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Tensile fatigue behaviour of tightly woven carbon/carbon composites A. OZTURK* and R.E. MOOREt (* Middle East Technical University, Turkeyff University of Missouri-Rolla, USA) Received 1July 1991; accepted in revised form 7 November 1991 The tensile fatigue behaviour of a tightly woven carbon/carbon composite was investigated as a function of stress level. Load-controlled fatigue tests were performed in tension-tension mode with a stress ratio, R, of 0.1 under ambient laboratory conditions. Results of composite behaviour are discussed in terms of the relationship of the stress/strain behaviour to the fatigue life of these composites as well as the effects of applied stress levels. It is shown that these composites exhibit ~ood resistance to cyclic loading. No fatigue failures were obtained after 10° cycles when the maximum tensile load in the fatigue cycle is less than or equal to 80% of the static tensile strength. Evidence of textural changes related to fatigue was observed in the matrix region of these composites.
Key words: carbon/carbon composites; fatigue; static tensile strength; elastic modulus; stress~strain hysteresis; fatigue failure; fatigue run out; striations Carbon fibre-reinforced carbon-matrix composites, the so-called carbon/carbon composites, are of great importance since they possess improved mechanical and fracture mechanics properties compared with ceramics and superalloys. Monolithic carbons can be extremely brittle. However, carbon fibre-reinforced carbons exhibit a high fracture toughness and pseudoplasticity a-3. The most attractive properties of these composites are their high specific strength and elastic modulus which are the result of the high degree of anisotropy in the graphite crystal lattice and therefore in the fibres and tows4. Their small thermal expansion coefficient and good strength retention at high temperatures, as well as high thermal shock resistance, make them useful for structural materials 1-7. Moreover, they retain their high thermal and chemical stability in inert environments r -8. Because of these superior thermal and mechanical properties which persist at high temperatures, carbon/carbon composites find special applications such as in exit nozzles for rockets, nose caps and leading edges for missiles and space shuttle, nuclear reactors and especially for fusion devices 1-7. Newer applications such as hot press dies, wind tunnel models, racing car components, commercial disc brakes and sporting goods are being developed 1'2"5. Studies on the development of carbon/carbon composites with different carbon fibres and matrix precursors have been reported in the literature 1'9-14. These studies
have been mainly concerned with the effect of processing parameters on the matrix microstructure and its relation to the mechanical and thermal properties of these composites. The fatigue behaviour of carbon/ carbon composites has remained an area of research relatively overlooked due to the fact that they show limited crack-tip plasticity 15"16. Therefore, their mechanical properties will not degrade significantly when cyclically loaded. However, recent studies by Prewo and co-workers 17-~9have shown that evidence exists for the presence of real mechanical fatigue effects in ceramics and ceramic-matrix composites. In other studies 2°-23 cyclic loads were found to result in a lower time to failure than static loads. Therefore, the fracture resistance of carbon/carbon composites under cyclic loads is a subject of considerable scientific and engineering interest. Consequently, a clear understanding of the resistance to mechanical fatigue conditions is essential because repeated application of loads may affect the service life of these composites. The purpose of this investigation was to determine the influence of cyclic fatigue damage on the mechanical properties of a tightly woven carbon/carbon composite. Mechanical testing was supplemented with microstrucrural characterization to provide information necessary to explain observed behaviour. Fractographic analysis of fracture surfaces was conducted to locate fracture sources and to study the fracture propagation patterns.
0010-4361/92/010039-08 © 1992 Butterworth-Hei nemann Ltd COMPOSITES. VOLUME 23. NUMBER 1. JANUARY 1992
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EXPERIMENTAL PROCEDURE
Microstructure analysis
Specimen preparation
Scanning electron microscopy (SEM) (JEOL JSMT330A) techniques were used for analysis of fracture surfaces and for microstructural analysis. Specimens were prepared for fractographic analysis by cutting the fracture end from the bar. No coating was necessary for the specimen to be analysed.
The two carbon/carbon composite panels used in this study (designated panel 1 and panel 2) were commercial materials (K-Karb) obtained from Kaiser Aerotech Company, San Leandro, CA, USA. The composite consists of a graphite fibre-reinforced graphite matrix developed for aerospace applications. Composites are formed by laying up cut-out patterns of plain weave woven fabrics in moulds. Then, the formed structure is impregnated with resin under pressure. Application of heat and pressure results in curing of the matrix and fixing of the carbon fibres in the desired shape. This is followed by a densification cycle involving carbonization in an inert atmosphere and re-impregnation by resin to obtain the desired density. Finally, the resultant product is graphitized at 2500°C. Fatigue test specimens were prepared by cutting the composite panel into small rectangular bars along the warp direction (x-direction) using a diamond saw. The nominal dimensions of these bars were 101.6 by 7.6 by 6.3 mm. Rectangular-shaped composite bars were then machined with a surface grinder using a 220 grit diamond wheel to reduce the gauge section. The length and the width of the parallel-sided gauge section were approximately 25.4 and 5.0 mm, respectively. A typical fatigue test specimen geometry is shown in Fig. 1. Finally a glued-on strain gauge, supplied by MTS Measurement Group Inc, was mounted on the specimen to measure the applied strain.
Fatigue testing Load-controlled tension-tension fatigue tests were performed in air at room temperature using a servo hydraulic testing machine (MTS model 810 Materials Test System). Two special grip tabs were desinged to prevent specimens from crushing during gripping. Specimen ends were clamped into the steel tabs. The tabs were then connected to a double set of universal joints through the holes in the tabs with steel pins. The entire unit was attached to the hydraulic grips of the MTS machine. Once the test assembly was set up, loads were applied to the specimen between a maximum (Pmax) and a minimum tension load (Pmin) following a sinusoidal waveform at 10 Hz frequency. The stress ratio, R, was maintained at 0.1 for all tests. Fatigue tests were run at different stress levels until specimen failure or fatigue run out (completion of 106 cycles). Specimen failure was taken as the point at which the specimen fractured into two pieces.
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RESULTS AND DISCUSSION At least five specimens from each composite panel were broken in tension to determine the average ultimate tensile strength (AUTS) of the corresponding panels. Results of all uniaxial static tensile tests are presented in Table 1 along with the average values and the + 1 standard deviation of each set of tests. Although the fabrication and testing procedures are exactly the same, the AUTS and fai!ure strains of the panels varied from 62.3 _+ 3.9 MPa to 104.8 + 5.2 MPa and from 0.41 _ 0.01% to 0.90 + 0.01%, respectively. The plus and minus signs indicate the + I standard deviation of the averages. However, the uncertainties of the data were calculated using error analysis as + 0.1 MPa, + 0.01 GPa and _+ 0.01% for the static tensile strength, the elastic modulus and the failure strain, respectively, for individual tensile tests and as + 0.5 MPa for the maximum applied stress level for individual fatigue tests. The static tensile strength of specimens broken in tension ranged from 99.5 + 0.1 MPa to 112.8 _+ 0.1 MPa for composite panel 1 and from 56.3 + 0.1 MPa to 66.1 + 0.1 MPa for composite panel 2. The variation of mechanical properties from panel to panel is attribTable carbon
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Static composite
Panel no.
Tensile Elastic Failure strength modulus strain (MPa+0.1) (GPa_+0.01) (%+0.01) 99.5
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-
105.8
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1
100.9
1
105.0
1
112.8 Average
Standard deviation Variance
104.8
-
11.66 -
-
0.90 -
11.66
0.90
15.59
0.37
5.2 27.1
2
57.5
2
64.8
2
64.8
2
62.1
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-
2 2
64.6
15.10
0.43
66.1
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2
56.3
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Variance
COMPOSITES. J A N U A R Y 1992
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1
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properties
panels
1
Average
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tensile
62.3 3.9 15.2
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15.38
15.36
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0.43
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0.41
uted primarily to variation of the control parameters (temperature, pressure and time) during fabrication. Despite differences in the AUTSs of both composite panels, no major changes were observed in the microstructure. Debonding of the fibre/matrix interface, fibre pull-out, spearation of matrices and delamination of plies are observed in composites. Although the maximum values for stresses and strains varied, the stress/strain curves of the composite panels exhibited similar behaviour. A representative tensile stress vs. strain curve for these composites is depicted in Fig. 2. Composites exhibited nearly linear stress/ strain behaviour up to a maximum stress. After the maximum stress is achieved, a small non-linearity in the form of a slight decrease in stress with increasing strain was followed by an abrupt drop of the stress. Failure was catastrophic and occurred in the gauge section. All specimens failed in a fibrous mode with some shear parallel to the fibres in the warp direction. Results of fatigue tests are presented as a plot of applied stress levels vs. the number of cycles in Fig. 3 for panel 1 and in Fig. 4 for panel 2. The specimens fatigue tested at stress levels less than 90 _+ 0.5 MPa survived after 106 cycles for panel 1. However, all the specimens but two tested at greater stress levels failed during fatigue. The residual strength of the surviving specimens fatigued at different stress levels was comparable to the static strength. The average residual strength of nine survivors was 103.6 _+ 5.0 MPa while the AUTS of specimens which had not undergone fatigue tests was 104.8 _+ 5.2 MPa. In the case of composite panel 2, it was possible to see fatigue failure with a stress as low as 53 _+ 0.5 MPa. Specimens fatigued at lower maximum applied stress levels survived after 106 cycles. When the stress level was increased to 53 + 0.5 MPa, the composite exhibited either fatigue failure or fatigue run out. For those specimens that survived after 106 cycles, residual strengths were equivalent to or even greater than those of unfatigued specimens. The average residual strength of seven surviving specimens was 63.8 _ 3.4 MPa while the AUTSwas 62.3 _+_3.9 MPa. Three of the tests were interrupted after 105 cycles to see whether there 150
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was a difference in the composite strength after 105 and 106 cycles. The average strength of the three specimens after 105 cycles was 62.6 + 3.6 MPa which is within the error limits of the AUTS and the residual strength of composite panel 1. It appears that fatigue testing does not have an effect on the residual strength of these composites regardless of the stress levels. The independence of the residual strength to fatigue stress levels is shown in Fig. 5 for composite panel 2. Similar results were reported for SiC fibre-reinforced lithium aluminosilicates by Prewo 17. The independence of the residual strength to fatigue stress levels is attributed to some form of 60
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COMPOSITES. J A N U A R Y 1992
41
fibre/matrix decoupling of the composite which reduces fibre damage during fatigue 18.
14/27. This indicates that the composite undergoes considerable fatigue damage for the cyclic loads applied between these stress levels. Specimens fatigue tested at stress levels greater than 90 PAS failed during testing.
All fatigued specimens failed in the gauge section, exhibiting a fibrous type of failure similar to that of the static tensile specimens. However, examination of the fracture surfaces revealed that fatigued specimens showed a greater degree of fibre pull-out and less shear than the specimens tested under static loading. This greater degree of fibre pull-out was also reported for SiC fibre-reinforced lithium aluminosilicates by Minford and Prewo TM,who suggested that more extensive fibre deterioration or fibre/matrix decoupling is occurring during fatigue cycles. The difference between the appearance of both test specimens was also observed for those specimens which survived after 10 6 cycles and were then broken in tension. The change in the appearance of both test specimens was one of the evidences of fatigue for these composites.
Composite elastic moduli were measured after 1, 10, 102, 103, i04, 105 and 106 cycles to determine whether any structural deterioration had occurred with increasing number of cycles. It is shown that increasing number of cycles does not have a significant effect on the composite elastic modulus. Elastic modulus of a specimen after 10 6 cycles was approximately 2% less than the modulus of the original specimen. The composite elastic modulus decreases within the first 100 cycles as indicated in Fig. 7. Beyond that elastic modulus seems to remain unchanged through 10 6 cycles. This behaviour was also observed in the stress/ strain hysteresis loop of these composites. The shape of the stress/strain hysteresis changed considerably indicating a change in the composite elastic modulus. As shown in Fig. 8, the stress/strain hysteresis loops become progressively narrower giving smaller moduli as the number of cycles increases. The change in the shape of the stress/strain hysteresis loop would imply the gradual accumulation of fatigue damage in the composites.
It is apparent from Figs 3 and 4 that the number of cycles to failure decreases with increasing applied stress levels. The composite fails at a lower number of cycles for greater stress levels. In other words, the resistance to cyclic fatigue decreases with increasing stress levels. Hence, the composite fatigue life decreases at greater applied stress levels.
Specimens exhibited two different lines on their stress/strain hysteresis loop corresponding to each loading and unloading. The slope of the lines indicated two different moduli for the composite. The two moduli are referred to as primary and secondary moduli. The primary modulus is the slope of the line obtained during loading, while the secondary modulus is that obtained during unloading. At the beginning of the fatigue test there are no cracks present. Fibres in both the warp and fill direction and the matrix contribute to the stiffness of the composite. As the load increases, microcracks begin to form in the fill direction. Thus, the contribution of the fibres in this direction and the matrix to the stiffness decrease. The decrease in stiffness eventually results in a decrease in the composite elastic modulus. Agarwal and Broutman 24 have reported that the composite elastic modulus decreases first in the presence of cross-ply cracks and then longitudinal-ply cracks are followed by delamination cracks. Therefore, with fatigue exposures, the stress/strain hysteresis of the composite
In order to compare the tension fatigue behaviour of the two composite panels, it was necessary to normalize the applied stress levels. Normalization of the data was done by taking the ratio of the repeated maximum stress level to the average ultimate tensile strength for each composite panel. The ratio was then multiplied by a factor of 100 in order to obtain the percentages. The sum was defined as percent average strength (PAS) for convenience. The variation in the number of cycles to failure as a function of PAS is plotted in Fig. 6, which also summarizes the cyclic fatigue data for these composites. Failure did not occur for stress levels less than 80 PAS. All the specimens tested below this stress level survived a f t e r 10 6 cycles. Specimens fatigue tested between 80 and 90 PAS either failed during testing or survived after 10 6 cycles. This range of PAS is considered as a critical range for cyclic fatigue of these composites. The ratio of the number of survival specimens to the total number of specimens fatigue tested in this range is
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42
C O M P O S I T E S . J A N U A R Y 1992
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becomes linear with a modulus close to the secondary modulus. As seen in Fig. 8, the primary and secondary moduli are the same at 105 cycles giving a straight line in the stress/strain hysteresis. The primary modulus becomes less than the secondary modulus at 106 cycles. This has been attributed to the development of longitudinal cracks and delamination cracks due to the fatigue loading 24. The SEM analysis of fatigue tested specimens revealed a
textural change associated with increasing number of cycles. The textural changes occurred in the matrix regions, the most likely regions for origination of cracks and their extension. Two representative fractographs of the matrix regions of static tensile and fatigue tested specimens are shown in Fig. 9. The surface of the matrix of a static tensile tested specimen is smooth and appears to be glassy (Fig. 9(a)). However, the surface of the matrix of a fatigue tested specimen has a rough topography (Fig. 9(b)). The striations, which indicate the advance of crack growth after each cycle, are also observed on the fracture surfaces of a fatigue tested specimen. These striations were not detected on the fracture surfaces of the static tensile tested specimens. However, smooth and rough surfaces in the matrix regions were observed in both test specimens. In general, the surface of the matrix parallel to the fracture surface seemed to be glassy while the surfaces of the matrix normal and/or at a high angle to the fracture surface had crystalline-type character, as shown in Fig. 10. Furthermore, the surfaces normal and/or at a high angle to the fracture surface had more topography and were rougher than the surfaces parallel to the fracture surface. Although the striations were detected in the surfaces parallel to the fracture surface, they are mostly observed on the surfaces normal and/or at a high angle to the fracture surface for the fatigue tested specimens as indicated in Fig. 11. A number of fatigue tests was conducted to determine the threshold cycle number, Nt, which is defined as the number of fatigue cycles below which no textural change occurs in various microstructures. Eleven specimens of composite panel 1 were fatigue tested for various numbers of cycles ranging from 1000 to 40 000. A stress level of 80 PAS was held constant for all fatigue tests. None of the specimens failed during testing. Upon completion of the fatigue tests, specimens were broken in tension to determine the residual strengths. The residual strengths of these specimens ranged from 99.2 + 0.1 M P a t o 107.9 + 4.9 MPa. The average residual strength of the eleven specimens was 104.1 + 4.9 MPa. Resultant data show that the residual strengths of the surviving specimens were comparable to those of unfatigued specimens. However, microstructural changes were observed in
Fig. 9 SEM fractographs of carbon/carbon composites showing the matrix region of the fracture surfaces: (a) static tensile tested; (b) fatigue tested and survived after 106 cycles
Fig. 10 SEM fractograph of a carbon/carbon composite fatigue
tested and failed after 1240 cycles
C O M P O S I T E S . J A N U A R Y 1992
43
from the matrix, resulting in debonding of the fibre/ matrix interfaces and pull-out of fibres. The debonded fibre/matrix interface serves as a path for a single large crack which propagates through the cross-section area of the gauge section.
Fig. 11 SEM fractograph of a carbon/carbon composite fatigue tested and failed after 3618 cycles
the matrix region of the specimens as the number of fatigue cycles increased. The striations become more obvious on the surfaces normal to or at a high angle to the fracture surface. It was not, however, possible to determine a precise value for Nt since all the specimens had some degree of microstructural change during fatigue cycles. The development of damage under tensile fatigue loading may be explained phenomenologically based on direct observations of fatigue tests and SEM fractographs of the test specimens. Within the first 10 load cycles, several non-critical radial and circumferential microcracks form in the matrix and at the fibre/matrix interface provided that applied stress exceeds the local ply strength. The critical stress for the formation of matrix cracking has been evaluated by Aveston et al. 25 who reported that large stress and strain concentrations at the fibre/matrix interface are responsible for the initiation of these cracks. Manocha and Bahl 7 observed that these microcracks develop beneath the bundle cross-over points due to differences in thermal expansion characteristics of sharply bent fibres and the inhomogeneous matrix distribution around these crossed bundles. The cracking process may continue until the cracks in each ply have attained an equilibrium. The equilibrium state in which the matrix cracking pattern is stabilized forms by 100 cycles. The number of cracks increases with either the number of cycles or an increase in the applied stress level 26'27. As the matrix cracks develop, fibre failure begins to initiate preferentially in adjacent plies in regions of stress concentration created by the primary cracks. After initiation the crack propagates between fibres primarily along the fibre/matrix interface 28-31.The radial cracks propagate up to the fibre surface. They either stop there or deflect and branch into small secondary cracks which extend short distances away from the interface between the ply in which the primary crack has occurred and the one in which the secondary crack occurs 26. These secondary cracks are generally perpendicular to the primary cracks (or as nearly so as possible) and are caused by the tensile stresses along the crack axis ahead of the primary cracks 27. The circumferential cracks cause the separation of the fibres
44
COMPOSITES • JANUARY
1992
The microcracks perpendicular to the applied load propagate through the entire width of the ply but are unable to propagate into the adjacent ply. They terminate at the interface of two plies. However, the crack tip produces a stress concentration ahead of itself. The resulting strong interlaminar stresses in regions where primary and secondary cracks cross produce favourable conditions for starting an interlaminar crack along the ply interface. This phenomenon, known as delamination, is caused by a mixed-mode growth of interlaminar cracks 26. More delamination cracks start and propagate through the interface as the number of cycles increases. Deflection of the crack path along the boundary between fibre bundles and cloth layers is controlled by the nature of the fibre/matrix interface. If the bond strength is relatively weak, the crack propagating through the matrix deflects and branches when it reaches a boundary between the fibre bundles and cloth layers where the delamination cracks already exist. Part of the deflection is due to debonding of the fibre/matrix interface, pull-out of the fibres from the matrix and the associated increase in crack opening 32. The crack propagation path may also deviate along the planes of circumferential microcracks which are likely to behave as new interfaces 6. The deflection of the crack reduces the stress concentrations at the crack tip 6"28. In the case of interfaces with relatively strong bonds containing both fibre/matrix interfaces and planes of circumferential microcracks in the matrix, it is difficult to effect relaxation of the stress concentration of the composites around the fibre effectively because of the low degree of interface debonding 6. In this case, cracks propagate with relatively few impediments by circumferential microcracking through the matrix. The composite fractures in the primary crack plane with no significant branching or deflection. The composite undergoes final fracture when it is sufficiently weakened by delamination cracks. The presence of delamination cracks prevents load distribution between plies, and the composite is essentially reduced to a number of independent plies acting in parallel to support the applied load. The weakest of these plies fails and starts failure of the remaining plies. Development of delamination cracks is slow, and hence the loss of strength in the later part of the fatigue life is slow. Rapid loss of strength occurs in the last few cycles of fatigue life when the stronger plies fail. Prior to this rapid loss of strength, individual plies do become weakened, but the overall strength reduction is slow. At the time when delamination cracks appear, fibres within a short distance of the crack tip remain intact and bridge the crack whereas the remainder of the fibres in the crack wake start fracturing and debonding 33. Failure of the composite may advance with either matrix crack growth or fibre failure. The crack covers the whole cross-sectional area of the gauge section when the maximum stress achieved. Beyond this point
applied load is s u p p o r t e d by the intact fibres bridging the crack. R e m a i n i n g bridging fibres break due to the large stress concentrations at the surface. C o m p o s i t e failure occurs w h e n the applied stress exceeds the strength of fibre bundles. Fibre bundle failure is the f u n d a m e n t a l process that dictates the ultimate strength, and axial crack extension is then merely a c o n s e q u e n c e of fibre bundle failure 34. T h e later stage of d a m a g e d e v e l o p m e n t is a rapidly increasing rate of progression o f the mixture of all d a m a g e modes. It must be n o t e d that individual d a m a g e m o d e s do not f o r m in distinct regions but rather occur m o r e or less simultaneously. H o w e v e r , each d a m a g e m o d e m a y have a range of d o m i n a n c e .
and wave pattern on the development of 2D carbon-carbon composites' Carbon 26 No 1 (1988) pp 13-21 8 Ozturk, A. 'The mechanical properties of tightly woven carboncarbon composites' PhD thesis (University of Missouri-Rolla, Rolla, MO, USA, 1991) 9 Yasuda, E., Tanabe, Y., Manoeha, L.M. and Kumura, S.
10 11 12 13
CONCLUSIONS 1) C a r b o n / c a r b o n composites investigated show g o o d fatigue resistance to fluctuating loads. Stress to survive 106 cycles was a p p r o x i m a t e l y 80% of the ultimate tensile strength. A notable, but not significant, c h a n g e in composite residual strength and elastic m o d u l u s was o b s e r v e d as a result of fatigue loading. These results are e n c o u r a g i n g since m a n y potential applications of fibre-reinforced ceramiccomposites require a high fatigue life. 2) T e n s i o n - t e n s i o n cyclic fatigue loading caused a substantial c h a n g e in the stress/strain hysteresis loop of these composites, indicating the progressive accumulation of fatigue d a m a g e . This c h a n g e in behaviour, p r e s u m e d to be due to the o c c u r r e n c e of matrix microcracking, did not a p p e a r to have a significant effect o n composite residual strength. 3) S E M analysis of b r o k e n specimens revealed that these composites show textural changes related to fatigue. Matrix cracking and fibre bundle failure are the d o m i n a n t failure m e c h a n i s m s during fatigue. Since matrix cracks are not deleterious to the fibres, a fibre-dominant fracture is p r o n e in these c o m p o s ites. 4) This investigation d e m o n s t r a t e s that fluctuating stresses induce fatigue d a m a g e even at r o o m temp e r a t u r e in air. T h e r e f o r e , the cyclic fatigue behaviour o f a d v a n c e d structural ceramic composites needs to be e x a m i n e d if they are to be used in critical applications involving r e p e a t e d load conditions.
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14 15 16 17 18
19 20 21
'Matrix modification by graphite powder additives in carbon fiber/carbon composite with thermosetting resin precursor as a matrix' Carbon 26 No 2 (1988) pp 225-227 Oh, S.-M. and Lee, J.-Y. 'Structures of pyrolytic carbon matrices in carbon/carbon composites' Carbon 26 No 6 (1988) pp 763-768 Oh, S.-M. and Lee, J.-Y. 'Effects of matrix structure on mechanical properties of carbon/carbon composites' Carbon 26 No 6 (1988) pp 769-776 Manocha, L.M., Yasuda, E., Tanabe, Y. and Kimura, S. 'Effect of carbon fiber surface-treatment on mechanical properties of C/C composites' Carbon 26 No 3 (1988) pp 333-337 Sato, S., Kurumada, A., Iwaki, H. and Komatsu, Y. 'Tensile properties and fracture toughness of carbon-fiber felt reinforced carbon composites at high temperature' Carbon 27 No 6 (1989) pp 791-801 Manocha, L.M., Bahl, O.P. and Singh, Y.K. 'Mechanical behavior of carbon-carbon composites made with surface treated carbon fibers' Carbon 27 No 3 (1989) pp 381-387 Evans, A.G. 'Fatigue in ceramics' lntJ Fract 16 No 6 (1980) pp 486-489 Lewis,D. and Rice, R.W. 'Comparison of static, cyclic and thermal-shock fatigue in ceramics composites' Ceram Engng Sci Proc 3 Nos 9-10 (1982) pp 714-721 Prewo, K.M. 'Fatigue and stress rupture of silicon carbide fibre reinforced glass ceramics' J Mater Sci 22 (1987) pp 2695-2701 Minford, E. and Prewo, K.M. 'Fatigue behavior of silicon carbide fiber reinforced lithium-alumino-silicate glass ceramic' in Patano, C.G. and Messing, R.E. (Eds) TailoringMultiphase and Composite Ceramics (Plenum Publishing Corp, 1986) pp 561570 Nardone, V.C. and Prewo, K.M. 'Tensile performance on carbon-fibre-reinforced glass' J Mater Sci 23 (1988) pp 168-180 Suresh, S. and Han, L.X. 'Fracture of Si3N4-SiCwhisker composites under cyclic loads' J Amer Ceram Soc 71 No 3 (1988) pp C158-161 Holmes,J.W., Kotil, T. and Folds, W.T. 'High temperature fatigue of SiC fiber-reinforced Si3N4 ceramic composites' in Proc ASC Symp on High Temperature Composites, Dayton, OH, USA, June 1989 (Technomic Publishing Co, Lancaster, PA,
USA) pp 176-186 22 Zawada, L.P. and Butkus, L.M. 'Room temperature tensile and fatigue properties of silicon carbide fiber-reinforced aluminosilicate glass' Ceram Engng Sci Proc 11 Nos 9-10 (1990) pp 15921606 23 Morrone, A.A., Nntt, S.R. and Suresh, S. 'Fracture toughness and fatigue crack growth behavior of an A1203-SiC composite' J Mater Sci 23 (1988) pp 3206-3213 24 Agarwal, B.D. and Brontmau, L.J. Analysis and Performance of Fiber Composites (John Wiley, New York, USA, 1980) pp 223246 25 Aveston, J., Cooper, G.A. and Kelly, A. 'Single and multiple fracture' in The Properties of Fiber Composites (Conference Proceedings of the National Physical Laboratory, IPC Science and Technology Press Ltd, Surrey, UK, 1971) pp 15-26 26 Reifsnider, K.L., Henneke, E.G., Stinchcomb, W.W. and Duke, J.C. 'Damage mechanics and NDE of composite laminates' in Hashin, Z. and Herakovich, C.T. (Eds) Mechanics of Composite Materials, Recent Advances (Pergamon Press, New York, USA, 1983) pp 399-420 27 Talreja, R. Fatigueof Composite Materials (Technomic Publishing Co, Lancaster, PA, USA, 1987) pp 59-70 28 Dharan, C.K.H. 'Fatigue failure mechanisms in a unidirectionally reinforced composite material' ASTM STP 569 (American Society for Testing and Materials, 1975) pp 171-188 29 Lorenzo, L. and Hahn, H.T. 'Fatigue failure mechanisms in unidirectional composites' ASTM STP 907 (American Society for Testing and Materials, 1987) pp 210-232 30 Salkind, M.J. 'Fatigue of composite materials' ASTM STP 497 (American Society for Testing and Materials, 1982) pp 143-169 31 Curtis, P. 'An investigation of tensile fatigue behavior of carbon fibre composite laminates' ICCM6 Proceedings, Imperial College of Science and Technology, London, UK (Elsevier Applied Science Publishers, 1987) pp 54-64
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. JANUARY
1992
45
32 Marshall, D.B. and Evans, A.G. 'Failure mechanisms in ceramicfiber/ceramic-matrix composites' J Amer Cer Soc 68 No 5 (1985) pp 225-231 33 Marshall, D.B. and Cox, B.N. 'Tensile fracture of brittle matrix composites: influence of fiber strength' Carbon 35 No 26 (1987) pp 2607-2619 34 Sbaizero, O. and Evans, A.G. 'Tensile and shear properties of laminated ceramic matrix composites' J A m e r Ceram Soc 69 No 6 (1986) pp 481--486
46
COMPOSITES
. JANUARY
1992
A U THORS
A. Ozturk is with the Metallurgical Engineering Department, Middle East Technical University, Ankara, Turkey. R.E. Moore is with the Ceramic Engineering Department, University of MissouriRolla, Rolla, MO, USA.
Key words: carbon/carbon composites; fatigue; static tensile strength; elastic modulus; stress~strain hysteresis; fatigue failure; fatigue run out; striations Carbon fibre-reinforced carbon-matrix composites, the so-called carbon/carbon composites, are of great importance since they possess improved mechanical and fracture mechanics properties compared with ceramics and superalloys. Monolithic carbons can be extremely brittle. However, carbon fibre-reinforced carbons exhibit a high fracture toughness and pseudoplasticity a-3. The most attractive properties of these composites are their high specific strength and elastic modulus which are the result of the high degree of anisotropy in the graphite crystal lattice and therefore in the fibres and tows4. Their small thermal expansion coefficient and good strength retention at high temperatures, as well as high thermal shock resistance, make them useful for structural materials 1-7. Moreover, they retain their high thermal and chemical stability in inert environments r -8. Because of these superior thermal and mechanical properties which persist at high temperatures, carbon/carbon composites find special applications such as in exit nozzles for rockets, nose caps and leading edges for missiles and space shuttle, nuclear reactors and especially for fusion devices 1-7. Newer applications such as hot press dies, wind tunnel models, racing car components, commercial disc brakes and sporting goods are being developed 1'2"5. Studies on the development of carbon/carbon composites with different carbon fibres and matrix precursors have been reported in the literature 1'9-14. These studies
have been mainly concerned with the effect of processing parameters on the matrix microstructure and its relation to the mechanical and thermal properties of these composites. The fatigue behaviour of carbon/ carbon composites has remained an area of research relatively overlooked due to the fact that they show limited crack-tip plasticity 15"16. Therefore, their mechanical properties will not degrade significantly when cyclically loaded. However, recent studies by Prewo and co-workers 17-~9have shown that evidence exists for the presence of real mechanical fatigue effects in ceramics and ceramic-matrix composites. In other studies 2°-23 cyclic loads were found to result in a lower time to failure than static loads. Therefore, the fracture resistance of carbon/carbon composites under cyclic loads is a subject of considerable scientific and engineering interest. Consequently, a clear understanding of the resistance to mechanical fatigue conditions is essential because repeated application of loads may affect the service life of these composites. The purpose of this investigation was to determine the influence of cyclic fatigue damage on the mechanical properties of a tightly woven carbon/carbon composite. Mechanical testing was supplemented with microstrucrural characterization to provide information necessary to explain observed behaviour. Fractographic analysis of fracture surfaces was conducted to locate fracture sources and to study the fracture propagation patterns.
0010-4361/92/010039-08 © 1992 Butterworth-Hei nemann Ltd COMPOSITES. VOLUME 23. NUMBER 1. JANUARY 1992
39
EXPERIMENTAL PROCEDURE
Microstructure analysis
Specimen preparation
Scanning electron microscopy (SEM) (JEOL JSMT330A) techniques were used for analysis of fracture surfaces and for microstructural analysis. Specimens were prepared for fractographic analysis by cutting the fracture end from the bar. No coating was necessary for the specimen to be analysed.
The two carbon/carbon composite panels used in this study (designated panel 1 and panel 2) were commercial materials (K-Karb) obtained from Kaiser Aerotech Company, San Leandro, CA, USA. The composite consists of a graphite fibre-reinforced graphite matrix developed for aerospace applications. Composites are formed by laying up cut-out patterns of plain weave woven fabrics in moulds. Then, the formed structure is impregnated with resin under pressure. Application of heat and pressure results in curing of the matrix and fixing of the carbon fibres in the desired shape. This is followed by a densification cycle involving carbonization in an inert atmosphere and re-impregnation by resin to obtain the desired density. Finally, the resultant product is graphitized at 2500°C. Fatigue test specimens were prepared by cutting the composite panel into small rectangular bars along the warp direction (x-direction) using a diamond saw. The nominal dimensions of these bars were 101.6 by 7.6 by 6.3 mm. Rectangular-shaped composite bars were then machined with a surface grinder using a 220 grit diamond wheel to reduce the gauge section. The length and the width of the parallel-sided gauge section were approximately 25.4 and 5.0 mm, respectively. A typical fatigue test specimen geometry is shown in Fig. 1. Finally a glued-on strain gauge, supplied by MTS Measurement Group Inc, was mounted on the specimen to measure the applied strain.
Fatigue testing Load-controlled tension-tension fatigue tests were performed in air at room temperature using a servo hydraulic testing machine (MTS model 810 Materials Test System). Two special grip tabs were desinged to prevent specimens from crushing during gripping. Specimen ends were clamped into the steel tabs. The tabs were then connected to a double set of universal joints through the holes in the tabs with steel pins. The entire unit was attached to the hydraulic grips of the MTS machine. Once the test assembly was set up, loads were applied to the specimen between a maximum (Pmax) and a minimum tension load (Pmin) following a sinusoidal waveform at 10 Hz frequency. The stress ratio, R, was maintained at 0.1 for all tests. Fatigue tests were run at different stress levels until specimen failure or fatigue run out (completion of 106 cycles). Specimen failure was taken as the point at which the specimen fractured into two pieces.
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RESULTS AND DISCUSSION At least five specimens from each composite panel were broken in tension to determine the average ultimate tensile strength (AUTS) of the corresponding panels. Results of all uniaxial static tensile tests are presented in Table 1 along with the average values and the + 1 standard deviation of each set of tests. Although the fabrication and testing procedures are exactly the same, the AUTS and fai!ure strains of the panels varied from 62.3 _+ 3.9 MPa to 104.8 + 5.2 MPa and from 0.41 _ 0.01% to 0.90 + 0.01%, respectively. The plus and minus signs indicate the + I standard deviation of the averages. However, the uncertainties of the data were calculated using error analysis as + 0.1 MPa, + 0.01 GPa and _+ 0.01% for the static tensile strength, the elastic modulus and the failure strain, respectively, for individual tensile tests and as + 0.5 MPa for the maximum applied stress level for individual fatigue tests. The static tensile strength of specimens broken in tension ranged from 99.5 + 0.1 MPa to 112.8 _+ 0.1 MPa for composite panel 1 and from 56.3 + 0.1 MPa to 66.1 + 0.1 MPa for composite panel 2. The variation of mechanical properties from panel to panel is attribTable carbon
1.
Static composite
Panel no.
Tensile Elastic Failure strength modulus strain (MPa+0.1) (GPa_+0.01) (%+0.01) 99.5
-
-
105.8
-
-
1
100.9
1
105.0
1
112.8 Average
Standard deviation Variance
104.8
-
11.66 -
-
0.90 -
11.66
0.90
15.59
0.37
5.2 27.1
2
57.5
2
64.8
2
64.8
2
62.1
-
-
2 2
64.6
15.10
0.43
66.1
-
2
56.3
-
Variance
COMPOSITES. J A N U A R Y 1992
of carbon/
1
Standard deviation
40
properties
panels
1
Average
Fig. 1 S c h e m a t i c illustration o f test s p e c i m e n (all dimensions in m m )
tensile
62.3 3.9 15.2
-
15.38
15.36
-
0.43
-
0.41
uted primarily to variation of the control parameters (temperature, pressure and time) during fabrication. Despite differences in the AUTSs of both composite panels, no major changes were observed in the microstructure. Debonding of the fibre/matrix interface, fibre pull-out, spearation of matrices and delamination of plies are observed in composites. Although the maximum values for stresses and strains varied, the stress/strain curves of the composite panels exhibited similar behaviour. A representative tensile stress vs. strain curve for these composites is depicted in Fig. 2. Composites exhibited nearly linear stress/ strain behaviour up to a maximum stress. After the maximum stress is achieved, a small non-linearity in the form of a slight decrease in stress with increasing strain was followed by an abrupt drop of the stress. Failure was catastrophic and occurred in the gauge section. All specimens failed in a fibrous mode with some shear parallel to the fibres in the warp direction. Results of fatigue tests are presented as a plot of applied stress levels vs. the number of cycles in Fig. 3 for panel 1 and in Fig. 4 for panel 2. The specimens fatigue tested at stress levels less than 90 _+ 0.5 MPa survived after 106 cycles for panel 1. However, all the specimens but two tested at greater stress levels failed during fatigue. The residual strength of the surviving specimens fatigued at different stress levels was comparable to the static strength. The average residual strength of nine survivors was 103.6 _+ 5.0 MPa while the AUTS of specimens which had not undergone fatigue tests was 104.8 _+ 5.2 MPa. In the case of composite panel 2, it was possible to see fatigue failure with a stress as low as 53 _+ 0.5 MPa. Specimens fatigued at lower maximum applied stress levels survived after 106 cycles. When the stress level was increased to 53 + 0.5 MPa, the composite exhibited either fatigue failure or fatigue run out. For those specimens that survived after 106 cycles, residual strengths were equivalent to or even greater than those of unfatigued specimens. The average residual strength of seven surviving specimens was 63.8 _ 3.4 MPa while the AUTSwas 62.3 _+_3.9 MPa. Three of the tests were interrupted after 105 cycles to see whether there 150
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was a difference in the composite strength after 105 and 106 cycles. The average strength of the three specimens after 105 cycles was 62.6 + 3.6 MPa which is within the error limits of the AUTS and the residual strength of composite panel 1. It appears that fatigue testing does not have an effect on the residual strength of these composites regardless of the stress levels. The independence of the residual strength to fatigue stress levels is shown in Fig. 5 for composite panel 2. Similar results were reported for SiC fibre-reinforced lithium aluminosilicates by Prewo 17. The independence of the residual strength to fatigue stress levels is attributed to some form of 60
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COMPOSITES. J A N U A R Y 1992
41
fibre/matrix decoupling of the composite which reduces fibre damage during fatigue 18.
14/27. This indicates that the composite undergoes considerable fatigue damage for the cyclic loads applied between these stress levels. Specimens fatigue tested at stress levels greater than 90 PAS failed during testing.
All fatigued specimens failed in the gauge section, exhibiting a fibrous type of failure similar to that of the static tensile specimens. However, examination of the fracture surfaces revealed that fatigued specimens showed a greater degree of fibre pull-out and less shear than the specimens tested under static loading. This greater degree of fibre pull-out was also reported for SiC fibre-reinforced lithium aluminosilicates by Minford and Prewo TM,who suggested that more extensive fibre deterioration or fibre/matrix decoupling is occurring during fatigue cycles. The difference between the appearance of both test specimens was also observed for those specimens which survived after 10 6 cycles and were then broken in tension. The change in the appearance of both test specimens was one of the evidences of fatigue for these composites.
Composite elastic moduli were measured after 1, 10, 102, 103, i04, 105 and 106 cycles to determine whether any structural deterioration had occurred with increasing number of cycles. It is shown that increasing number of cycles does not have a significant effect on the composite elastic modulus. Elastic modulus of a specimen after 10 6 cycles was approximately 2% less than the modulus of the original specimen. The composite elastic modulus decreases within the first 100 cycles as indicated in Fig. 7. Beyond that elastic modulus seems to remain unchanged through 10 6 cycles. This behaviour was also observed in the stress/ strain hysteresis loop of these composites. The shape of the stress/strain hysteresis changed considerably indicating a change in the composite elastic modulus. As shown in Fig. 8, the stress/strain hysteresis loops become progressively narrower giving smaller moduli as the number of cycles increases. The change in the shape of the stress/strain hysteresis loop would imply the gradual accumulation of fatigue damage in the composites.
It is apparent from Figs 3 and 4 that the number of cycles to failure decreases with increasing applied stress levels. The composite fails at a lower number of cycles for greater stress levels. In other words, the resistance to cyclic fatigue decreases with increasing stress levels. Hence, the composite fatigue life decreases at greater applied stress levels.
Specimens exhibited two different lines on their stress/strain hysteresis loop corresponding to each loading and unloading. The slope of the lines indicated two different moduli for the composite. The two moduli are referred to as primary and secondary moduli. The primary modulus is the slope of the line obtained during loading, while the secondary modulus is that obtained during unloading. At the beginning of the fatigue test there are no cracks present. Fibres in both the warp and fill direction and the matrix contribute to the stiffness of the composite. As the load increases, microcracks begin to form in the fill direction. Thus, the contribution of the fibres in this direction and the matrix to the stiffness decrease. The decrease in stiffness eventually results in a decrease in the composite elastic modulus. Agarwal and Broutman 24 have reported that the composite elastic modulus decreases first in the presence of cross-ply cracks and then longitudinal-ply cracks are followed by delamination cracks. Therefore, with fatigue exposures, the stress/strain hysteresis of the composite
In order to compare the tension fatigue behaviour of the two composite panels, it was necessary to normalize the applied stress levels. Normalization of the data was done by taking the ratio of the repeated maximum stress level to the average ultimate tensile strength for each composite panel. The ratio was then multiplied by a factor of 100 in order to obtain the percentages. The sum was defined as percent average strength (PAS) for convenience. The variation in the number of cycles to failure as a function of PAS is plotted in Fig. 6, which also summarizes the cyclic fatigue data for these composites. Failure did not occur for stress levels less than 80 PAS. All the specimens tested below this stress level survived a f t e r 10 6 cycles. Specimens fatigue tested between 80 and 90 PAS either failed during testing or survived after 10 6 cycles. This range of PAS is considered as a critical range for cyclic fatigue of these composites. The ratio of the number of survival specimens to the total number of specimens fatigue tested in this range is
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V a r i a t i o n in t h e n u m b e r o f cycles t o failure as a f u n c t i o n percent average s t r e n g t h o f t h e composite panels
42
C O M P O S I T E S . J A N U A R Y 1992
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becomes linear with a modulus close to the secondary modulus. As seen in Fig. 8, the primary and secondary moduli are the same at 105 cycles giving a straight line in the stress/strain hysteresis. The primary modulus becomes less than the secondary modulus at 106 cycles. This has been attributed to the development of longitudinal cracks and delamination cracks due to the fatigue loading 24. The SEM analysis of fatigue tested specimens revealed a
textural change associated with increasing number of cycles. The textural changes occurred in the matrix regions, the most likely regions for origination of cracks and their extension. Two representative fractographs of the matrix regions of static tensile and fatigue tested specimens are shown in Fig. 9. The surface of the matrix of a static tensile tested specimen is smooth and appears to be glassy (Fig. 9(a)). However, the surface of the matrix of a fatigue tested specimen has a rough topography (Fig. 9(b)). The striations, which indicate the advance of crack growth after each cycle, are also observed on the fracture surfaces of a fatigue tested specimen. These striations were not detected on the fracture surfaces of the static tensile tested specimens. However, smooth and rough surfaces in the matrix regions were observed in both test specimens. In general, the surface of the matrix parallel to the fracture surface seemed to be glassy while the surfaces of the matrix normal and/or at a high angle to the fracture surface had crystalline-type character, as shown in Fig. 10. Furthermore, the surfaces normal and/or at a high angle to the fracture surface had more topography and were rougher than the surfaces parallel to the fracture surface. Although the striations were detected in the surfaces parallel to the fracture surface, they are mostly observed on the surfaces normal and/or at a high angle to the fracture surface for the fatigue tested specimens as indicated in Fig. 11. A number of fatigue tests was conducted to determine the threshold cycle number, Nt, which is defined as the number of fatigue cycles below which no textural change occurs in various microstructures. Eleven specimens of composite panel 1 were fatigue tested for various numbers of cycles ranging from 1000 to 40 000. A stress level of 80 PAS was held constant for all fatigue tests. None of the specimens failed during testing. Upon completion of the fatigue tests, specimens were broken in tension to determine the residual strengths. The residual strengths of these specimens ranged from 99.2 + 0.1 M P a t o 107.9 + 4.9 MPa. The average residual strength of the eleven specimens was 104.1 + 4.9 MPa. Resultant data show that the residual strengths of the surviving specimens were comparable to those of unfatigued specimens. However, microstructural changes were observed in
Fig. 9 SEM fractographs of carbon/carbon composites showing the matrix region of the fracture surfaces: (a) static tensile tested; (b) fatigue tested and survived after 106 cycles
Fig. 10 SEM fractograph of a carbon/carbon composite fatigue
tested and failed after 1240 cycles
C O M P O S I T E S . J A N U A R Y 1992
43
from the matrix, resulting in debonding of the fibre/ matrix interfaces and pull-out of fibres. The debonded fibre/matrix interface serves as a path for a single large crack which propagates through the cross-section area of the gauge section.
Fig. 11 SEM fractograph of a carbon/carbon composite fatigue tested and failed after 3618 cycles
the matrix region of the specimens as the number of fatigue cycles increased. The striations become more obvious on the surfaces normal to or at a high angle to the fracture surface. It was not, however, possible to determine a precise value for Nt since all the specimens had some degree of microstructural change during fatigue cycles. The development of damage under tensile fatigue loading may be explained phenomenologically based on direct observations of fatigue tests and SEM fractographs of the test specimens. Within the first 10 load cycles, several non-critical radial and circumferential microcracks form in the matrix and at the fibre/matrix interface provided that applied stress exceeds the local ply strength. The critical stress for the formation of matrix cracking has been evaluated by Aveston et al. 25 who reported that large stress and strain concentrations at the fibre/matrix interface are responsible for the initiation of these cracks. Manocha and Bahl 7 observed that these microcracks develop beneath the bundle cross-over points due to differences in thermal expansion characteristics of sharply bent fibres and the inhomogeneous matrix distribution around these crossed bundles. The cracking process may continue until the cracks in each ply have attained an equilibrium. The equilibrium state in which the matrix cracking pattern is stabilized forms by 100 cycles. The number of cracks increases with either the number of cycles or an increase in the applied stress level 26'27. As the matrix cracks develop, fibre failure begins to initiate preferentially in adjacent plies in regions of stress concentration created by the primary cracks. After initiation the crack propagates between fibres primarily along the fibre/matrix interface 28-31.The radial cracks propagate up to the fibre surface. They either stop there or deflect and branch into small secondary cracks which extend short distances away from the interface between the ply in which the primary crack has occurred and the one in which the secondary crack occurs 26. These secondary cracks are generally perpendicular to the primary cracks (or as nearly so as possible) and are caused by the tensile stresses along the crack axis ahead of the primary cracks 27. The circumferential cracks cause the separation of the fibres
44
COMPOSITES • JANUARY
1992
The microcracks perpendicular to the applied load propagate through the entire width of the ply but are unable to propagate into the adjacent ply. They terminate at the interface of two plies. However, the crack tip produces a stress concentration ahead of itself. The resulting strong interlaminar stresses in regions where primary and secondary cracks cross produce favourable conditions for starting an interlaminar crack along the ply interface. This phenomenon, known as delamination, is caused by a mixed-mode growth of interlaminar cracks 26. More delamination cracks start and propagate through the interface as the number of cycles increases. Deflection of the crack path along the boundary between fibre bundles and cloth layers is controlled by the nature of the fibre/matrix interface. If the bond strength is relatively weak, the crack propagating through the matrix deflects and branches when it reaches a boundary between the fibre bundles and cloth layers where the delamination cracks already exist. Part of the deflection is due to debonding of the fibre/matrix interface, pull-out of the fibres from the matrix and the associated increase in crack opening 32. The crack propagation path may also deviate along the planes of circumferential microcracks which are likely to behave as new interfaces 6. The deflection of the crack reduces the stress concentrations at the crack tip 6"28. In the case of interfaces with relatively strong bonds containing both fibre/matrix interfaces and planes of circumferential microcracks in the matrix, it is difficult to effect relaxation of the stress concentration of the composites around the fibre effectively because of the low degree of interface debonding 6. In this case, cracks propagate with relatively few impediments by circumferential microcracking through the matrix. The composite fractures in the primary crack plane with no significant branching or deflection. The composite undergoes final fracture when it is sufficiently weakened by delamination cracks. The presence of delamination cracks prevents load distribution between plies, and the composite is essentially reduced to a number of independent plies acting in parallel to support the applied load. The weakest of these plies fails and starts failure of the remaining plies. Development of delamination cracks is slow, and hence the loss of strength in the later part of the fatigue life is slow. Rapid loss of strength occurs in the last few cycles of fatigue life when the stronger plies fail. Prior to this rapid loss of strength, individual plies do become weakened, but the overall strength reduction is slow. At the time when delamination cracks appear, fibres within a short distance of the crack tip remain intact and bridge the crack whereas the remainder of the fibres in the crack wake start fracturing and debonding 33. Failure of the composite may advance with either matrix crack growth or fibre failure. The crack covers the whole cross-sectional area of the gauge section when the maximum stress achieved. Beyond this point
applied load is s u p p o r t e d by the intact fibres bridging the crack. R e m a i n i n g bridging fibres break due to the large stress concentrations at the surface. C o m p o s i t e failure occurs w h e n the applied stress exceeds the strength of fibre bundles. Fibre bundle failure is the f u n d a m e n t a l process that dictates the ultimate strength, and axial crack extension is then merely a c o n s e q u e n c e of fibre bundle failure 34. T h e later stage of d a m a g e d e v e l o p m e n t is a rapidly increasing rate of progression o f the mixture of all d a m a g e modes. It must be n o t e d that individual d a m a g e m o d e s do not f o r m in distinct regions but rather occur m o r e or less simultaneously. H o w e v e r , each d a m a g e m o d e m a y have a range of d o m i n a n c e .
and wave pattern on the development of 2D carbon-carbon composites' Carbon 26 No 1 (1988) pp 13-21 8 Ozturk, A. 'The mechanical properties of tightly woven carboncarbon composites' PhD thesis (University of Missouri-Rolla, Rolla, MO, USA, 1991) 9 Yasuda, E., Tanabe, Y., Manoeha, L.M. and Kumura, S.
10 11 12 13
CONCLUSIONS 1) C a r b o n / c a r b o n composites investigated show g o o d fatigue resistance to fluctuating loads. Stress to survive 106 cycles was a p p r o x i m a t e l y 80% of the ultimate tensile strength. A notable, but not significant, c h a n g e in composite residual strength and elastic m o d u l u s was o b s e r v e d as a result of fatigue loading. These results are e n c o u r a g i n g since m a n y potential applications of fibre-reinforced ceramiccomposites require a high fatigue life. 2) T e n s i o n - t e n s i o n cyclic fatigue loading caused a substantial c h a n g e in the stress/strain hysteresis loop of these composites, indicating the progressive accumulation of fatigue d a m a g e . This c h a n g e in behaviour, p r e s u m e d to be due to the o c c u r r e n c e of matrix microcracking, did not a p p e a r to have a significant effect o n composite residual strength. 3) S E M analysis of b r o k e n specimens revealed that these composites show textural changes related to fatigue. Matrix cracking and fibre bundle failure are the d o m i n a n t failure m e c h a n i s m s during fatigue. Since matrix cracks are not deleterious to the fibres, a fibre-dominant fracture is p r o n e in these c o m p o s ites. 4) This investigation d e m o n s t r a t e s that fluctuating stresses induce fatigue d a m a g e even at r o o m temp e r a t u r e in air. T h e r e f o r e , the cyclic fatigue behaviour o f a d v a n c e d structural ceramic composites needs to be e x a m i n e d if they are to be used in critical applications involving r e p e a t e d load conditions.
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32 Marshall, D.B. and Evans, A.G. 'Failure mechanisms in ceramicfiber/ceramic-matrix composites' J Amer Cer Soc 68 No 5 (1985) pp 225-231 33 Marshall, D.B. and Cox, B.N. 'Tensile fracture of brittle matrix composites: influence of fiber strength' Carbon 35 No 26 (1987) pp 2607-2619 34 Sbaizero, O. and Evans, A.G. 'Tensile and shear properties of laminated ceramic matrix composites' J A m e r Ceram Soc 69 No 6 (1986) pp 481--486
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A U THORS
A. Ozturk is with the Metallurgical Engineering Department, Middle East Technical University, Ankara, Turkey. R.E. Moore is with the Ceramic Engineering Department, University of MissouriRolla, Rolla, MO, USA.