(d) and Table l), can be attributed to the microstructural differences in the composite samples, since the testing conditions were the same for all the fatigue samples. This sample specific behavior was not seen in monotonic tensile loading for the same sample size and design, and testing facility [16]. Therefore, it is reasonable to suggest that the sensitivity of the fatigue behavior to the structural variations in the cross woven C/Sic composite is greatly increased by the high loading rate (about 640760 in fatigue versus 3-5 MPa s - ’ in the monotonic tension) and the load repetitions in the fatigue test. A high loading rate in fatigue was found to cause more delaminations in another cross-woven C/Sic composite (X. Wu. and J.W. Holmes, personal communication) and in addition, the repeated loading of the composite was also observed to lead to the growth of the dalamination and bundle splitting in all the fatigue samples in the present study. Therefore, the combined effect of high loading rate and repeated loading produced a large amount of ply delamination and bundle splitting in the cyclic samples, increasing their influence on the fatigue behavior of the C/Sic composite, as described above. Since ply delamination and bundle splitting can most readily occur in areas of structural irregularity, different densities of these structural irregularities can lead to different amounts of ply delamination and bundle splitting. Due to the ‘random-phase lay up’ of the cross woven fabric in the present composite (Fig. 2), different densities of structural irregularities can exist
among fatigue samples. Therefore, different amounts of ply delamination and bundle splitting can be generated for all the fatigue samples, leading to sample specific fatigue behavior, and thus, a large scatter in the experimental results. Similar phenomena were observed in the compression behavior of the same C/Sic composite, in which ply delamination was one of the main damage modes [18].
5. Conclusions Studies of the tension-tension fatigue behavior of a C/Sic composite and its damage and fracture mechanisms lead to the following conclusions: (1) The observed fatigue stress versus life relationship (S-N diagram) showed the usual trend that as the fatigue stress increases, the fatigue life decreases. Runouts without fracture which occurred at IO6 cycles of fatigue are concluded to represent a fatigue limit in the area of 320-340 MPa, above which fatigue failure took place by a number of mechanisms but particularly the the mechanism of continued fiber fractures occurring mostly in the longitudinal bundle crossover sections (weak sections) during fatigue. (2) The fatigue behaviors of the cycled samples were found to result from the continued development of the cyclic damage, including the formation of ply delamination, bundle splitting, matrix damage zones in the fiber bundles, matrix wear, and fiber fracture. For fatiguefailed samples (cycled above the fatigue limit), the development of all the cyclic damage modes caused the elapsed strain to increase and the composite modulus to decrease continuously with cycling until fatigue failure occurred. For the fatigue-survived samples (cycled at fatigue limit), the development of the cyclic damage modes, except for continued fiber fracture, caused the elapsed strain to increase and the composite modulus to decrease initially. At high cycles up to lo6 cycles, structural adjustments associated with ply delamination and bundle splitting produced modulus recovery and even high residual strength in fatigue-survived samples. This behavior was caused by more even load distribution in the composite without fiber fracture. (3) The high loading rate and repeated loading in fatigue are believed to increase the sensitivity of the fatigue behavior to microstructural inhomogeneities existing in the C/Sic composite, by producing more ply delamination and bundle splitting. Different densities of structural irregularities, at which ply delamination and bundle splitting take place, exist in different samples due to the ‘random-phase’ lay up of the fabric in the composite, causing the sample specific nature of the fatigue behavior in the C/Sic composite.
IS2
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?Vutzg, C. Laird,/Materials
Science
Acknowledgements
This work was supported by both the Benjamin Franklin Program of the Commonwealth of Pennsylvania and the Materials Science Corporation. The authors are deeply grateful for this support. The authors also thank Dr Alex Radin for his expert help in mechanical testing of the composite samples, and MS Xuqing Wang for her help in characterization studies. The Laboratory for Research on the Structure of Matter of the University of Pennsylvania both supported the work and provided experimental facilities. The authors acknowledge as well a crucial equipment grant from the Research Foundation of the University of Pennsylvania.
:l] K.M. Prewo, Fatigue and stress rupture of silicon carbide fiberreinforced glass ceramics, J. Marer. Sci., 22 (1987) 2695-2701. 121E. Minford and K.M. Prewo, in CC. Patano et al. (eds.), Fatigue behavior of silicon carbide fiber reinforced lithium-alumino-silicate glass-ceramics, Tailoring Mzdfiphase and Composite Ceramic, Plenum, New York, 1986, pp. 561-570. 131D. Lewis III, CycIic mechanical fatigue in ceramic-ceramic composites-an update. Ceratn. Eng. Sic Proc., 77-S) (1986) 874-881. ‘41 C.Q. Rousseau, in J.M. Kennedy, H.H. Moeller and W.S. Johnson (eds.), Monotonic and cyclic behavior of a silicon carbidecalcium-aluminosilicate ceramic composite. Thermal and Mechanical Behaoior of Metal Morrix and Ceramic k1atri.v Composites. ASTM STP 1080, American Society of Testing and Materials, Philadelphia, 1990, pp. 136- 151. :5] R.F. Allen and P. Bowen. Fatigue and fracture of a Sic-CAS continuous fiber reinforced glass ceramic matrix composite at elevated temperature, Ceram. Sci. Eng. Proc., 9- 10 (1993) 265272. 141J.W. Holmes, T. Kotil and W.T. Foulds, High temperature fatigue of Sic fiber-reinforced Si, N, ceramic composites, in Symposiutn on High Temperature Composites, Technomic, Lancaster, PA, 1989, pp. 176-182.
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I71 L.P. Zawada, L.M. Butkus and G.A. Hartman, Tensile and fatigue behavior of silicon carbide fiber-reinforced aluminosilicate glass, J. Am. Ceram. Sot.. 74fll) (1991) 2851-2858. PI P.G. Karandikar and T.-W. Chou, Damage development and moduli reductions in nicalon-calcium aluminosilicate composites under static fatigue and cyclic fatigue, J. Am. &ram. Sot., 76(7j (1993) 1720- 1728. [9] F.A. Habib, R.A.J. Taylor, R.G. Cooke and B. Harris. Fatigue damage in Sic-CAS composites, Composiles, N(2) (1993) 157165. [lOI A.W. Pryce and P.A. Smith, Matrix cracking in cross-ply ceramic matrix composites under quasi-static and cyclic loading, Acta Metall. ,Ifater.. 42(3j (1994) 861-870. s11 T. Kotil, J.W. Holmes and M. Comniou, Origin of hysteresis observed during fatigue of ceramic-matrix composites, J. Am. Ceratn. Sot. 73(7) (1990) 1879-1883. [I21 C. Cho, J.W. Holmes and J.R. Barber, Estimation of interfacial shear in ceramic composites from frictional heating measurements, J. Am. Ceram. Sot., 74(11) (1991) 2802-2808. [I31 A.G. Evans, F.W. Zok and R.M. McMeeking, Fatigue of ceramic matrix composites, Acta Metali. niiaier., 43{3] (1995) 859-874. [I41 2. Wang, C. Laird, 2. Hashin, B.W. Rosen and C.-F. Yen, Mechanical behavior of a cross-woven ceramic composite, part II: repeated loading, J. Mater. Sk, 26 (1991) 5335-5341. L151 S.F. Shuler, J.W. Holmes, X. Wu and D. Roach, Influence of loading frequency on the room-temperature fatigue of a carbonfiber-sic-matrix composite, J. Am. Ceram. Sot., 76<9,) (1993) 2321-2336. [I61 M. Wang and C. Laird, Characterization of microstructure and tensile behavior of a cross-woven C/Sic composite, Aura Melallwgica et Mareriaiia, 44 (1996) 1371-1388. M. Wang, Damage and fracture in monotonic and cyclic!loading of cross-woven C/Sic composite, Ph.D. Thesis, University of Pennsylvania, Philadelphia, 1994. M. Wang and C. Laird. Damage and fracture of a cross woven C,/SiC composite subject to compression loading, J. Mar, Sci., 31 (1996) 2065-2069. W.L. Morris, B.N. Cox, D.B. Marshall, R.V. Inman and M.R. James, Fatigue mechanisms in graphite-sic composites at room and high temperature, J. Am. Ceram. Sot., 77(3) (1994) 792800.
ELSEVIER
Materials Scienceand Engineering A230 (1997) 171- 182
Tension-tension
fatigue of a cross-woven C/Sic composite Mingde Wang, Campbell
Department
of M&vkds
Science
md Etlgiueering,
Unicersit~~
Laird *
of Penng~lrn~~in,
Philadelphia,
PA 191046272,
C;S’A
Received2 November1996;receivedin revisedform 12November1996
Abstract The tension-tension fatigue behavior of a crosswoven C/Sic compositewas studiedin termsof damagemodesand damage development.The fatigue stressversuslife diagram(S-N) curve and an endurancelimit of 320-340 MPa (about SO’% tensileUTS) for IO6cycleswereobtained for the C/Sic composite.Different fatigue behaviorswerefound for samplesthat failed during fatigue and for samplesthat survived 10” cycles. Sevenfatigue damagemodeswere observed,the developmentof which were usedto explain the different fatigue behaviors.For the fatigue-failed samples,the degreesof damageof the sevenmodesincreasedwith increaseof cycles,leadingto an increasein elapsedstrain and a decreasein compositemodulus.For fatigue-survivedsamples,the developmentof all the damagemodesexcept for fiber breaking causedan initial increaseof elapsedstrain and decreaseof compositemodulus,but at high cycles,fiber bundle realignmentand straighteningin thesesamplesled to partial recovery of the modulusand cessationof the damagedevelopment.0 1997Elsevier ScienceS.A. Kqwords:
C/Sic composite; Fatigue;Fiberbundle
1. Introduction
but also to predict the cyclic stress-strain loops [lo? 111, and to estimate interFacial friction by C. Cho et al. [12].
Fiber reinforced ceramic matrix composites (CMCs) have been under extensive development as potential high temperature materials because of their possible high strength and toughness at both room and high temperatures. Since the applications of these composites are likely to include jet engines and heat engines, the cyclic fatigue behavior of these composites has been the focus of a number of studies, especially in composite systems with unidirectional or cross-ply fiber arrangements such as Sic fiber (Nicalon)/lithium-aluminum-silicate (LAS) composite [l-7]. In those unidirectional or cross-ply ceramic composites, there seems to exist a definite fatigue limit, which approximately coincides with the proportional limit of the 0” plies in monotonic loading. In those composite systems, the most studied fatigue damage mode has been the same as in monotonic loading, i.e., matrix cracking. Because of the simple fiber arrangement in the unidirectionalcross-ply composites, it has been possible not only to relate matrix cracking during fatigue to the reduction of composite stiffness through mechanical modeling [8,9],
The fatigue failure mechanisms in the fiber reinforced CMCs are associated with fiber-matrix interface degra-
*Correspondingauthor. Tel.: + 1 215 8986664;fax: + 1 215 5732128; e-mail:[email protected] 0921-5093/97/$17.00 0 1997ElsevierScienceS.A. All rightsrcsrrved. HISO921-5093(97)00018-X
dation and degradation of fiber strength. A recent overview paper by A.G. Evans et al. [13] deals with interface degradation mechanism, giving a detailed
analysis and modeling of the fatigue behavior of unidirectional and cross-ply CMCs. In ceramic composites with fiber reinforcements in the form of textile preforms such as cross woven C/Sic composites, cyclic damage is a more complex process. For example; in an exploratory study of a cross woven C/Sic composite by Wang et al. [14], ply delamination and matrix wear between plies and fiber bundles were observed and found to be important in fatigue failure
of this composite, in addition to matrix cracking. These damage modes were also found to be more gradual and localized during fatigue than in monotonic loading. Such fatigue damage modes were identified in a study on a similar type of cross woven C/Sic composite by Shuler et al. [15]: and the mechanisms for these damage modes were proposed to result from non-uniform stress-strain distributions produced by the cross weaving of the fiber bundles in the composite. As in unidirectional CMCs, the modulus reduction was related
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largely to the matrix cracking in that study, and other fatigue damage modes were observed from the fracture surfaces. Damage development under cyclic loading was demonstrated as it developed. In a recent study on a slightly different cross woven C/Sic composite by Wang et al. 1161, six damage modes were identified during monotonic tension due to the complex composite microstructures, and they were found to be responsible for the mostly nonlinear stress-strain behavior of the composite, which is shown in Fig. 1. As an extension to the monotonic tensile behavior, the purpose of the present study is to investigate how the complex damage develops under cyclic loading conditions and the effect of the cyclic damage on the fatigue behavior of a typical textile ceramic composite, In the present study, the fatigue stress versus life relation was first obtained for an unbalanced cross woven C/Sic composite. The fatigue behavior of the composite, represented by using elapsed strain and modulus, was found to be different for fatigue stresses above the fatigue limit and at the fatigue limit. Fatigue damage and damage evolution were examined and used to explain the differences in the fatigue behavior at stresses above and below the fatigue limit.
2. Materials and experimental procedures
The experimental composite used in this study is an unbalanced cross woven C/Sic composite supplied by DuPont. The microstructure of the composite was characterized in detail previously [16]. The interested reader is referred to that paper for details of the material. Here? only those microstructural features relevant to the present study are briefly summarized as follows.
Fatigued
T g
350 300 250
< Virgin
-I :: 2
200
G
150
./“’
Sampl
Sample
..J’
50 t :, 0 f 0
I'/ I 0.2
0.4 strain
0.6
0.6
(%f
Fig. 1. The stress--strain curve of a typical virgin C/Sic composite. together with the stress-strain curve of a fatigued survived sample. The off-set strain in the strain axis for the latter is equal to the residual strain accumulated through fatigue.
and
Etq@wit~gA230(1997)I71-18-3
Fig. 2. Longitudinal cross section of C&C composite, showing coating layers on both sample surfaces and interbundle voids (A and B). Bundles X and Y were laid up mostly through contacts between transverse bundles while Y and Z have extensive degrees of longitudinal contacts. Notice that the travel of bundle C ends at D, indicating bundle misalignment with respect to the polished plane. E and F indicate bundle sections with differeni matrix distributions.
The typical microstructure of the C/Sic composite can be seen in a longitudinal cross section: as shown in Fig. 2. Several features can be seen from this figure. First, the undulations of the longitudinal bundles in different plies are not in synchrony with each other (e.g., the first and second bundles from the top surface in Fig. 1). a situation termed ‘random phase lay up’, Second, as a consequence of the random phase lay up, the contacts between the plies in the composite are highly irregular, i.e., some plies contact through transverse bundles iX and Y in Fig. l)? and other plies are laid up through various degrees of contact between longitudinal bundles (Y and Z in Fig. 2). Third, this random phase lay up also caused a non-uniform distribution and size variation in the interbundle voids (e.g., A and 3 in Fig. 2-white areas). Most of these voids are oriented in the longitudinal direction (e.g., A in Fig. 2)> and their sharp tips can act as sites for local ply delamination under cyclic stresses. Lastly, a less obvious feature in the composite is the misalignment of some plies (and bundles in those plies) with respect to the polished plane of the sample (the loading direction in the fatigue tests was parallel to this plane), as evidenced in Fig. 2 by the discontinuation of the travel of the bundle marked C from right to left at point D. In fact, all the plies have a certain degree of misalignment due to the random phase lay-up of the fabrics. The longitudinal bundles in the composite consist of three typical kinds of sections with different matrix distributions, as described in detail 1161. The first section, which exists between crossover of the bundle and is not in contact of the adjunct plies (e.g., Section E in
Fig. 2), has a continuous matrix distribution, and is thus termed ‘matrix-continuous’ section. The second cross section, which is often positioned in the crossover of the fabric and is also pressed by the adjunct ply (e.g., Section F in Fig. 2), has a discontinuous matrix distribution and thus, is termed ‘matrix-discontinuous’ section. In this section, the fibers are densely packed (i.e., mostly fiber-fiber contact) and the matrix exists as isolated ‘matrix pockets’, filling the interstitials of the dense fiber pack. The third section is a ‘mixed section’ consisting of the above two, i.e., one part of the section has continuous matrix distribution and the other has discontinuous matrix distribution. Because of the different matrix distributions in a fiber bundle with the same number of 3000 fibers, the cross section area of the ‘matrix-continuous’ section is larger than that of the ‘matrix-discontinuous’ section (e.g., see Sections E and F in Fig. 2). Therefore, the present composite possesses a high degree of variability in microstructure at both composite level and fiber bundle level, which can lead to complex damage modes such as ply delamination. Because of the complex microstructures of the composite, the volume fractions of the fiber, matrix and void can only be approximated (see Ph.D. Thesis by M. Wang [17]) and are about 53, 34 and 13%, respectively. In addition to its complex microstructure, the composite is coated with thick Sic coatings on both surfaces of the composite plate (Fig. 2) and thin SIC coatings over fiber bundle sections which are open to the interbundle voids (e.g., Section E in Fig. 2). Matrix cracks were found in these coatings and in the transverse fiber bundles in the as-received composite [16,17]. Since these cracks resulted from the CVI processing, they are termed processing-induced cracks. The crack populations can be quantified as ‘surface coating crack density’ and ‘bundle coating crack density’. The measurements of such densities, which were also described in [17], are represented by numbers of cracks per unit length (no. mn- ‘). These crack densities were obtained both for as-received samples and for samples fatigued at all the stress levels conducted in this study.
For cyclic fatigue tests, straight-sided composite specimens 65 mm long, 4 mm wide and 4.4 mm thick were used. Both ends of the specimen were tabbed using clear polyester resin for gripping. For tension-tension fatigue tests with R = 0, in which tension was applied in the ‘longitudinal’ direction, an Instron 1361 screwdriven universal testing machine was used, which was equipped with a MTS hydraulic grip system. Strains were measured using an Instron dynamic extensometer with a gage length of 10 mm. To avoid extensometer slippage during a test, the extensometer knife edges
were glued to the specimen using a 5 min epoxy. The load and uniaxial displacement during a fatigue test were recorded using a X- Y recorder and a data acquisition system running in an IBM PC. Pulsating tension fatigue tests were conducted in load control using a half sinusoidal wave at 1 Hz. Based on the monotonic strength of 405 MPa for the experimental composite (see Fig. l), the controlled load levels were chosen to be 380, 360, 340 and 320 MPa. A typical test was run continuously until the sample failed (i.e., fractured in the gage section, otherwise the result was considered invalid) or survived 10” cycles. The fatigue life of a failed sample was counted to the last cycle number at which the sample fractured. For the fatigue-survived samples, a tensile test was finally performed at a loading rate of about 5 MPa s- ’ (i.e., the same as used in monotonic tests) to determine the residual strength. During fatigue tests, individual stress-strain loops were recorded for the following cycles: l> 2, 4, lo> 20, 10, lo’, 2 x lo’, 4 x lo’, lo”, 2 x lo”, 4 x lo’, lo”, 2 x lo’, 4 x lOa, 105, 2 x lo”, 4 x 10’ and 10h. From these loops, the elapsed strains, which are defined as the accumulative peak strain at the above selected cycles, can be obtained. The composite modulus can also be obtained by linear fitting of the loading portion of the stress-strain loops except for the first cycle. The cycle intervals chosen were found to describe the fatigue behavior in a representative f’ashion. In addition, the strain versus time during fatigue was recorded continuously on a chart recorder. From this recording, the elapsed strain for the last cycle before fatigue failure was obtained for comparison among the fatigued samples and with the tensile failure strain. In order to understand the fatigue behavior of the composite, fatigue damage evolution was monitored at two stress levels, i.e., 360 and 320 MPa, which are above and at the fatigue (or endurance) limit respectively. Side-polished samples were examined and areas of interest were selected in the SEM before loading. Subsequently, the same areas of interest were observed (1) at cycles 1, 10 and 10’ for a sample fatigued at 360 MPa; (2) at cycles 1, lo”, 10” and lo6 for the sample fatigued at 320 MPa. The fracture surfaces of fatiguefailed samples were examined using SEM and the results for a typical fracture are also reported here.
3. Tension-tension composite
fatigue behavior of C/SC
The fatigue stress versus life diagram (S-N diagram), obtained from a total of 12 specimens, is shown in Fig.
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Science and Engineering A230 (1997) I71- 182
3. From this figure, it is seen that a fatigue limit’ (or more accurately endurance limit) for lo6 cycles exists at stresses 320-340 MPa for the subject C/Sic composite. For present purposes, we define this stress as the IO6 cycle endurance limit. Specifically, the two samples fatigued at 320 MPa plus one sample fatigued at 340 MPa survived lOh cycles. These three samples are referred to as ‘fatigue-survived’ samples. All the other samples, which failed before lo6 cycles, are referred to as ‘fatigue-failed’ samples. It is also seen from Fig. 3 that, in general, the average fatigue life decreases as the fatigue stress increases. However, and as usual, in the S-N diagram, there is some degree of scatter, which indicates the sensitivity of the fatigue life to the microstructure variability in the C/Sic composite samples.
3
250
52 -
200
I 2
tj
150
0
O.‘l
0.2
0.3
Strain
3.2, Cyclic stress-stmin behavior
In order to understand the fatigue behavior of the C/Sic composite in detail, it is necessary to study the hysteresis stress-strain loops of the fatigued samples. Fig. 4 shows the typical hysterisis stress’strain loops of a fatigued sample at selected cycles. From Fig. 4, it is seen that the loading portion of the first cycle is largely non-linear, similar in shape to that of a monotonic tensile test (Fig. 1). The first cycle shows an open loop with large residual (irreversible) strain when the sample was unloaded. In the subsequent cycles, the stress-strain
4201
32-I
*
-
IO
0
10 1
10
2
10
3
10
4
10
5
10
Q-80%
6
10
7
LIFE (Cycles) Fig. 3. S-N curve (fatigue stress vs, life curve) for C/Sic composites fatigued at 1 Hz. The solid lines indicate the possible range of experimental data, and show the trend of life behavior. The point at 10° cycles represents the monotonic strength.
’ The solid lines in this figure were drawn to indicate the trend of the relation between fatigue stress and life, and are based on the scatter of the experimental data, as well as an assessment of fatigue scatter in many material systems. Clearly, -there are insufficient specimens to establish high cycle endurance behavior on a proper statistical basis.
0.4
0.5
0.6
0.7
(%)
Fig. 4. Stress-strain loops at selected cycles for a fatigue stress of 360 MPa (sample F-4). This behavior is typical of the stress-strain loops for all the fatigue samples. Definition of elapsed strain is indicated in this figure.
loops appear closed. In contrast to the first loading, the typical loading curves of the subsequent cycles are mostly ‘linear’ except for the very beginning. It is also seen from Fig. 4 that fatigue loading produces cyclic creep* in the composite (which is indicated by rachetting of the loop peaks). In the meantime, the loops tend to tilt downwards as the cycle count increases. Therefore, it is useful to choose the elapsed strain to describe the cyclic creep, and the composite modulus to describe cyclic damage for the fatigued samples because the former represents the dimensional changes, and the latter is both an indication of damage as well as an important design parameter for practical engineering applications. In order to compare the changes of the modulus among the fatigue samples, the modulus was normalized with respect to the average tensile modulus of 137 GPa and then used in the modulus versus cycles plots (Fig. 5(a)-(d)). This approach was adopted because the tensile modulus is characteristic of the C/Sic composite [16,18], and because the limited linear part of the first fatigue loading curve did not contain enough data to obtain a reliable value of the modulus for the individual specimens. 3.3. Variations of elapsedstrai~z and modztlzts lrlith cycling
Fig. 5(a)-(d) show the elapsed strains and normalized composite moduli versus cycles for samples fatigued at 380, 3609 340 and 320 MPa, respectively. From Fig. 5(a)-(d), it can be seen that, for the fktiguefailed samples, in general, the elapsed strains increased continuously from cycle one, while the normalized ’ We employ this term as conventionally used for metal specimens
M.
1.1
Wang,
C. Laird
/ Marevials
,
Science
and
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modulus decreased steadily after a large drop at cycle one. For the fatigue-survived samples (two at 320 MPa in Fig. 5(d) and one at 340 MPa (F17) in Fig. 5(c)), the elapsed strain increased only modestly in the beginning, and then, stabilized at high cycles (lo4 for F17 and 340 MPa and lo5 at 320 MPa. Also note that in Fig. 4(c), the elapsed strain for samples F17 and F20 appear to decrease after 10” and 10’ cycles, and this is considered to be experimental error caused by possible slippage of the extensometer during fatigue). For these samples, the relationship between the normalized modulus and fatigue cycles is very interesting, as indicated for sample F17 in Fig. 5(c) and the two samples in Fig. 5(d), i.e., the modulus decreased rather modestly after the big drop in the first cycle, but subsequently the modulus recovered at high cycles ( > lo5 cycles).
It is interesting to note that the elapsed strains of all the fatigue-failed samples at one cycle before failure seem to approach a value of about 0.7%, regardless of fatigue stress levels and fatigue life (Fig. 5(a)-(c)). Since this critical elapsed strain of 0.7% for fatigue failure is comparable to the average failure strain of the composite under monotonic tensile loading (which is about 0.69%, see Fig. l), it is reasonable to suggest that the failure of the C/Sic composite is strain-controlled, i.e., no matter what the loading condition is, once the failure strain is reached, the composite will fail. On the other hand, for the fatigue-survived samples, since the elapsed failure strain of about 0.7% was not reached during fatigue, they survived the fatigue test for lo6 cycles. Therefore, the elapsed strain is a useful parameter which can be used to describe cyclic creep and to
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monitor whether the composite will fail under other testing conditions such as creep. Fig. 5(a)-(d) also show that there exists a relatively large data scatter in the change of both elapsed strain and normalized modulus with cycles at the same fatigue stress. Since all the fatigue tests were carefully conducted using the same sample size and test procedures, this large data scatter can only indicate that the fatigue behavior of CiSiC composite is sample specific, i.e., the fatigue behavior is very sensitive to microstructurally-induced differences that may exist from sample to sample. 3.4. Damage and fktwe
during cyclic iondhg of
C/Sic composites
In order to understand the different fatigue behaviors of the C/Sic composite described above. damage evolution and fracture during tension-tension fatigue were examined at two stress levels. The first stress level was 320 MPa, representing fatigue-survived samples, and the second, 360 MPa. representing fatigue-failed samples. In general, a total of seven fatigue damage modes were observed in both samples. These damage modes are: (1) matrix cracking; (2) transverse bundle cracking; (3) interfacial debonding; (4) fiber breaking; (5) ply delamination; (6) bundle splitting and (7) matrix wear. All the damage modes except matrix wear (item 7) were observed after the first fatigue cycle at both stress levels. Since the six damage modes produced in the first cycle are the same as those observed in monotonic tensile tests r16], they are not shown here in detail, instead they will be pointed out in the following discussion. After the first cycle, all the damage modes accumulate in a subtle manner: leading to a gradual increase in elapsed strain and decrease in modulus, as indicated in Fig. 5(a)-(d). Matrix cracking, observed during fatigue at the two stress levels, includes matrix cracks in the longitudinal bundles and in the composite surface coating. Since the longitudinal bundles are the major load carrying elements in the composite for the fatigue tests, matrix cracking behavior was examined mainly in these bundles. As described in Section 2.1, the longitudinal bundles are composed of ‘matrix-continuous’ sections. ‘matrix-discontinuous’ sections and ‘mixed’ sections, thus, matrix cracking in these typical sections was monitored during fatigue. In the matrix-continuous section, matrix cracks run across most of the bundle in a mostly continuous and planar manner as shown in Fig. 6. In the matrix-discontinuous section, matrix cracking occurs in the form of matrix damage zones, which consist of small separated matrix cracks inside the bundle, as shown in Fig. 7(a). The matrix cracking behavior in the ‘mixed’ section is a mixture of the above two mechanisms 1171. These types of matrix cracking formed in the first cycle, which are the same as those observed in the monotonic tensile test [16]. With in-
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crease of cycles, however, matrix cracking continues to occur mainly in terms of new matrix damage zones in the ‘matrix-discontinuous! sections and ‘mixed’ sections of the longitudinal bundles, as shown in Fig. 7(b). Matrix crack multiplication was found to play an important role in the decrease of the composite modulus in the tensile test of the present composite [16], and in the fatigue of ceramic composites including unidirectional Nicaloniglass ceramics [8,9], and across woven composite [15]. To determine whether a similar relationship exists in the fatigue of the present composite, the densities of bundle coating cracks and surface coating cracks were measured as a function of cycles for stress levels of 380. 360. 340 and 320 MPa. The results are plotted in Fig. 8, from which it is seen that the crack densities for the longitudinal bundle coating cracks and the sample surface coating cracks saturate in the first two cycles with the biggest increase in the first cycle. The results in Fig. 8 also confirm the observation that the development of matrix cracking after the first two cycles took place inside the longitudinal bundles, i.e., in terms of matrix damage zone formation (e.g., Fig. 7), since the formation of matrix damage zones is not reflected in the bundle coating crack densities. Fig. 8 also shows a large data scatter? indicating again that the fatigue behavior of the composite is sample specific. Longitudinal bundle splitting, which was seen in monotonic tensile loading to the high stress level of 385 MPa [16], was observed at both fatigue stress levels of 320 and 360 MPa. This type of damage was found to occur mostly in the crossover sections of the longitudinal bundle. A typical example is shown in Fig. 9(a) for a sample fatigued for 1000 cycles at 320 MPa, With cycling, the splitting was seen to grow (Fig. 9(a)). Also in Fig. 9(b), matrix wear (in terms of matrix fragmentation) is seen to accompany the longitudinal bundle splitting.
Fig. 6. Matrix cracking in a mostly ‘matrix continuous’ section of a longitudinal bundle. The matrix cracks are continuous and straight across the bundle. The loading direction is horizontal.
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The fracture surface of a typical fatigue-failed sample is shown in Fig. 12(a), which indicates that failure of the composite consisted of fractures of the longitudinal bundles. From Fig. 12(a), it is seen that most of the fractures of the longitudinal bundles appeared to occur at or near the crossover sections of the bundle (e.g., A and B) with few bundles fracturing between crossover sections (e.g., C). Examination of the longitudinal bundle fracture surfaces which occurred at or near the crossover sections of the bundle revealed a ragged fracture surface with limited fiber pull-outs, while those which occurred between the crossovers showed either fracture surfaces with long fiber pull-outs or flat fracture surfaces with little fiber pull-out. Since these types of bundle fracture surfaces are the same as those observed in monotonic tensile fracture [16], they are now shown in detail here. However, in the flat fracture surface of a longitudinal bundle (Fig. 1 l(b)), an interesting feature of the matrix wear, which was closely associated with fatigue loading, was observed. This appears as the rough area on the left of the fractograph, indicated by W. 3.5. Residuul strer?gth
(b) Fig. 7. Matrix cracking in a ‘matrix-discontinuous’ sections of a longitudinal bundle at 320 MPa: (a) after 1 cycle, two matrix damage zones formed (N); (b) after 1000 cycles, another matrix damage zone developed (M).
Ply delamination, which was seen in monotonic tensile loading to the high stress level of 385 MPa [16], was observed at both fatigue stress levels of 320 and 360 MPa. An example of this damage mode is shown in Fig. 10, which was taken after 1 cycle at 320 MPa. Ply delamination was also observed to grow with increase of fatigue cycles for both stress levels of 320 and 360 MPa. Fiber breaking was observed after the first cycle in both samples fatigued at 320 and 360 MPa. However, only the higher stressed sample caused continued fiber fracture with increase of cycles. Again, like bundle splitting, the continued fiber fracture was seen to occur mostly in the crossover section of the longitudinal bundles. A typical example is shown in Fig. 1l(a) and (b), in which continued fiber fractures were seen with increase of cycles, together with the growth of bundle splitting and matrix wear in the sample fatigued at the stress of 360 MPa.
The residual strength of the fatigue-survived samples (two at 320 and one at 340 MPa) is listed in Table 1, together with the accumulative failure strains (i.e., the sum of the residual strain at lo6 cycles and the subsequent tensile failure strain). The tensile stress-strain curve of a typical fatigue-survived sample is shown in Fig. 1. From Table 1 and Fig. 1, it is interesting to see that the residual strength of these fatigued samples increased by an amount ranging from about 10% (F9) to 20% (F15), and the proportional limit of the fatigue sample is higher than that of the virgin sample. Again, the residual strength is sample specific. However, if we 6
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380MPa
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360 MPa (90% of T . S.)
(95% of T . S.)
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340 MPa (85% of T . S.) 320 MPa (80% of T . S.)
2 0
20
40
60
80
100
Cycles Fig. 8. Crack densities vs. cycles for the four fatigue stress levels studied. The crack densities were measured on composite surface coating and longitudinal bundle coating. The horizontal arrows positioned at 100 cycles indicate that the samples did not fail at the maximum cycle plotted in this figure.
Fig. 9. Growth of longitudinal bundle splitting at 320 MPa: (a) after 10” cycle, longitudinal splitting appeared (S) and N indicates matrix cracking: (b) after 10” cycles, the splitting grew with the associated production of matrix fragments.
The present C/Sic composite exhibits seven fatigue damage modes at all the stress levels studied here. These damage modes include: (1) matrix cracking, (2) transverse bundle cracking, (3) mterf’acial debonding, (4) fiber breaking, (5) ply delamination, (6) longitudinal bundle splitting, and (7) matrix wear. During the first fatigue cycle, all these modes of damage except for matrix wear (item 7) were observed to take place. For example, the crack densities of the sample surface coating and the longitudinal bundle coating (Fig. 8), which indicate matrix crack multiplication, increased to a level close to saturation soon after the first cycle for all the fatigue stress levels. The damage and damage accumulation. therefore, are believed to have caused the largely nonlinear stress-strain behavior of the first cycle (Fig. 4) leading to a large reduction of the composite modulus (Fig. 5(a)-(d)), similar to behavior during monotonic tensile loading (Fig. 1). Subsequent cyclic loading of the composite caused gradual damage developments in the composite, leading to nearly closed stress-strain loops and gradual changes in modulus and elapsed strain (Fig. 4 and Fig. 5(a)-(d)). The damage developments observed in this study include increases of bundle splitting, ply delamination and matrix cracking. continued fiber breaking, and matrix wear. Of these damage developments, bundle splitting and ply delamination are the unique and important damage mechanisms in the present composite. which influence and are enhanced by other damage mechanisms. Longitudinal bundle splitting is a damage mode that was observed 1.0 increase during fatigue (Fig. 9(a) and (b), and Fig. IO). Longitudinal bundle splitting is a result (also an .indication) of the ‘permanent’ straightening of longitudinal bundles. The mechanism for this is closely related to the cross weave of the C/Sic com-
consider the accumulative failure strain for the fatigue survived samples as the ‘failure strain’. the values of this ‘failure strain’ are about the critical value of 0.7%. comparable to monotonic tensile failure strain and fatigue elapsed failure strain of the CjSiC composite (see Section 3.3), confirming that the fracture of the C/Sic composite is controlled by the total strain. 4. Discussion
The fatigue behaviors of the C/Sic composite reported in the previous section can be understood qualitatively in terms of damage and damage development occurring in the composite during cyclic loading. Through this understanding, the failure behaviors can be related to the complex microstructure of the C/Sic composite. Thus. we point out key microstructural features responsible for important fatigue properties, such as the fatigue limit.
Fig. 10, Ply deiam:nations (II) between longitudinal bundles in C/Sic composites after 1 cycle at 320 MPa.
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bundle shear strength, both of which make it easier for the bundle to split. Like bundle splitting, ply delamination was also seen after the first loading cycle in the composite (Fig. 9). Ply delamination can be caused by either the shear stress generated from unequal displacement between contacting bundles in neighboring plies, or tensile stress normal to the longitudinal bundle due to the straightening of undulated bundles. These two mechanisms can take place in the present cross woven C/Sic composite in which the bundle contacts are highly irregular (Fig. 2) so that both shear stress and normal stress can be generated in these irregularly contacted areas, and both stresses can also be increased by the interbundle voids which have sharp tips oriented in the bundle direction, as shown in Fig. 2. The occurrence of ply delamination and bundle splitting decreases the composite integrity, thus reducing its
Fig. Il. Continued fiber breaking during fatigue at 360 MPa: (a) after 1 cycle, fiber breaking (B), debonding (D) and new cracks (N) are formed; (b) 100 cycles, other fiber breaking (B), matrix wear (W). Bundle splitting growth is indicated by S in (a) and (b).
posite as follows: Under applied stress, the undulation of the longitudinal bundles causes complex stresses which include tension and shear inside the bundle, especially in bundle sections that crossover transverse bundles. In these bundle sections, the fibers most likely contact each other and the matrix is discontinuous, as described in Section 2. l., so that the shear strength is very low. Therefore, when the shear stress caused by bundle undulation under the applied tensile stress exceeds this low shear strength, fracture occurs in the fiber direction, causing the bundle section to split. The bundle splitting damage allows the undulated section to straighten ‘permanently’. With cycling, the stress state in the bundle can change due to damage such as ply delamination and matrix wear, resulting in the growth of the splitting. For example, bundle splitting can be enhanced by ply dalamination which relieves the constraints of the bundles, while matrix wear decreases
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@I Fig. 12. (a) Fracture surface of a sample fatigued at 360; (b) Flat fracture surface of a longitudinal bundle with little fiber pull-out looking down the fiber direction (b). Notice the matrix wear at the bundle fracture surface.
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Table 1 Residual strength of the fatigue survivedsamples Sample no.
Fatigue stress (MPa)
Residual strength(MPa)
Residual-tensile
Failure strain (‘X0)
Tensile test F-9 F-15 F-17
N.A. 320 320 340
405 466 491 453
1 l.l5 1.21 1.1’
0.70 0.67” 0.73” 0.69’
“Failure strain for those samples = the sum of the tensile Failure strain obtained in the final tensile test and the residual strain at 10’ cycles.
modulus and increasing its extension. However, their influences on the fatigue behavior of the present CjSiC composite lie in the associated structural adjustments of the longitudinal bundles. For example, the ‘permanent’ straightening of undulated fiber bundles, which is associated with bundle splitting, can redistribute the load amongst the fibers in the bundle. And ply delamination can relieve constraint in bundles that are misaligned with respect to the loading direction (Fig. 2), permitting them to realign locally to the stress axis, and thus allowing the bundle section to bear higher load. These structural adjustments and the associated stress redistribution are the main causes for the increase of matrix cracking and continued fiber breaking during fatigue of the cross woven C/Sic composite. Similar damage mechanisms. termed geometric stiffening, were proposed in Ref. [5] to explain the fatigue modulus recovery in fatigue survived samples of a cross woven composite. Matrix cracking was observed to increase after the first few cycles as the formation of new matrix damage zones developed inside the fiber bundles (Fig. 7). This continued matrix cracking is associated with bundle splitting and ply delamination. During cyclic loading, the structural adjustments caused by bundle splitting and ply delamination can subject the affected bundle sections to higher stress, as described above. These affected bundle sections are most likely in cross-over sections of longitudinal bundles (Fig. 10 and Fig. 11) where the matrix distribution is often discontinuous. The high stress in matrix discontinuous sections leads to more matrix damage zones during fatigue. The development of these matrix damage zones during fatigue contributed to the gradual decrease of the modulus and the increase of elapsed strain. The continued matrix cracking during tensile fatigue was also observed in many unidirectional-cross ply glass ceramic composites [4,8,9] as well as in a balanced cross woven C/Sic composite [15], although the mechanisms for these composites are different from those in the present composite. For example, stress corrosion of the matrix is proposed for glass ceramic composites [8]: while a high loading rate was found to slow down matrix cracking for a cross woven C/Sic composite [15].
Continued fiber fracture in the fatigue-failed samples is also related to bundle splitting and ply delamination, which subject fibers in the affected bundle sections to a higher stress due to the aforementioned structural adjustments. In addition. the affected bundle sections are often at the bundle crossover (‘weak’ sections) e.g., Fig. 11 in which both fiber fracture and bundle splitting can be seen. Thus. the local stress can be raised to higher levels than the average applied stress, because these sections often have smaller cross section areas and the stress distribution can be quite non-uniform, leading to continued fiber fracture during fatigue. The continued fiber fracture leads not only to modulus decrease and increase in elapsed strain (Fig. 5(d)--(d)), but also to reduction in composite residual strength. Therefore, the plausible fatigue failure mechanism is that, because of the continued fiber fracture, the composite residual strength is gradually reduced to the level of the applied stress, until finally the longitudinal bundle fails in the ‘weak’ bundle sections. Once one bundle fails, it triggers other bundles to fracture due to their having to carry increased load, leading to the composite fatigue fracture. This mechanism is supported by the observation of the fatigue fracture surl&e of the composite which consists of mostly bundle fractures near the crossover poinls, as indicated in Fig. 12(a). Continued fiber fracture can occur as a result of fiber-matrix interface degradation [13] and by attrition of bridging fibers at the fiber-matrix interface as reported 1191, and can be expected to occur here also. even at the low stress of 320 MPa. Such a mechanism of failure could possibly induce failure at stress levels below the fatigue limit reported here. but if operable, was too gradual to be observed. Based on the above discussions, the behavioral differences between the fatigue-Pdiled samples and fatiguesurvived samples can be explained as follows: For the fatigue-Failed samples, the development of all seven damage modes caused a continuous increase in elapsed strain and decrease in modulus till sample failure (Fig. 5(a)-6(c)). For the Fatigue-survived samples, on the other hand, the development of these damage modes except continued fiber fracture led to an initial modest increase in elapsed strain and decrease in modulus {Fig. 5(c) and (d)). For these samples, after a certain number
of cycles, the structural adjustments associated with the increase of ply delamination and bundle splitting become the dominant damage mechanisms for the composite. These structural adjustments, i.e., the straightening of fibers and the realignment of the fiber bundles, can redistribute the load amongst the fiber bundles more evently without breaking the fibers. Therefore, the composite experienced saturation of the elapsed strain and a modulus recovery at high cycle numbers. These mechanisms are also responsible for the increase in residual strength, compared to that of the monotonic tensile strength (Table 1 and Fig. l), because a higher stress can be taken by the realigned and straightened bundles in the final tensile test if the fibers carry it more evenly. A similar increase of the residual strength in a balanced cross-woven C/Sic composite after lo6 cycles was reported in [15]. In that study, the increase in residual strength was proposed to result from the relief of the cross concentration near fiber bundle crossovers by a process of matrix fracture and interfacial debonding. Given the complexity in the damage and fracture mechanisms observed here, such a stress relief mechanism could very well operate in the subject composite. The sample specific fatigue behavior, indicated by the relatively large data scatter in the fatigue test results (Fig. 3 and Fig. 5(a)-(d) and Table l), can be attributed to the microstructural differences in the composite samples, since the testing conditions were the same for all the fatigue samples. This sample specific behavior was not seen in monotonic tensile loading for the same sample size and design, and testing facility [16]. Therefore, it is reasonable to suggest that the sensitivity of the fatigue behavior to the structural variations in the cross woven C/Sic composite is greatly increased by the high loading rate (about 640760 in fatigue versus 3-5 MPa s - ’ in the monotonic tension) and the load repetitions in the fatigue test. A high loading rate in fatigue was found to cause more delaminations in another cross-woven C/Sic composite (X. Wu. and J.W. Holmes, personal communication) and in addition, the repeated loading of the composite was also observed to lead to the growth of the dalamination and bundle splitting in all the fatigue samples in the present study. Therefore, the combined effect of high loading rate and repeated loading produced a large amount of ply delamination and bundle splitting in the cyclic samples, increasing their influence on the fatigue behavior of the C/Sic composite, as described above. Since ply delamination and bundle splitting can most readily occur in areas of structural irregularity, different densities of these structural irregularities can lead to different amounts of ply delamination and bundle splitting. Due to the ‘random-phase lay up’ of the cross woven fabric in the present composite (Fig. 2), different densities of structural irregularities can exist
among fatigue samples. Therefore, different amounts of ply delamination and bundle splitting can be generated for all the fatigue samples, leading to sample specific fatigue behavior, and thus, a large scatter in the experimental results. Similar phenomena were observed in the compression behavior of the same C/Sic composite, in which ply delamination was one of the main damage modes [18].
5. Conclusions Studies of the tension-tension fatigue behavior of a C/Sic composite and its damage and fracture mechanisms lead to the following conclusions: (1) The observed fatigue stress versus life relationship (S-N diagram) showed the usual trend that as the fatigue stress increases, the fatigue life decreases. Runouts without fracture which occurred at IO6 cycles of fatigue are concluded to represent a fatigue limit in the area of 320-340 MPa, above which fatigue failure took place by a number of mechanisms but particularly the the mechanism of continued fiber fractures occurring mostly in the longitudinal bundle crossover sections (weak sections) during fatigue. (2) The fatigue behaviors of the cycled samples were found to result from the continued development of the cyclic damage, including the formation of ply delamination, bundle splitting, matrix damage zones in the fiber bundles, matrix wear, and fiber fracture. For fatiguefailed samples (cycled above the fatigue limit), the development of all the cyclic damage modes caused the elapsed strain to increase and the composite modulus to decrease continuously with cycling until fatigue failure occurred. For the fatigue-survived samples (cycled at fatigue limit), the development of the cyclic damage modes, except for continued fiber fracture, caused the elapsed strain to increase and the composite modulus to decrease initially. At high cycles up to lo6 cycles, structural adjustments associated with ply delamination and bundle splitting produced modulus recovery and even high residual strength in fatigue-survived samples. This behavior was caused by more even load distribution in the composite without fiber fracture. (3) The high loading rate and repeated loading in fatigue are believed to increase the sensitivity of the fatigue behavior to microstructural inhomogeneities existing in the C/Sic composite, by producing more ply delamination and bundle splitting. Different densities of structural irregularities, at which ply delamination and bundle splitting take place, exist in different samples due to the ‘random-phase’ lay up of the fabric in the composite, causing the sample specific nature of the fatigue behavior in the C/Sic composite.
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Acknowledgements
This work was supported by both the Benjamin Franklin Program of the Commonwealth of Pennsylvania and the Materials Science Corporation. The authors are deeply grateful for this support. The authors also thank Dr Alex Radin for his expert help in mechanical testing of the composite samples, and MS Xuqing Wang for her help in characterization studies. The Laboratory for Research on the Structure of Matter of the University of Pennsylvania both supported the work and provided experimental facilities. The authors acknowledge as well a crucial equipment grant from the Research Foundation of the University of Pennsylvania.
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