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carbon composites prepared by chemical vapor deposition
OCW6223/88 $3.00 + .OO Copyright 0 1988 Pergamon Journals Ltd.
Carbon Vol. 26, No. 2, pp. 157-162. 1988 Printed in Great Britain.
FAILURE BEHAVIOR OF CARBON/CARBON COMPOSITES PREPARED BY CHEMICAL VAPOR DEPOSITION K. Y. SOHN, SEH-MIN OH, and JAI-YOUNG LEE Department
of Materials Science and Engineering, Korea Advanced Institute of Science and Technology, P. 0. Box 131, Cheongryang, Seoul, Korea (Received 20 Apr 1987: Accepted in revised form 25 July 1987)
Abstrad-Tensile tests have been performed on carbon/carbon composites made from carbon cloths and a pyrolytic carbon matrix and their failure surfaces were examined by scanning electron microscopy. The load-extension curves are characterized by three distinct regions: the applied load needed to maintain a constant extention rate increases initially (region I), which is maintained at an almost constant level (region II), then the load increases again until the abrupt failure of the composite occurs (region III). Cracks are initiated in region II and the composite becomes deformed plastically before failure. Circumferential microcracks formed due to the thermal anisotropic contraction of the pyrolytic carbon
matrix deviated the path of the advancing cracks and therefore increased the toughness of the composites. Fractography showed a terraced fracture surface in the matrix of the composites in which distinct circumferential
microcracks are formed in front of the advancing cracks.
Key Words-Carbon/carbon circumferential microcracks.
composites, chemical vapor deposition, carbon cloths, failure behavior,
1. INTRODUCTION
Carbon/carbon composites (C/C composites) have received considerable interest as a material for aerospace applications due to their high heat of ablation, small thermal expansion, and good strength retention at high temperatures, as well as high thermal shock resistance[l,2]. Also, due to their compatibility with blood and tissues, application of C/C composites is gradually extending to biomaterials such as heart valves and skeletons[3]. C/C composites can be produced from fibers by liquid impregnation using resins or pitch, or by chemical vapor deposition (CVD). One shortcoming of the impregnation technique is a volumetric shrinkage of the impregnant during carbonization, resulting in residual stresses that are difficult to handle, whereas in the CVD technique such stresses are smaller[4]. Several papers on C/C composites prepared by CVD are concerned with the matrix microstructure and effects of substrate and matrix fabrication on the properties of the resulting composite[5-91. There are a few review papers on the failure behavior of ceramic/ceramic (i.e., brittle fiber/brittle matrix composites and on their toughening mechanisms[lO-131). C/C composites prepared by the CVD technique are unique among such composites because the pyrolytic carbon matrix is highly anisotropic. However, little work is done on the failure behavior of the composites in spite of its importance, considering their application as structural materials. The objective of this work was, therefore, to examine the failure behavior of C/C composites prepared by CVD. The load-extension characteristics
and the fracture surfaces of the composites were investigated and the role of circumferential microcracks formed in the carbon matrix during cooling after deposition were discussed.
2. EXPERIMJTWAL The substrates used in this study were commercial* carbon cloths (Torayca cloth #6151B) made by bidirectionally weaving the polyacrylonitrile (PAN)based carbon yarns containing 1000 filaments that had an average fiber diameter of approximately 7 pm and a bulk density of 1.76 g/cm). The carbon yarn had an average strength, modulus, and ultimate elongation of 3500 MPa, 230 GPa, and 1.5%, respectively. The isothermal CVD technique[lrl] using a resistance furnace was used to infiltrate the pyrolytic carbon into the substrates. The apparatus for the deposition of pyrolytic carbon is shown schematically in Fig. 1. Propane was used as a source of carbon and argon as a carrier gas and the process of deposition was performed under atmospheric pressure. The concentration of propane was varied from 5 to 30% with the total flow rate of 0.5 to 4 Urnin. The deposition temperature was maintained at 1000°C and was continuously monitored using a type-R thermocouple. The density of the composites was measured using the following equation, which is based on the Ar-
*Toray Industries, Inc., Tokyo, Japan. 157
K. Y. SOHNet al.
158
E
t
B
B
A
A
Ar
A : Drierite B : Flow Meter C : Thermocouple D : Alumina Tube E : Resistance Furnace F : Subdrate G : Trap
CaH,
3 G
Fig. 1. Illustration of the experimental apparatus. chimedean principle:
PC=
3. RESULTS AND DISCUSSION
w*:wB x PI
where WA and W, are the weights of the composite measured in air and liquid, respectively, and p, is density of the liquid. Absolute ethyl alcohol with a density of 0.79 g/cm3 was used. Tensile tests were carried out at room temperature. The test specimens were prepared by machining the as-deposited composites. The tests were made on an Instron test machine. The dimensions of the test specimens are shown in Fig. 2a. In order to prevent breaking the specimens in the grips, epoxy resins were used to reinforce the specimens at their ends and for obvious observation of the crack initiation in the composites, tensile specimens were notched using a low speed diamond saw at a location in the specimen as shown in Fig. 2b.
Notch
(single
fiber bundle 1
(b)
Fig. 2. Illustrations of the test specimen showing (a) the dimension and (b) the location of notch.
Figure 3 shows typical microstructures of C/C composites under polarized light and bright field illuminations, respectively. The pyrolytic carbon matrix has a smooth laminar structure that displays a Maltese cross pattern in Fig. 3a. Several circumferential microcracks within the matrix, as seen in Fig. 3b, may be caused by the anisotropic thermal contraction of the pyrolytic carbon during cooling after deposition[l5]. Figure 4 depicts a typical load-extension curve for the composites in tension. The curve can be characterized by three regions: the applied load needed to maintain a constant extension rate increase initially up to a fairly high value (region I), which is maintained at an almost constant level (region II); then the load increases again until abrupt failure of the specimen occurs (region III). This specific behavior of composites may be probably ascribed to the characteristics of the carbon fiber bundles forming the carbon cloth substrate. Figure 5 shows scanning electron micrographs of the surfaces of test specimens when the load was released in regions II and III, respectively. A macroscopic crack initiated at the tip of the machined notch, propagates into the inner region of the specimen. No crack was observed in the specimen released in region I. The second region is of interest in the response of the specimen because of the occurrence of initial fracture. Figure 6 shows the load-extension curve for the specimen in which the load was released in region III. The unloading curve has almost the same length as the constant load region but a lower load value when compared to the loading curve. Moreover, the unloading curve does not return to the initial point. This means that the specimen becomes deformed plastically before failure, even if the applied load is removed. Pyrolytic carbon has lower failure strain than carbon fiber. Therefore, under loading, the matrix fails first in region II as is observed in the fractography
Carbon/carbon composites
159
(a)
(b) Fig. 3. Optical micrographs of C/C composites under (a) polarized light and (b) bright field illumination. studies, whereas the fiber is still far from its failure strain and supports the additional load up to region III. Scanning electron micrographs of various fracture surfaces are given in Fig. 7. A debonding of fiber/ matrix (F/M) interfaces and pullout of fibers from the matrix are observed in Fig. 7a, whereas the crack propagates without any debonding of F/M interfaces in Fig. 7b. Figure 7c shows the terraced propagation of the crack in the matrix, which may result from the circumferential microcracks within the matrix. It is generally known that the failure behavior of
Fig. 5. SEM photographs showing the cracks near the tip of the notch in the specimen when the load was released within region II (a) and within region III (b).
/
1 EXTENSION
EXTENSION
Fig. 4. Apical tensile behavior of the C/C composites.
Fig. 6. The load-extension curve of the specimen released in region III.
the composites is significantly affected by the properties of F/M interfaces. In composites having relatively weak bond strength at the interface, the crack propagating through the matrix will cause some debonding at the F/M interfacesI 161. Such an interface failure causes the crack to deviate from the original path. In C/C composites, there are not only FIM interfaces but also planes of circumferential microcracks parallel to the fibers in front of the advancing crack. Therefore, the propagation path of the crack can deviate along the planes of circumferential microcracks as well as at F/M interfaces. These circumferenti~ microcracks, which are likely to behave as “new interfaces,” interrupt the path of the cracks advancing through the matrix, and hence relax the stress concentration within the composites. In the
case of interfaces with relatively strong bonds containing both F/M interfaces and planes of circumferential microcracks in the matrix, it is difficult to relax the stress concentration of the composites around the fiber effectively because of little interface debonding. The cracks propagating through such a material wili have little or no deviation at the interface. In order to investigate the effects of circumferential microcracks on the mechanical properties of the composites, the number of microcracks was controlled by varying the cooling rate of the composites from the deposition tem~rat~e. Specimen A is cooled slowly to a room temperature in a furnace and specimen B is cooled more rapidly as shown in Fig. 8. The mechanical properties of both specimens are compared in Table 1. It is evident that specimen B has lower values of both tensile strength and elongation to fracture, though both specimens have similar bulk densities. Figure 9 shows both the optical micrographs under bright field illumination and the fracture surfaces of the specimens. Specimen A has more distinct circumferential microcracks in the carbon matrix than specimen B as shown in Figs. 9a and b respectively. A closer examination reveals that the microcracks with wide openings almost completely surround the fibers in specimen A, whereas narrow microcracks with short lengths are formed in specimen B. On rapid cooling, a large number of microcracks may be initiated instantaneously because of the need for rapid a~comodation of the residual thermal stresses. The microcracks do not grow to large sizes since almost all the stresses can be removed only by such crack formation. On slow cooling, however, the limited number of microcracks initiated at the beginning of cooling relaxes the stresses by further propagation or by being linked together on continued cooling, resulting in more opening and more extension around the fiber. The cracks advancing through the matrix deviate at each plane of circumferential microcracks in spec-
(b)
Fig. 7. SEM fractographies of the C/C composites showing (a) fiber/matrix interface failure, (b) little interface failure, and (c) terraced fracture surfaces.
Fig. 8. Cooling rates after deposition for two specimens; A, furnace cooling; B, air quenching in argon atmosphere.
Carbon/carbon
161
composites
Table 1. The mechanical properties of specimens A and B.
\ Cooling
No. of
“df”
R
d
E
e,
a
E
Sf
Rate
Spec.
fiber
(g/cm31
HAPa)
(GPa)
WI
Wa)
(G&d
(o/o)
A
B
S-l
15
1.91
98.85
16.94
1.30
s-2
15
1 .Ql
79.48
16.26
1.03
F-l
15
1.91
75.75
21.63
0.89
F-2
15
1.92
50.37
19.48
0.59
F-3
16
1.91
44.00
17.43
0.33
89.2
16.6
1.17
56.7
19.5
0.60
I
(b)
(d)
Fig. 9. Optical micrographs and SEM fractographs (a) and (c) specimen A (slow cooling); (b) and (d) specimen B (rapid cooling).
162
K. Y. SOHNet al.
imen A (slow cooling) and hence terraced fracture surfaces are formed in the matrix as shown in Fig. 9c. On the other hand, relatively flat fracture surfaces are observed in specimen B (rapid cooling) (Fig. 9d), which is due to an insufficient separation or debonding of the planes of microcracks to deviate the paths of the advancing crack. In this case, after initiation, a crack can propagate with relatively few impediments by circumferential microcracks through the matrix, the composites being fractured in the primary crack plane with no significant contribution to the total work of fracture. The higher value in the toughness of specimen A is, therefore, ascribed to the behavior of circumferential microcracks. Pierson and Northrop[l7] observed in carbonfelt/pyrolytic carbon-matrix composites that the strengths and moduli of composites, whose pyrolytic carbon matrix had a rough laminar structure, were higher than those with a smooth laminar structure. This may seem to contradict our results because no circumferential microcracks were found in the rough laminar structure. The influence of microcracks on C/C composite properties is thought to also depend on the fiber orientation within the substrate. Since the axes of all the fibers and thus the planes of microcracks are oriented perpendicular to the advancing crack plane in composites using a carbon cloth substrate, the microcracks are always expected to increase the composite properties as described previously. In the case of a carbon felt substrate, however, fibers have a quasirandom array and the microcrack planes that may be formed in a smooth laminar structure meet the advancing crack plane at various angles. Therefore, the microcracks perpendicular to the crack have an advantage of improving the properties whereas those parallel to it serve as a rapid path of crack propagation. If the latter effect is more predominant than the former, the presence of circumferential microcracks can have an undesirable influence on the mechanical properties of C/C composites. The decreased strength and modulus of a heattreated composite compared to one with an as-de-
posited matrix of pyrolytic carbon[5] can be similarly understood. 4. CONCLUSION The load-extension curve in tension tests of a composite of carbon cloth/pyrolytic carbon showed three distinct regions including a constant load (second) region, in which cracks were initiated. Circumferential microcracks due to an anisotropic thermal contraction of pyrolytic carbon deflected the path of the advancing cracks parallel to the fibers and, hence, resulted in the terraced fracture surface in the matrix. The circumferential microcracks also increase the toughness of the composites.
REFERENCES
1. K. R. Carnahan and R. W. Kiger, In: Advanced Materials: Composites and Carbon (Edited by J. D. Buckley), p. 171. American Ceramic Society Inc., Colombus, Ohio (1973). 2. A. J. Klein, Ad;. Mater. Prog. l30(5), 64 (1986). 3. W. Huettner and G. Keuscher. Proc. 16th Biennial Co& Carbon, American Carbon Society, p. 486. San Diego, CA (1983). 4. E. R. Frye, Nucl. Tech. 12, 93 (1971). 5. B. Granoff, H. 0. Pierson and D. M. Schuster, Carbon 11, 177 (1973). 6. B. Granoff, H. 0. Pierson and D. M. Schuster, J. Comp. Mater. 7, 36 (1973).
7. H. 0. Pierson and M. L. Lieberman, Carbon l3, 159 (1975).
8. B. Granoff, Carbon 12,681 (1974). 9. M. L. Lieberman, R. M. Curlee,.F. H. Braaten and G. T. Noles. J. Come. Mater. 9, 337 (1975). 10. I. W. Donald and P.-W. McMilian, J.‘M&r. Sci. 11, 949 (1976). 11. R. W. Rice, Ceram. Eng. Sci. Proc. 2,661 (1981). 12. R. Naslain, J. de Phys. 47, Cl-703 (1986). 13. D. B. Marshall and J. E. Ritter, Am. Ceram. Sot. Bull. 66, 309 (1987).
14. W. V. Kotlensky, In: Chemistry and Physics of Carbon (Edited by P. L. Walker, Jr.), Vol. 9, pp. 173-262. Marcel Dekker. New York (1973). 15. L. F. Coffin, Jr:, J. Am. Cekm. koc. 47,473 (1964). 16. J. Cook and J. E. Gordon, Proc. Roy. Sot. A282,508 (1964).
17. H. 0. Pierson and D. A. Northrop, J. Comp. Mater. 9, 118 (1975).
Carbon Vol. 26, No. 2, pp. 157-162. 1988 Printed in Great Britain.
FAILURE BEHAVIOR OF CARBON/CARBON COMPOSITES PREPARED BY CHEMICAL VAPOR DEPOSITION K. Y. SOHN, SEH-MIN OH, and JAI-YOUNG LEE Department
of Materials Science and Engineering, Korea Advanced Institute of Science and Technology, P. 0. Box 131, Cheongryang, Seoul, Korea (Received 20 Apr 1987: Accepted in revised form 25 July 1987)
Abstrad-Tensile tests have been performed on carbon/carbon composites made from carbon cloths and a pyrolytic carbon matrix and their failure surfaces were examined by scanning electron microscopy. The load-extension curves are characterized by three distinct regions: the applied load needed to maintain a constant extention rate increases initially (region I), which is maintained at an almost constant level (region II), then the load increases again until the abrupt failure of the composite occurs (region III). Cracks are initiated in region II and the composite becomes deformed plastically before failure. Circumferential microcracks formed due to the thermal anisotropic contraction of the pyrolytic carbon
matrix deviated the path of the advancing cracks and therefore increased the toughness of the composites. Fractography showed a terraced fracture surface in the matrix of the composites in which distinct circumferential
microcracks are formed in front of the advancing cracks.
Key Words-Carbon/carbon circumferential microcracks.
composites, chemical vapor deposition, carbon cloths, failure behavior,
1. INTRODUCTION
Carbon/carbon composites (C/C composites) have received considerable interest as a material for aerospace applications due to their high heat of ablation, small thermal expansion, and good strength retention at high temperatures, as well as high thermal shock resistance[l,2]. Also, due to their compatibility with blood and tissues, application of C/C composites is gradually extending to biomaterials such as heart valves and skeletons[3]. C/C composites can be produced from fibers by liquid impregnation using resins or pitch, or by chemical vapor deposition (CVD). One shortcoming of the impregnation technique is a volumetric shrinkage of the impregnant during carbonization, resulting in residual stresses that are difficult to handle, whereas in the CVD technique such stresses are smaller[4]. Several papers on C/C composites prepared by CVD are concerned with the matrix microstructure and effects of substrate and matrix fabrication on the properties of the resulting composite[5-91. There are a few review papers on the failure behavior of ceramic/ceramic (i.e., brittle fiber/brittle matrix composites and on their toughening mechanisms[lO-131). C/C composites prepared by the CVD technique are unique among such composites because the pyrolytic carbon matrix is highly anisotropic. However, little work is done on the failure behavior of the composites in spite of its importance, considering their application as structural materials. The objective of this work was, therefore, to examine the failure behavior of C/C composites prepared by CVD. The load-extension characteristics
and the fracture surfaces of the composites were investigated and the role of circumferential microcracks formed in the carbon matrix during cooling after deposition were discussed.
2. EXPERIMJTWAL The substrates used in this study were commercial* carbon cloths (Torayca cloth #6151B) made by bidirectionally weaving the polyacrylonitrile (PAN)based carbon yarns containing 1000 filaments that had an average fiber diameter of approximately 7 pm and a bulk density of 1.76 g/cm). The carbon yarn had an average strength, modulus, and ultimate elongation of 3500 MPa, 230 GPa, and 1.5%, respectively. The isothermal CVD technique[lrl] using a resistance furnace was used to infiltrate the pyrolytic carbon into the substrates. The apparatus for the deposition of pyrolytic carbon is shown schematically in Fig. 1. Propane was used as a source of carbon and argon as a carrier gas and the process of deposition was performed under atmospheric pressure. The concentration of propane was varied from 5 to 30% with the total flow rate of 0.5 to 4 Urnin. The deposition temperature was maintained at 1000°C and was continuously monitored using a type-R thermocouple. The density of the composites was measured using the following equation, which is based on the Ar-
*Toray Industries, Inc., Tokyo, Japan. 157
K. Y. SOHNet al.
158
E
t
B
B
A
A
Ar
A : Drierite B : Flow Meter C : Thermocouple D : Alumina Tube E : Resistance Furnace F : Subdrate G : Trap
CaH,
3 G
Fig. 1. Illustration of the experimental apparatus. chimedean principle:
PC=
3. RESULTS AND DISCUSSION
w*:wB x PI
where WA and W, are the weights of the composite measured in air and liquid, respectively, and p, is density of the liquid. Absolute ethyl alcohol with a density of 0.79 g/cm3 was used. Tensile tests were carried out at room temperature. The test specimens were prepared by machining the as-deposited composites. The tests were made on an Instron test machine. The dimensions of the test specimens are shown in Fig. 2a. In order to prevent breaking the specimens in the grips, epoxy resins were used to reinforce the specimens at their ends and for obvious observation of the crack initiation in the composites, tensile specimens were notched using a low speed diamond saw at a location in the specimen as shown in Fig. 2b.
Notch
(single
fiber bundle 1
(b)
Fig. 2. Illustrations of the test specimen showing (a) the dimension and (b) the location of notch.
Figure 3 shows typical microstructures of C/C composites under polarized light and bright field illuminations, respectively. The pyrolytic carbon matrix has a smooth laminar structure that displays a Maltese cross pattern in Fig. 3a. Several circumferential microcracks within the matrix, as seen in Fig. 3b, may be caused by the anisotropic thermal contraction of the pyrolytic carbon during cooling after deposition[l5]. Figure 4 depicts a typical load-extension curve for the composites in tension. The curve can be characterized by three regions: the applied load needed to maintain a constant extension rate increase initially up to a fairly high value (region I), which is maintained at an almost constant level (region II); then the load increases again until abrupt failure of the specimen occurs (region III). This specific behavior of composites may be probably ascribed to the characteristics of the carbon fiber bundles forming the carbon cloth substrate. Figure 5 shows scanning electron micrographs of the surfaces of test specimens when the load was released in regions II and III, respectively. A macroscopic crack initiated at the tip of the machined notch, propagates into the inner region of the specimen. No crack was observed in the specimen released in region I. The second region is of interest in the response of the specimen because of the occurrence of initial fracture. Figure 6 shows the load-extension curve for the specimen in which the load was released in region III. The unloading curve has almost the same length as the constant load region but a lower load value when compared to the loading curve. Moreover, the unloading curve does not return to the initial point. This means that the specimen becomes deformed plastically before failure, even if the applied load is removed. Pyrolytic carbon has lower failure strain than carbon fiber. Therefore, under loading, the matrix fails first in region II as is observed in the fractography
Carbon/carbon composites
159
(a)
(b) Fig. 3. Optical micrographs of C/C composites under (a) polarized light and (b) bright field illumination. studies, whereas the fiber is still far from its failure strain and supports the additional load up to region III. Scanning electron micrographs of various fracture surfaces are given in Fig. 7. A debonding of fiber/ matrix (F/M) interfaces and pullout of fibers from the matrix are observed in Fig. 7a, whereas the crack propagates without any debonding of F/M interfaces in Fig. 7b. Figure 7c shows the terraced propagation of the crack in the matrix, which may result from the circumferential microcracks within the matrix. It is generally known that the failure behavior of
Fig. 5. SEM photographs showing the cracks near the tip of the notch in the specimen when the load was released within region II (a) and within region III (b).
/
1 EXTENSION
EXTENSION
Fig. 4. Apical tensile behavior of the C/C composites.
Fig. 6. The load-extension curve of the specimen released in region III.
the composites is significantly affected by the properties of F/M interfaces. In composites having relatively weak bond strength at the interface, the crack propagating through the matrix will cause some debonding at the F/M interfacesI 161. Such an interface failure causes the crack to deviate from the original path. In C/C composites, there are not only FIM interfaces but also planes of circumferential microcracks parallel to the fibers in front of the advancing crack. Therefore, the propagation path of the crack can deviate along the planes of circumferential microcracks as well as at F/M interfaces. These circumferenti~ microcracks, which are likely to behave as “new interfaces,” interrupt the path of the cracks advancing through the matrix, and hence relax the stress concentration within the composites. In the
case of interfaces with relatively strong bonds containing both F/M interfaces and planes of circumferential microcracks in the matrix, it is difficult to relax the stress concentration of the composites around the fiber effectively because of little interface debonding. The cracks propagating through such a material wili have little or no deviation at the interface. In order to investigate the effects of circumferential microcracks on the mechanical properties of the composites, the number of microcracks was controlled by varying the cooling rate of the composites from the deposition tem~rat~e. Specimen A is cooled slowly to a room temperature in a furnace and specimen B is cooled more rapidly as shown in Fig. 8. The mechanical properties of both specimens are compared in Table 1. It is evident that specimen B has lower values of both tensile strength and elongation to fracture, though both specimens have similar bulk densities. Figure 9 shows both the optical micrographs under bright field illumination and the fracture surfaces of the specimens. Specimen A has more distinct circumferential microcracks in the carbon matrix than specimen B as shown in Figs. 9a and b respectively. A closer examination reveals that the microcracks with wide openings almost completely surround the fibers in specimen A, whereas narrow microcracks with short lengths are formed in specimen B. On rapid cooling, a large number of microcracks may be initiated instantaneously because of the need for rapid a~comodation of the residual thermal stresses. The microcracks do not grow to large sizes since almost all the stresses can be removed only by such crack formation. On slow cooling, however, the limited number of microcracks initiated at the beginning of cooling relaxes the stresses by further propagation or by being linked together on continued cooling, resulting in more opening and more extension around the fiber. The cracks advancing through the matrix deviate at each plane of circumferential microcracks in spec-
(b)
Fig. 7. SEM fractographies of the C/C composites showing (a) fiber/matrix interface failure, (b) little interface failure, and (c) terraced fracture surfaces.
Fig. 8. Cooling rates after deposition for two specimens; A, furnace cooling; B, air quenching in argon atmosphere.
Carbon/carbon
161
composites
Table 1. The mechanical properties of specimens A and B.
\ Cooling
No. of
“df”
R
d
E
e,
a
E
Sf
Rate
Spec.
fiber
(g/cm31
HAPa)
(GPa)
WI
Wa)
(G&d
(o/o)
A
B
S-l
15
1.91
98.85
16.94
1.30
s-2
15
1 .Ql
79.48
16.26
1.03
F-l
15
1.91
75.75
21.63
0.89
F-2
15
1.92
50.37
19.48
0.59
F-3
16
1.91
44.00
17.43
0.33
89.2
16.6
1.17
56.7
19.5
0.60
I
(b)
(d)
Fig. 9. Optical micrographs and SEM fractographs (a) and (c) specimen A (slow cooling); (b) and (d) specimen B (rapid cooling).
162
K. Y. SOHNet al.
imen A (slow cooling) and hence terraced fracture surfaces are formed in the matrix as shown in Fig. 9c. On the other hand, relatively flat fracture surfaces are observed in specimen B (rapid cooling) (Fig. 9d), which is due to an insufficient separation or debonding of the planes of microcracks to deviate the paths of the advancing crack. In this case, after initiation, a crack can propagate with relatively few impediments by circumferential microcracks through the matrix, the composites being fractured in the primary crack plane with no significant contribution to the total work of fracture. The higher value in the toughness of specimen A is, therefore, ascribed to the behavior of circumferential microcracks. Pierson and Northrop[l7] observed in carbonfelt/pyrolytic carbon-matrix composites that the strengths and moduli of composites, whose pyrolytic carbon matrix had a rough laminar structure, were higher than those with a smooth laminar structure. This may seem to contradict our results because no circumferential microcracks were found in the rough laminar structure. The influence of microcracks on C/C composite properties is thought to also depend on the fiber orientation within the substrate. Since the axes of all the fibers and thus the planes of microcracks are oriented perpendicular to the advancing crack plane in composites using a carbon cloth substrate, the microcracks are always expected to increase the composite properties as described previously. In the case of a carbon felt substrate, however, fibers have a quasirandom array and the microcrack planes that may be formed in a smooth laminar structure meet the advancing crack plane at various angles. Therefore, the microcracks perpendicular to the crack have an advantage of improving the properties whereas those parallel to it serve as a rapid path of crack propagation. If the latter effect is more predominant than the former, the presence of circumferential microcracks can have an undesirable influence on the mechanical properties of C/C composites. The decreased strength and modulus of a heattreated composite compared to one with an as-de-
posited matrix of pyrolytic carbon[5] can be similarly understood. 4. CONCLUSION The load-extension curve in tension tests of a composite of carbon cloth/pyrolytic carbon showed three distinct regions including a constant load (second) region, in which cracks were initiated. Circumferential microcracks due to an anisotropic thermal contraction of pyrolytic carbon deflected the path of the advancing cracks parallel to the fibers and, hence, resulted in the terraced fracture surface in the matrix. The circumferential microcracks also increase the toughness of the composites.
REFERENCES
1. K. R. Carnahan and R. W. Kiger, In: Advanced Materials: Composites and Carbon (Edited by J. D. Buckley), p. 171. American Ceramic Society Inc., Colombus, Ohio (1973). 2. A. J. Klein, Ad;. Mater. Prog. l30(5), 64 (1986). 3. W. Huettner and G. Keuscher. Proc. 16th Biennial Co& Carbon, American Carbon Society, p. 486. San Diego, CA (1983). 4. E. R. Frye, Nucl. Tech. 12, 93 (1971). 5. B. Granoff, H. 0. Pierson and D. M. Schuster, Carbon 11, 177 (1973). 6. B. Granoff, H. 0. Pierson and D. M. Schuster, J. Comp. Mater. 7, 36 (1973).
7. H. 0. Pierson and M. L. Lieberman, Carbon l3, 159 (1975).
8. B. Granoff, Carbon 12,681 (1974). 9. M. L. Lieberman, R. M. Curlee,.F. H. Braaten and G. T. Noles. J. Come. Mater. 9, 337 (1975). 10. I. W. Donald and P.-W. McMilian, J.‘M&r. Sci. 11, 949 (1976). 11. R. W. Rice, Ceram. Eng. Sci. Proc. 2,661 (1981). 12. R. Naslain, J. de Phys. 47, Cl-703 (1986). 13. D. B. Marshall and J. E. Ritter, Am. Ceram. Sot. Bull. 66, 309 (1987).
14. W. V. Kotlensky, In: Chemistry and Physics of Carbon (Edited by P. L. Walker, Jr.), Vol. 9, pp. 173-262. Marcel Dekker. New York (1973). 15. L. F. Coffin, Jr:, J. Am. Cekm. koc. 47,473 (1964). 16. J. Cook and J. E. Gordon, Proc. Roy. Sot. A282,508 (1964).
17. H. 0. Pierson and D. A. Northrop, J. Comp. Mater. 9, 118 (1975).