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carbon composite under uniaxial tensile and shear loads
PII:
Acta mater. Vol. 46, No. 18, pp. 6631±6643, 1998 # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain S1359-6454(98)00269-9 1359-6454/98 $19.00 + 0.00
DAMAGE AND FAILURE MECHANISMS OF A 3-DIRECTIONAL CARBON/CARBON COMPOSITE UNDER UNIAXIAL TENSILE AND SHEAR LOADS O. SIRON and J. LAMON{ Laboratoire des Composites Thermostructuraux, UMR 5801 (CNRS-SEP-UB1), 3 AlleÂe de La BoeÂtie, 33600 Pessac, France (Received 5 May 1997; accepted 22 July 1998) AbstractÐThe mechanical behavior of a three-directional carbon/carbon (C/C) composite under tensile and shear loads is investigated in relation with the failure mechanisms and, the ®ber architecture. This three-directional C/C composite was produced by Chemical Vapor In®ltration of a needled ®ber preform of multiple layers of satin woven tows. The C/C composite exhibited several interesting features including an essentially non-linear stress±strain behavior and permanent deformations. Three families of matrix cracks were identi®ed under tensile and shear loads, including microcracks in the tows, intertow delamination and cracks across the longitudinal tows. It was found that the delamination cracks aect preponderantly the stress±strain behavior and the mechanical properties. Similar features in the mechanical behavior and the failure mechanisms were highlighted under tension and under shear loading. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. ReÂsumeÂÐLe comportement meÂcanique d'un composite Carbone/Carbone 3D sous chargement de traction et de cisaillement est eÂtudieÂ. Ce mateÂriau composite C/C a eÂte obtenu par densi®cation par C.V.I. (Chemical Vapor In®ltration) d'une preÂforme ®breuse aiguilleteÂe normalement au plan du tissu. Les relations entre les courbes de comportement meÂcanique, les meÂcanismes de deÂgradation identi®eÂs par microscopie optique et l'architecture ®breuse de la structure ont eÂte eÂtablies. Ce mateÂriau composite C/C preÂsente un comportement meÂcanique non lineÂaire jusqu'aÁ rupture avec deÂformations aneÂlastiques. Les trois meÃmes familles de ®ssures matricielles ont eÂte identi®eÂes sous chargement de traction et de cisaillement, aÁ savoir la ®ssuration intra-®l, la ®ssuration inter-®ls, et, la ®ssuration intra et inter-®ls. La ®ssuration inter-®l, et, la ®ssuration intra et inter-®ls aectent majoritairement le comportement meÂcanique. Les similitudes importantes observeÂes entre le comportement meÂcanique et les meÂcanismes de rupture sous sollicitations de traction et de cisaillement sont discuteÂes. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.
1. INTRODUCTION
Carbon/carbon (C/C) composites display several advantageous properties for structural applications at high temperatures up to 30008C, including: a low density, a good strength retention at high temperatures, a high thermal and chemical stability in inert environments, a high thermal shock resistance, a high resistance to wear and good friction properties at high temperature. Because of these superior thermal and mechanical properties which persist at very high temperatures, C/C composites are used in many areas including disc brakes, wing edges for missiles and space shuttle, solid propulsion applications, furnace tools, fasteners, biomedical, nuclear reactors, etc. The C/C composites are manufactured by in®ltration of the carbon matrix into the carbon ®ber preform using either a liquid or a gaseous route. Unlike the C/C composites manufactured by liquid {To whom all correspondence should be addressed.
impregnation of the matrix [1±9] or the textile composites with a polymer matrix [10], the mechanical behavior of the CVI±C/C (Chemical Vapor In®ltration (CVI) [11]) composites has not been subject of extensive studies. The main trends in the stress±strain behavior and the failure mechanisms that were reported in the literature for those C/C composites obtained by liquid impregnation, may be summarized as follows. The stress±strain behavior depends on the composite structure. The laminated and the braided C/C composites generally exhibit a linear stress±strain behavior under tension or compression, when the load is applied parallel to a ®ber tow direction [12]. By contrast, the stress± strain behavior of the two-dimensional woven composites is essentially non-linear. The non-linear behavior is attributed to matrix and ®ber damage. Some investigations of two-dimensional C/C laminates in which damage was observed during tensile loading, indicated that damage occurred ®rst at the locations of the interface where the woven yarns crossed one another [8]. The failure mechanisms
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SIRON and LAMON: CARBON/CARBON COMPOSITE
Fig. 1. 8-H satin weave pattern and tow directions.
under tension identi®ed in a C/C composite with a two-dimensional satin weave reinforcement have been grouped into three families [2]: (i) microcracks located in the transverse tows, (ii) cracks in the transverse tows, and (iii) interply delamination. It
was recently shown that in a satin weave C/C laminate, delamination produces permanent change of the stress±strain behavior [9]. Very few studies have been directed to the local mechanisms of failure under shear loading [4 and references therein]. Several types of two-dimensional laminated C±C composites under shear loading showed crack initiation at the crimp edges situated on the interply interface [4]. The Chemical Vapor In®ltration (CVI)±C/C composites have been reported to dier from those manufactured by liquid impregnation in that the pyrocarbon matrix is characterized by highly anisotropic properties and lower residual stresses [13]. In the present paper, the mechanical behavior of a three-directional CVI±C/C composite under tensile and shear loading conditions at room temperature, is investigated in relation to the damages observed during the tests.
Fig. 2. Crimp angle de®nition (a) and distribution (b).
SIRON and LAMON: CARBON/CARBON COMPOSITE
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2. EXPERIMENTAL PROCEDURE
2.1. Material The three-directional C/C composite is a needled laminate reinforced with ex-PAN C ®bers and manufactured by CVI (by SEP, Le Haillan, France). The ®ber architecture consists of needled layers of 8-harness satin weave. The weave structure is depicted in Fig. 1. In the plies [planes (1, 2)], the weft tows (parallel to direction 1) are nearly straight, whereas the warp tows (parallel to direction 2) are locally crimped. The average crimp angle is around 288 (Fig. 2). Fibers were randomly needled perpendicular to the plies (i.e. parallel to direction 3). The volume fraction of ®bers along the directions 1, 2 and 3 is respectively 0.11, 0.11 and 0.03. The pyrocarbon matrix is rough laminar-type (i.e. highly anisotropic) [14]. The total porosity was measured by helium pycnometry on powdered samples: 14.7 2 0.8% (7 measurements). The apparent density of the specimens is 0.93 2 0.05. As discussed in a following section, intra- and inter-tow cracks were present in the as-received samples. 2.2. Uniaxial on-axis tensile tests Twenty one tensile test specimens were cut from a single part so as to make either the weft bundles (direction 1) or the warp bundles (direction 2) the loading axis. The specimen dimensions are given in Fig. 3. The front face of eleven samples (i.e. the face parallel to the plies, as opposed to the lateral faces which are perpendicular to the plies) was polished for microscope examination. The test spe-
Fig. 4. Schematic diagram showing the loading mode of the Iosipescu shear test specimen and the front face which was examined during the tests.
cimen ends were stuck to aluminium pieces that were gripped in the load train of the tensile machine. Special attention was paid to specimen alignment which was ensured on an aluminium test specimen. The specimens were loaded at a strain rate of 0.05% minÿ1 by means of a closed-loop servo hydraulic testing machine. Longitudinal deformations were measured using extensometers mounted respectively on the front and the lateral faces (gauge length: 20 mm). Unloading±reloading cycles were also carried out. Acoustic emission was recorded during the tests. Calibration of the recording system permitted selection of the appropriate parameters to eliminate extraneous noise from the testing device. The polished front surface was observed under an optical microscope before the tests, during the tests and after the ultimate failure. The examination of the specimens under load used a high magni®cation microscope (up to 416) built on a stage with three displacement axes and coupled with a CCD camera, a video recorder and a control monitor. The same region in the gauge area was inspected at various loads according to a grid of 50±70 views, covering roughly one-®fth of the total gauge area. The load was kept constant during specimen examination. The load steps correspond to 0.05% deformation increments. An alternative method was used to observe through thickness failure modes. For this purpose, a mold was mounted on the specimens. Under a load close to the ultimate failure, the specimens were vacuum impregnated with an epoxy resin to ®ll in the open cracks. The load was kept constant during resin polymerization. Further, the specimens were sectioned and polished parallel to the three reference planes de®ned by the directions 1, 2 and 3, and examined using an optical microscope. This method was applied to two specimens. 2.3. Intralaminar Iosipescu shear tests
Fig. 3. Dimensions (mm) of the test specimens: (a) tensile specimens, (b) specimens for Iosipescu shear tests.
The Iosipescu test specimens were cut out of the previously mentioned part, so as to make either the weft bundle (direction 1) or the warp bundles (direction 2) the loading axis. The specimen dimensions are given in Fig. 3. Figure 4 shows the front face and the orientation of the plies with respect to the loading axis. The tests were performed under
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SIRON and LAMON: CARBON/CARBON COMPOSITE
Fig. 5. Typical stress±strain curves obtained during the tensile tests under: (a) a load parallel to the direction 1 of tows; and (b) a load parallel to the direction 2 of tows.
displacement control at a strain rate of 0.05 mm minÿ1. The deformations were measured on the rear face by means of a 2458 strain gauge (Vishay CEA06-125-WT-350). The front face was polished for microscope examination of the area located between the ``V'' notches (Fig. 4). The two previously described methods of damage observation were used. About 70 micrographs covering roughly one-quarter of the ``V'' notch area were recorded. The load was kept constant during surface examination. The load was increased by 5 MPa increments. The resin impregnation technique was applied to one sample under a load close to ultimate failure. The shear tests were interrupted at a stress of 145 MPa. 3. EXPERIMENTAL RESULTS
3.1. Stress/strain behavior 3.1.1. On-axis tensile tests. Figure 5 shows typical stress±strain curves which evidence several features. First, the three-directional C/C composite is essentially non-linear, from very low loads, whatever the loading direction. No signi®cant initial linear domain was identi®ed. Second, the strain at a given stress is much larger (roughly two times), when the direction 2 is aligned with the tensile axis. Third, the unloading±reloading hysteresis loops are very narrow. Fourth, signi®cant permanent deformations are observed after unloading. Table 1 summarizes the results of the on-axis tensile tests. Young's modulus was determined from the slope of the tangent to the curve at the origin. It is around 30 GPa, and the strain-to-failure is around 0.5%. The higher Young's modulus and failure stress are obtained in direction 1, which suggests that the crimped tows exert a certain in¯uence on the mechanical properties. Comparison of
the Poisson's ratios indicates larger lateral contractions (about two times) when the load is applied parallel to the direction 2. This feature also indicates an eect of the presence of crimped tows in this direction. Figure 6 describes stiness changes during the tensile tests. The stiness was determined from the slope of the unloading±reloading hysteresis loops. This slope provides a satisfactory approximation of the composite stiness, since the loops width was generally rather small (Fig. 5). It may be noticed from Fig. 6 that the stiness seems to increase ®rst. Such stiening must not be attributed to an artefact but instead it is related to the pre-existing matrix cracks, as discussed in a following section. Acoustic emission was detected from 0.05% applied deformations (Fig. 5). It develops exponentially from 0.2% deformations, so that most of the acoustic emission is recorded close to the ultimate failure. Figure 5 also shows that no acoustic emission is observed during the unloading±reloading cycles suggesting the occurrence of limited damage phenomena. Comparable numbers of counts are obtained whatever the loading direction. Figure 7 shows a linear relationship between the residual strains at zero load and the degree of damage measured by a modulus loss [1-(Ei/E0i ), Ei=composite stiness, E0i =initial Young's modulus along the loading direction i(i = 1,2)]. Beyond Table 1. Main mechanical properties of the investigated C/C composite determined under on-axis tensile tests E0i (GPa) (i = 1,2)
n0
sr (MPa)
er (%)
1
33 22
126.6 24.6
0.482 0.06
2
22.2 20.5
n012=0.18 20.01 n013=0.21 20.02 n023=0.39 20.02
64.5 24.9
0.472 0.07
Direction of loading
SIRON and LAMON: CARBON/CARBON COMPOSITE
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Fig. 6. Stiness changes vs longitudinal strains during the on-axis tensile tests under loads parallel to direction 1 or to direction 2.
a threshold of around 0.02%, the residual strains seem to be proportional to the damage parameter, and independent of the loading direction. Below the threshold, the residual strains are not (or little) related to the damage parameter suggesting that they do not result from a stress-induced damage. The Poisson's ratios shown by Fig. 8 re¯ect a decrease in the lateral contractions as matrix cracking proceeds. This trend is generally observed with ceramic matrix composites (CMCs). It is attributed to cracking parallel to the loading direction (typically ®ber debonding in CMCs). Figure 8 also shows that the Poisson's ratio dependence on the loading direction becomes more pronounced during the tests. This indicates that the lateral contractions in direction 3, are larger by a
factor >2 when the load is applied parallel to direction 2, relative to those obtained when the load is applied parallel to direction 1. 3.1.2. Iosipescu in-plane shear tests. Figure 9 shows a typical stress±strain curve from the Iosipescu shear tests. The stress was derived from the ratio of the applied force to the cross section area between the notches. The mechanical properties did not need to be corrected, since the possible eect of ply anisotropy can be considered to be low: E01/ E0211.6. As a matter of fact, for this degree of anisotropy, the correction factor [15] is smaller than 5% for the shear modulus. The stress±strain curve exhibits features similar to those observed under the on-axis tensile loading conditions. First, the stress±strain relationship is essentially non-linear. No initial domain of linear behavior could be identi®ed. Second, the unload-
Fig. 7. Evolution of residual strains vs a damage parameter 1 ÿ (Ei/E0i ), i = 1,2.
Fig. 8. Poisson coecient changes, along direction 2 or 3 under tensile loads applied parallel to direction 1 or 2.
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SIRON and LAMON: CARBON/CARBON COMPOSITE
Fig. 9. Typical stress±strain curve obtained during the Iosipescu shear tests.
ing±reloading hysteresis loops are ®rst rather narrow. Then they become wider under increasing loads, suggesting the presence of friction phenomena near the ultimate failure. Third, permanent deformations are observed after unloading. Nevertheless, it is worth mentioning that the stress± strain curves did not display any dependence upon the loading direction. The shear modulus that was derived from the tangent to the stress±strain curves at the origin, is obtained to be G012=5.8 2 0.2 GPa. Figure 10 shows the shear modulus degradation determined from the slopes of the unloading±reloading hysteresis loops. As opposed to the stiness which decreased by <25% under tension, the shear modulus loss reaches about 70%. Acoustic emission was detected at a 0.05% shear strain (Fig. 9). The number of counts is quite proportional to the shear strain. No acoustic emission was observed during the unloading±reloading cycles, suggesting the occurrence of limited damage. As previously for the tensile tests, a linear relationship is obtained between the residual strains and the degree of damage (Fig. 11) measured using the shear modulus loss: 1 ÿ (G12/G012). This relationship indicates that the residual strains are related to
the stress-induced damage. The presence of a threshold is not obvious (<0.01%).
Fig. 10. Shear modulus changes vs shear stress during the Iosipescu shear test.
Fig. 11. Evolution of residual shear deformations e12p vs a damage parameter 1 ÿ (G12/G012).
3.2. Microscope observations of matrix damage 3.2.1. As-received samples. Two families of preexisting cracks were identi®ed in the as-received specimens (Fig. 12): (i) intra yarn microcracks located in the transverse tows, in the matrix, or in the ®bermatrix interface, parallel to the ®ber direction [referred to as a-cracks: Fig. 12(a)]; and (ii) inter yarns microcracks located between two perpendicular tows [referred to as b-cracks; Fig. 12(b)]. The microcracks formed during composite processing at high temperatures and the subsequent cooling down cycles. They result from thermally-induced stresses that arise because of a thermal expansion mismatch between ®bers and matrix (for the a-cracks involving ®ber-matrix debonding) and also between perpendicular tows (for the b-cracks). However, reliable data on the thermal expansion coecients of the CVI±pyrocarbon matrix and the C ®bers are not available yet. Estimates of the density of pre-existing microcracks were derived from the number of micro-
SIRON and LAMON: CARBON/CARBON COMPOSITE
6637
Fig. 12. Micrographs showing the pre-existing intra-yarn microcracks (a) and the inter-yarn microcracks (b) detected in the as-received C/C specimens.
cracks detected in the specimen area inspected with the microscope. The average densities are respectively 3.40 2 1.30 mmÿ2 for the a-cracks, and 0.45 2 0.27 mmÿ2 for the b-cracks. 3.2.2. Specimens tested under on-axis tension. The microcracks and cracks that have been detected in the specimens during the tests may be grouped into three families on the basis of their location, whatever the loading direction (1 or 2): (i) a-microcracks (Fig. 13) parallel and perpendicular to the loading direction. They are dictated by the high anisotropy of pyrocarbon (i.e. the orientation of the graphitic layers) and by the presence of pores. Besides, the pre-existing a-micro-
cracks did not generally propagate during the tensile tests. (ii) b-cracks (Figs 13 and 14) parallel and perpendicular to the loading direction. (iii) c-cracks (Figs 13 and 14) consisting of matrix cracks that extend ®rst between a longitudinal tow and either a transverse tow or the intertow matrix, and then across a longitudinal tow. They are then preferentially oriented perpendicular to the loading direction. However, cracks parallel to the loading direction as well as cracks that deviated at 908, were also observed. Depending on the applied load, the c-cracks were either shorter than two tow
Fig. 13. Schematic diagram summarizing the three families of cracks detected during the tensile and the Iosipescu shear tests.
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SIRON and LAMON: CARBON/CARBON COMPOSITE
Fig. 14. Optical micrographs showing the b- and c-cracks created in the matrix during the tensile tests.
diameters (c1-cracks), or as long as two or three tow diameters (c2-cracks). Table 2 shows that the a-, b- and c-cracks formed in chronological order. The a-cracks initiated ®rst
at a 0.05% applied deformation, whatever the loading direction. This deformation coincides with the onset of acoustic emission. Then, the b-cracks, and ®nally the c-cracks were observed at higher defor-
Table 2. Matrix cracking stresses and strains determined during on-axis tensile tests Direction of loading 1 2
strain stress strain stress
a-crack
b-crack
c1-crack
c2-crack
0.05±0.1% 16±30 MPa 0.05±0.1% 10±20 MPa
0.1±0.2% 30±60 MPa 0.05±0.15% 10±30 MPa
0.2±0.3% 60±90 MPa 0.1±0.2% 20±40 MPa
0.3±0.35% 90±100 MPa 0.2±0.3% 40±50/55 MPa
SIRON and LAMON: CARBON/CARBON COMPOSITE Table 3. Mean crack densities and sizes measured after on-axis tensile tests Crack family
Mean density (mmÿ2)
Mean length (mm)
1.00 0.50 0.55 0.25 0.15 0.05
0.14 0.20 0.20 0.35 0.35 1.00
a (perpendicular)* a (parallel)* b c1 (perpendicular)* c1 (parallel)* c2 (perpendicular)* *To loading direction
mations (Table 2). The matrix cracking stresses and deformations were generally smaller for those specimens loaded parallel to the direction 2, suggesting an eect of the crimped tows. Saturation of matrix cracking was not identi®ed. The crack densities increased steadily with the applied load. The three families of cracks were distributed homogeneously. Table 3 summarizes the crack characteristics obtained from the examination of the surface and the interior of the specimens. Further details are provided in [16]. The crack densities are generally <1 crack/mm2. They must be regarded as small when comparing with the crack densities observed in ceramic matrix composites including the SiC/SiC composites [17]. Data on the crack lengths indicate that the a- and b-cracks are localized failures. The c-cracks are the longest and include one up to two or three yarn diameters. 3.2.3. Specimens tested under shear. The cracks created during the Iosipescu shear tests pertain to the three previously described families of a-, b- and c-matrix cracks (Fig. 13). No additional damage modes nor crack types were observed. Table 4 summarizes the matrix cracking stresses and deformations. Data in parentheses refer to the extension of the pre-existing cracks since contrary to the on-axis tensile tests, the pre-existing a- and b-cracks propagated. ``Mixed'' designates the cracks which were de¯ected to a perpendicular direction. As previously, the three families of cracks may be regarded as independent. The a- and b-cracks initiated under 20±25 MPa whatever their orientation. The c-cracks initiated under 25 MPa (when parallel to the loading direction) and around 35±40 MPa (when perpendicular to the loading direction). As previously, the matrix cracks were homogeneously
6639
Table 5. Mean crack densities (mmÿ2) measured after a maximum applied stress of 45 MPa under Iosipescu shear tests Crack orientation* Perpendicular Parallel Mixed
a-crack
b-crack
c-crack
0.53 20.40 1.43 20.44 0.30 20.07
0.252 0.13 0.592 0.42 Ð
0.092 0.06 0.282 0.17 0.192 0.14
*To loading direction
distributed, and the notch tips did not appear to be preferential sites of cracking. The densities of cracks that initiated in the surface of specimens during the Iosipescu tests (Table 5), are larger than those determined during the tensile tests (Table 3). Besides, comparable densities were generally obtained for both the cracks perpendicular or parallel to the tow direction at stresses <15 MPa. The parallel cracks were preponderant at stresses >15 MPa. Comparable data were determined for the cracks detected in the interior of specimens [16]. The lengths of the a-, b- and c1-cracks are close to those measured in the tensile specimens. They are, respectively, 0.16 2 0.05 mm, 0.19 2 0.09 mm and 0.37 2 0.08 mm. The length of the c2-cracks was dicult to determine because of crack deviation. A mean length of about 1 mm was estimated. 4. DISCUSSION
The stress±strain behavior was found to exhibit several features such as a non-linearity even at very low loads, residual deformations at zero load and a signi®cant anisotropy. Let us examine the contribution of various factors including the pre-existing and the stress-induced damage, the presence of crimped tows and the matrix anisotropy. 4.1. In¯uence of damage Figure 15 shows the in¯uence of stress-induced matrix cracks on the stiness. Under tension, the acracks do not in¯uence signi®cantly the stiness whereas a stiness loss can be noticed when the bcracks appear whatever the loading direction. The c-cracks which also involve a localized intertow delamination contribute to the stiness degradation. However, their transverse extension across the matrix in longitudinal tows may not be expected to
Table 4. Matrix cracking stresses and strains determined during the Iosipescu shear tests (data in parentheses refer to the extension of the pre-existing cracks) Crack orientation* Perpendicular Parallel Mixed *To loading direction
a-crack
b-crack
c-crack
(10±20) 25±45 MPa (0.1±0.2) 0.3±1.2% (10±20) 25±45 MPa (0.1±0.2) 0.2±1.2% (0±15) 25±45 MPa (0±0.14) 0.3±1.2%
(0±20) 20±45 MPa (0±0.2) 0.2±1.2% (0±20) 20±45 MPa (0±0.2) 0.2±1.2% (30±40 MPa) (0.6±0.87%)
(20±35) 40±45 MPa (0.2±0.6) 0.9±1.2% (10±20) 25±45 MPa (0.1±0.2) 0.3±1.2% (15±35) 35±45 MPa (0.14±0.6) 0.6±1.2%
6640
SIRON and LAMON: CARBON/CARBON COMPOSITE
Fig. 15. Stiness (a) and shear modulus (b) changes as a function of the failure mechanisms. n, normal to the axis of loading; p, parallel to the axis of loading; m, mixed (i.e. parallel and normal to the axis of loading); dv, development of a damage mode; pr, extension of a pre-existing damage mode.
aect tremendously the composite stiness, because of the limited load carrying capacity of the pyrocarbon matrix, which exhibits a low longitudinal elastic modulus (Em130 GPa) relative to that of ®bers (Ef1200 GPa). Therefore, transverse cracking of the pyrocarbon matrix does not increase the load on the ®bers nor the associated deformations. Furthermore, poor load transfers must be expected through the ®ber/matrix interfaces. As a matter of fact, Fig. 16 shows that ®bers have been pulled in. Under shear load, a shear modulus loss [Fig. 15(b)] coincides with onset of matrix cracking. Most of the stress-induced damage during the shear tests results from the extension of the pre-existing cracks at very low loads. Therefore, it is quite dicult to
determine the respective in¯uence of each family of stress-induced crack on the shear modulus. The stress-induced matrix cracks are also responsible for the residual strains at a zero load (Figs 7 and 11). Signi®cant crack opening displacements of a few microns were measured at zero load during the tensile tests. Furthermore, Figs 7 and 11 show that the residual strains are proportional to the damage-induced modulus loss. The residual strains are attributed to the non-closure of the cracks, favored by the presence of rough sliding surfaces (Fig. 14). A contribution of the pre-existing cracks prior to the onset of the stress-induced cracking must also be logically expected under low tensile loads, since
Fig. 16. Micrograph showing ®bers that have been pulled-in during the tensile test (direction of loading 2).
SIRON and LAMON: CARBON/CARBON COMPOSITE
these pre-existing cracks are similar to the a- and bstress induced cracks. During the tensile tests, an initial growth of the pre-existing cracks was not detected, whereas during the shear tests, these cracks propagated under very low loads. The contribution of the pre-existing cracks is indicated by the initial non-linear domain of the tensile stress± strain curves (Fig. 5) and the associated residual deformations (Fig. 7). The residual deformations can be attributed, as previously, to the non-closure of cracks as a result of roughness of the sliding crack surfaces (Fig. 12). Under shear, the threshold is not obvious because of the growth of the pre-existing cracks under very low loads (Fig. 11). 4.2. In¯uence of the crimped tows The non-linear stress±strain behavior is essentially dictated by the delamination cracks which appeared to allow easier expansion of the specimens (Figs 5 and 9). Under tension, the expansion was signi®cantly larger along the direction 2 that involves the crimped tows, relative to that along the direction 1 that involves nearly straight tows. The in¯uence of crimped bundles on delamination and deformations has been examined by several authors [3, 4, 9, 10, 18±21]. Under the applied tensile loads, the crimped ®bers, that take most of the load because of their high modulus relative to that of the matrix, try to stretch [9]. If there is a delamination, stretching of the crimped bundles is easier. This phenomenon is indicated by the comparison of Poisson's ratios which re¯ects much larger lateral contractions when the load is applied parallel to direction 2. The consequence of ®ber stretching in the presence of delaminations is that the material extends more easily, as shown by the signi®cant non-linearity exhibited by the stress± strain curves when the load is applied parallel to direction 2. Fiber stretching leads to a crimp angle decrease. Permanent crimp angle decreases have been measured by microscopy and image analysis in a similar three-directional C/C composite [22] and by scanning electron microscopy in two-dimensional C/SiC composite [23]. As the ®bers try to stretch, forces are applied to the layers below the crimp [9, 21]. If there is a delamination, the crimped ®bers act as a string which tries to straighten and stretch, and the rest of the layers behave like a beam which bends [9]. The stretching tow slides against the adjacent perpendicular tow, along the delamination. Roughness of the crack surfaces prevents complete reverse sliding during unloading, leading to permanent deformations, and the non-linearity of the stress±strain curve observed at very low loads prior to the onset of stress-induced cracking. The permanent elongation of the crimped bundles causes a decrease in the crimp angle, leading to the
6641
slight modulus increase observed at very low loads under tension, prior to onset of stress-induced cracking. Similar eect has been reported in the literature [21±23]. It is worth mentioning that this initial stiening eect does not result from an overestimation of the modulus by the slope of the loading±unloading hysteresis loops, since the loop width was very small under a low load. Instead it is related to the permanent local elongation of the misaligned (along direction 1) or the crimped (direction 2) tows. It seems logical to consider that the presence of a similar stress-induced damage under tension and under shear results essentially from the contribution of the crimped tows to the tensile damage, that caused the same type of cracks (delamination) as the in-plane shear loading is expected to do. However, certain experimental results suggest that the stress-state under shear may be aected either by the crimped tows or by the ®ber architecture. The rather large crack opening displacements (between 2 and 25 mm) under shear loading may indicate a mode I of opening instead of mode II. Furthermore the c-cracks that extended perpendicular to the adjacent ®bers, suggest the presence of normal tensile stresses parallel to the ®ber axis. Therefore, one may conjecture that the shear load is locally decomposed along the tows into tensile and compressive loads. This conjecture seems to be consistent with the results of the ®nite element study of the in¯uence of ®ber bundle crimps on the local stress-state by Anand et al., [4]. Normal tensile and compressive stresses of almost twice the magnitude of the shear stress in the vicinity of the crimped bundles were predicted. This stress-state suggested that the local failure was initiated by a tensile separation of the crimp boundaries. Anand and coworkers examined interlaminar shear loading conditions. Determination of the through thickness stress-state, and particularly at the crimp boundaries, is quite dicult under an in-plane shear loading. Other experimental data suggest that the in¯uence of the crimped tows is not obvious. Thus, it was found that the density of cracks under shear is larger than that obtained under tension, and that, the stress±strain behavior under shear did not exhibit a dependence on the loading direction. Determination of the in¯uence of the ®ber architecture and the crimped tows on the stress-state under shear loading conditions will require further investigations. 4.3. In¯uence of matrix anisotropy The high anisotropy of the pyrocarbon matrix probably favors matrix cracking in the C/C CVIcomposites. Most of the pre-existing (a- and bcracks) and the stress-induced cracks (a-, b-cracks and a portion of the c-cracks) are parallel to the graphite layers (i.e. parallel to the axis of the adjacent ®bers).
6642
SIRON and LAMON: CARBON/CARBON COMPOSITE
Some aspects of the anisotropic nature of the mechanical failure of pyrolitic carbon have been reported in the literature [24]. In addition to interlaminar cleavage, the authors mention two distinct types of layer fracture surfaces, corresponding to crack propagation parallel to or perpendicular to the layers [24]. Layer rupture parallel to the layers is characterized by a low toughness and a low strength when compared with layer rupture characteristics. Furthermore, a great anisotropy has been pointed out for the strength [24]. A limited number of cracks perpendicular to the graphitic layers (i.e. the transverse cracks perpendicular to the adjacent ®bers axis) was observed. Such cracks are involved in the c-cracks which appeared at rather high stresses. This may be related to the above mentioned larger layer rupture characteristics when compared with those pertinent to rupture parallel to the layers. It may also be related to the modulus mismatch between the matrix and the ®ber. As a result, the stress for transverse cracking must be very large. It is worth mentioning that such families of a- and b-cracks are not observed in the 2D woven composites with a sti and isotropic SiC matrix (SiC/SiC and C/SiC composites). Instead, matrix cracking involves essentially transverse cracks perpendicular to the adjacent ®bers, that aect the composite stiness because the load carrying capacity of the matrix is high (Em/Ef12).
(ii) inter-tow cracks (b-cracks) and (iii) partly interand partly intra-tow matrix cracks extending across a longitudinal tow (c-cracks). It was shown that the inter-tow cracks exert a preponderant in¯uence on the mechanical behavior and are responsible for the non-linearity of the stress±strain curves and the presence of permanent deformations. Such delamination cracks are favored by the presence of crimped tows (under tensile loading conditions) and the great anisotropy of pyrocarbon. Similarities between the matrix cracks observed under tension and under shear were evidenced. Furthermore, the in¯uence of the ®ber architecture on the stress-state under shear loading conditions was indicated by several experimental data. However, this feature will require further analysis. This work provides data on the mechanical behavior in relation with damage and microstructure for a three-directional CVI±C/C composite. These data, as well as the measured mechanical properties and the observed failure mechanisms will be useful for developing models for predicting the stress±strain behavior of such composites. In a ®rst step, phenomenological models based upon damage mechanics are preferred, due to the lack of data on constituent properties and particularly on the pyrocarbon matrix. Such a model, based on the results reported in this paper, will appear in a forthcoming paper. Sound micromechanics-based models require sound constituent properties, which are not available.
5. CONCLUSIONS
The mechanical behavior of a three-directional C/ C composite was investigated under tensile and shear loads. Microscope inspection of the surface of test specimens under load and of the interior of a few specimens allowed identi®cation of the matrix damage caused by loading. Then, relationships between features of the stress±strain behavior and the matrix cracks that were detected were proposed. The stress±strain curves consist of a single nonlinear domain from very low loads close to zero. No initial linear domain was observed, as a result of the presence of pre-existing cracks. The mechanical behavior under tension and the related main properties (including Young's modulus and fracture strength) exhibited a great anisotropy, which was attributed to the presence of crimped tows along the warp direction whereas the tows were nearly straight along the weft direction. Various eects including initial stiening, the initial non-linearity of the stress±strain curve, and the initial permanent deformations were attributed to the presence of pre-existing cracks. Three families of stress-induced cracks were detected in the matrix, whatever the loading direction, and whatever the loading conditions (tension or shear): (i) microcracks essentially in the transverse tows (a-cracks),
AcknowledgementsÐThis work has been supported by the University of Bordeaux (M.R.E. Grant) and SocieÂte EuropeÂenne de Propulsion (SEP, Le Haillan, France). The authors are indebted to SEP for the supply of the materials, to B. Humez and E. Inghels for valuable discussions.
REFERENCES 1. Gao, F., Patrick, J.W. and Walker, A., Carbon, 1993, 31, 103. 2. Yurgartis, S. W., McGibbon, B. S., Wallsh, J. and Moosbrugger, J. C., Conference extended abstracts, 21st biennal conference on carbon, American Carbon Society, Bualo, 1993, p. 36. 3. Gupta, V., Anand, K. and Kryska, M., Acta metall. mater., 1994, 42, 781. 4. Anand, K., Gupta, V. and Dartford, D., Acta metall. mater., 1994, 42, 797. 5. Hughes, J. and Yurgartis, S. W., Conference extended abstracts, 21st biennal conference on carbon, American Carbon Society, Bualo, 1993, p. 132. 6. Adams, D.F., Mater. Sci. Engng, 1976, 23, 55. 7. Robinson, C.T. and Francis, P.H., A.S.T.M. (S.T.P. 723), 1981, 85. 8. Stanton, E.L. and Kipp, T.E., AIAA J, 1985, 23, 1437. 9. Peterlik, H., Dommanovich, A., Kromp, W. and Kromp, K., Comput. Sci. Technol., 1992, 45, 65. 10. Cox, B. N. and Flanagan, G., in Handbook of Analytical Methods for Textile Composites, version 1.0. Rockwell Science Center, 1996.
SIRON and LAMON: CARBON/CARBON COMPOSITE 11. BroqueÁre, B., Buttazzoni, B. and Choury, J. J., in Introduction aux MateÂriaux Composites, Vol. 2, ed. Naslain. Editions du CNRS, 1985, p. 405. 12. Remond, Y., Ph.D. thesis, University of Paris VI, 1984. 13. Delhaes, P., in Introduction aux MateÂriaux Composites, Vol. 2, ed. Naslain. Editions du CNRS, 1985, p. 385. 14. Kaae, J.L., Carbon, 1985, 23, 665. 15. Ho, H., Tsai, M.Y., Morton, J. and Farley, G.L., Comp. Mater., 1993, 46, 115. 16. Siron, O., No. 1498. Ph.D. thesis, University of Bordeaux, 1996.
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17. Guillaumat, L. and Lamon, J., Rev. Comp. Mater. Avanc., 1993, 3, 159. 18. Pollock, P.B., Carbon, 1990, 28, 717. 19. Jortner, J., Carbon, 1992, 30, 153. 20. Hsiao, H.M. and Daniel, I.M., Composites Part A, 1996, 27A, 931. 21. Shuler, S.F., Holmes, J.W., Wu, X. and Roach, D., J. Am. Ceram. Soc., 1993, 76, 2327. 22. Tardif, F. to be published. . 23. Dalmaz, A., No. 97. Ph.D. thesis, ISAL 0085, INSA de Lyon, 1997. 24. Sakai, M., Bradt, R.C. and Fischbach, D.B., J. Mater. Sci., 1986, 21, 1491.
Acta mater. Vol. 46, No. 18, pp. 6631±6643, 1998 # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain S1359-6454(98)00269-9 1359-6454/98 $19.00 + 0.00
DAMAGE AND FAILURE MECHANISMS OF A 3-DIRECTIONAL CARBON/CARBON COMPOSITE UNDER UNIAXIAL TENSILE AND SHEAR LOADS O. SIRON and J. LAMON{ Laboratoire des Composites Thermostructuraux, UMR 5801 (CNRS-SEP-UB1), 3 AlleÂe de La BoeÂtie, 33600 Pessac, France (Received 5 May 1997; accepted 22 July 1998) AbstractÐThe mechanical behavior of a three-directional carbon/carbon (C/C) composite under tensile and shear loads is investigated in relation with the failure mechanisms and, the ®ber architecture. This three-directional C/C composite was produced by Chemical Vapor In®ltration of a needled ®ber preform of multiple layers of satin woven tows. The C/C composite exhibited several interesting features including an essentially non-linear stress±strain behavior and permanent deformations. Three families of matrix cracks were identi®ed under tensile and shear loads, including microcracks in the tows, intertow delamination and cracks across the longitudinal tows. It was found that the delamination cracks aect preponderantly the stress±strain behavior and the mechanical properties. Similar features in the mechanical behavior and the failure mechanisms were highlighted under tension and under shear loading. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. ReÂsumeÂÐLe comportement meÂcanique d'un composite Carbone/Carbone 3D sous chargement de traction et de cisaillement est eÂtudieÂ. Ce mateÂriau composite C/C a eÂte obtenu par densi®cation par C.V.I. (Chemical Vapor In®ltration) d'une preÂforme ®breuse aiguilleteÂe normalement au plan du tissu. Les relations entre les courbes de comportement meÂcanique, les meÂcanismes de deÂgradation identi®eÂs par microscopie optique et l'architecture ®breuse de la structure ont eÂte eÂtablies. Ce mateÂriau composite C/C preÂsente un comportement meÂcanique non lineÂaire jusqu'aÁ rupture avec deÂformations aneÂlastiques. Les trois meÃmes familles de ®ssures matricielles ont eÂte identi®eÂes sous chargement de traction et de cisaillement, aÁ savoir la ®ssuration intra-®l, la ®ssuration inter-®ls, et, la ®ssuration intra et inter-®ls. La ®ssuration inter-®l, et, la ®ssuration intra et inter-®ls aectent majoritairement le comportement meÂcanique. Les similitudes importantes observeÂes entre le comportement meÂcanique et les meÂcanismes de rupture sous sollicitations de traction et de cisaillement sont discuteÂes. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.
1. INTRODUCTION
Carbon/carbon (C/C) composites display several advantageous properties for structural applications at high temperatures up to 30008C, including: a low density, a good strength retention at high temperatures, a high thermal and chemical stability in inert environments, a high thermal shock resistance, a high resistance to wear and good friction properties at high temperature. Because of these superior thermal and mechanical properties which persist at very high temperatures, C/C composites are used in many areas including disc brakes, wing edges for missiles and space shuttle, solid propulsion applications, furnace tools, fasteners, biomedical, nuclear reactors, etc. The C/C composites are manufactured by in®ltration of the carbon matrix into the carbon ®ber preform using either a liquid or a gaseous route. Unlike the C/C composites manufactured by liquid {To whom all correspondence should be addressed.
impregnation of the matrix [1±9] or the textile composites with a polymer matrix [10], the mechanical behavior of the CVI±C/C (Chemical Vapor In®ltration (CVI) [11]) composites has not been subject of extensive studies. The main trends in the stress±strain behavior and the failure mechanisms that were reported in the literature for those C/C composites obtained by liquid impregnation, may be summarized as follows. The stress±strain behavior depends on the composite structure. The laminated and the braided C/C composites generally exhibit a linear stress±strain behavior under tension or compression, when the load is applied parallel to a ®ber tow direction [12]. By contrast, the stress± strain behavior of the two-dimensional woven composites is essentially non-linear. The non-linear behavior is attributed to matrix and ®ber damage. Some investigations of two-dimensional C/C laminates in which damage was observed during tensile loading, indicated that damage occurred ®rst at the locations of the interface where the woven yarns crossed one another [8]. The failure mechanisms
6631
6632
SIRON and LAMON: CARBON/CARBON COMPOSITE
Fig. 1. 8-H satin weave pattern and tow directions.
under tension identi®ed in a C/C composite with a two-dimensional satin weave reinforcement have been grouped into three families [2]: (i) microcracks located in the transverse tows, (ii) cracks in the transverse tows, and (iii) interply delamination. It
was recently shown that in a satin weave C/C laminate, delamination produces permanent change of the stress±strain behavior [9]. Very few studies have been directed to the local mechanisms of failure under shear loading [4 and references therein]. Several types of two-dimensional laminated C±C composites under shear loading showed crack initiation at the crimp edges situated on the interply interface [4]. The Chemical Vapor In®ltration (CVI)±C/C composites have been reported to dier from those manufactured by liquid impregnation in that the pyrocarbon matrix is characterized by highly anisotropic properties and lower residual stresses [13]. In the present paper, the mechanical behavior of a three-directional CVI±C/C composite under tensile and shear loading conditions at room temperature, is investigated in relation to the damages observed during the tests.
Fig. 2. Crimp angle de®nition (a) and distribution (b).
SIRON and LAMON: CARBON/CARBON COMPOSITE
6633
2. EXPERIMENTAL PROCEDURE
2.1. Material The three-directional C/C composite is a needled laminate reinforced with ex-PAN C ®bers and manufactured by CVI (by SEP, Le Haillan, France). The ®ber architecture consists of needled layers of 8-harness satin weave. The weave structure is depicted in Fig. 1. In the plies [planes (1, 2)], the weft tows (parallel to direction 1) are nearly straight, whereas the warp tows (parallel to direction 2) are locally crimped. The average crimp angle is around 288 (Fig. 2). Fibers were randomly needled perpendicular to the plies (i.e. parallel to direction 3). The volume fraction of ®bers along the directions 1, 2 and 3 is respectively 0.11, 0.11 and 0.03. The pyrocarbon matrix is rough laminar-type (i.e. highly anisotropic) [14]. The total porosity was measured by helium pycnometry on powdered samples: 14.7 2 0.8% (7 measurements). The apparent density of the specimens is 0.93 2 0.05. As discussed in a following section, intra- and inter-tow cracks were present in the as-received samples. 2.2. Uniaxial on-axis tensile tests Twenty one tensile test specimens were cut from a single part so as to make either the weft bundles (direction 1) or the warp bundles (direction 2) the loading axis. The specimen dimensions are given in Fig. 3. The front face of eleven samples (i.e. the face parallel to the plies, as opposed to the lateral faces which are perpendicular to the plies) was polished for microscope examination. The test spe-
Fig. 4. Schematic diagram showing the loading mode of the Iosipescu shear test specimen and the front face which was examined during the tests.
cimen ends were stuck to aluminium pieces that were gripped in the load train of the tensile machine. Special attention was paid to specimen alignment which was ensured on an aluminium test specimen. The specimens were loaded at a strain rate of 0.05% minÿ1 by means of a closed-loop servo hydraulic testing machine. Longitudinal deformations were measured using extensometers mounted respectively on the front and the lateral faces (gauge length: 20 mm). Unloading±reloading cycles were also carried out. Acoustic emission was recorded during the tests. Calibration of the recording system permitted selection of the appropriate parameters to eliminate extraneous noise from the testing device. The polished front surface was observed under an optical microscope before the tests, during the tests and after the ultimate failure. The examination of the specimens under load used a high magni®cation microscope (up to 416) built on a stage with three displacement axes and coupled with a CCD camera, a video recorder and a control monitor. The same region in the gauge area was inspected at various loads according to a grid of 50±70 views, covering roughly one-®fth of the total gauge area. The load was kept constant during specimen examination. The load steps correspond to 0.05% deformation increments. An alternative method was used to observe through thickness failure modes. For this purpose, a mold was mounted on the specimens. Under a load close to the ultimate failure, the specimens were vacuum impregnated with an epoxy resin to ®ll in the open cracks. The load was kept constant during resin polymerization. Further, the specimens were sectioned and polished parallel to the three reference planes de®ned by the directions 1, 2 and 3, and examined using an optical microscope. This method was applied to two specimens. 2.3. Intralaminar Iosipescu shear tests
Fig. 3. Dimensions (mm) of the test specimens: (a) tensile specimens, (b) specimens for Iosipescu shear tests.
The Iosipescu test specimens were cut out of the previously mentioned part, so as to make either the weft bundle (direction 1) or the warp bundles (direction 2) the loading axis. The specimen dimensions are given in Fig. 3. Figure 4 shows the front face and the orientation of the plies with respect to the loading axis. The tests were performed under
6634
SIRON and LAMON: CARBON/CARBON COMPOSITE
Fig. 5. Typical stress±strain curves obtained during the tensile tests under: (a) a load parallel to the direction 1 of tows; and (b) a load parallel to the direction 2 of tows.
displacement control at a strain rate of 0.05 mm minÿ1. The deformations were measured on the rear face by means of a 2458 strain gauge (Vishay CEA06-125-WT-350). The front face was polished for microscope examination of the area located between the ``V'' notches (Fig. 4). The two previously described methods of damage observation were used. About 70 micrographs covering roughly one-quarter of the ``V'' notch area were recorded. The load was kept constant during surface examination. The load was increased by 5 MPa increments. The resin impregnation technique was applied to one sample under a load close to ultimate failure. The shear tests were interrupted at a stress of 145 MPa. 3. EXPERIMENTAL RESULTS
3.1. Stress/strain behavior 3.1.1. On-axis tensile tests. Figure 5 shows typical stress±strain curves which evidence several features. First, the three-directional C/C composite is essentially non-linear, from very low loads, whatever the loading direction. No signi®cant initial linear domain was identi®ed. Second, the strain at a given stress is much larger (roughly two times), when the direction 2 is aligned with the tensile axis. Third, the unloading±reloading hysteresis loops are very narrow. Fourth, signi®cant permanent deformations are observed after unloading. Table 1 summarizes the results of the on-axis tensile tests. Young's modulus was determined from the slope of the tangent to the curve at the origin. It is around 30 GPa, and the strain-to-failure is around 0.5%. The higher Young's modulus and failure stress are obtained in direction 1, which suggests that the crimped tows exert a certain in¯uence on the mechanical properties. Comparison of
the Poisson's ratios indicates larger lateral contractions (about two times) when the load is applied parallel to the direction 2. This feature also indicates an eect of the presence of crimped tows in this direction. Figure 6 describes stiness changes during the tensile tests. The stiness was determined from the slope of the unloading±reloading hysteresis loops. This slope provides a satisfactory approximation of the composite stiness, since the loops width was generally rather small (Fig. 5). It may be noticed from Fig. 6 that the stiness seems to increase ®rst. Such stiening must not be attributed to an artefact but instead it is related to the pre-existing matrix cracks, as discussed in a following section. Acoustic emission was detected from 0.05% applied deformations (Fig. 5). It develops exponentially from 0.2% deformations, so that most of the acoustic emission is recorded close to the ultimate failure. Figure 5 also shows that no acoustic emission is observed during the unloading±reloading cycles suggesting the occurrence of limited damage phenomena. Comparable numbers of counts are obtained whatever the loading direction. Figure 7 shows a linear relationship between the residual strains at zero load and the degree of damage measured by a modulus loss [1-(Ei/E0i ), Ei=composite stiness, E0i =initial Young's modulus along the loading direction i(i = 1,2)]. Beyond Table 1. Main mechanical properties of the investigated C/C composite determined under on-axis tensile tests E0i (GPa) (i = 1,2)
n0
sr (MPa)
er (%)
1
33 22
126.6 24.6
0.482 0.06
2
22.2 20.5
n012=0.18 20.01 n013=0.21 20.02 n023=0.39 20.02
64.5 24.9
0.472 0.07
Direction of loading
SIRON and LAMON: CARBON/CARBON COMPOSITE
6635
Fig. 6. Stiness changes vs longitudinal strains during the on-axis tensile tests under loads parallel to direction 1 or to direction 2.
a threshold of around 0.02%, the residual strains seem to be proportional to the damage parameter, and independent of the loading direction. Below the threshold, the residual strains are not (or little) related to the damage parameter suggesting that they do not result from a stress-induced damage. The Poisson's ratios shown by Fig. 8 re¯ect a decrease in the lateral contractions as matrix cracking proceeds. This trend is generally observed with ceramic matrix composites (CMCs). It is attributed to cracking parallel to the loading direction (typically ®ber debonding in CMCs). Figure 8 also shows that the Poisson's ratio dependence on the loading direction becomes more pronounced during the tests. This indicates that the lateral contractions in direction 3, are larger by a
factor >2 when the load is applied parallel to direction 2, relative to those obtained when the load is applied parallel to direction 1. 3.1.2. Iosipescu in-plane shear tests. Figure 9 shows a typical stress±strain curve from the Iosipescu shear tests. The stress was derived from the ratio of the applied force to the cross section area between the notches. The mechanical properties did not need to be corrected, since the possible eect of ply anisotropy can be considered to be low: E01/ E0211.6. As a matter of fact, for this degree of anisotropy, the correction factor [15] is smaller than 5% for the shear modulus. The stress±strain curve exhibits features similar to those observed under the on-axis tensile loading conditions. First, the stress±strain relationship is essentially non-linear. No initial domain of linear behavior could be identi®ed. Second, the unload-
Fig. 7. Evolution of residual strains vs a damage parameter 1 ÿ (Ei/E0i ), i = 1,2.
Fig. 8. Poisson coecient changes, along direction 2 or 3 under tensile loads applied parallel to direction 1 or 2.
6636
SIRON and LAMON: CARBON/CARBON COMPOSITE
Fig. 9. Typical stress±strain curve obtained during the Iosipescu shear tests.
ing±reloading hysteresis loops are ®rst rather narrow. Then they become wider under increasing loads, suggesting the presence of friction phenomena near the ultimate failure. Third, permanent deformations are observed after unloading. Nevertheless, it is worth mentioning that the stress± strain curves did not display any dependence upon the loading direction. The shear modulus that was derived from the tangent to the stress±strain curves at the origin, is obtained to be G012=5.8 2 0.2 GPa. Figure 10 shows the shear modulus degradation determined from the slopes of the unloading±reloading hysteresis loops. As opposed to the stiness which decreased by <25% under tension, the shear modulus loss reaches about 70%. Acoustic emission was detected at a 0.05% shear strain (Fig. 9). The number of counts is quite proportional to the shear strain. No acoustic emission was observed during the unloading±reloading cycles, suggesting the occurrence of limited damage. As previously for the tensile tests, a linear relationship is obtained between the residual strains and the degree of damage (Fig. 11) measured using the shear modulus loss: 1 ÿ (G12/G012). This relationship indicates that the residual strains are related to
the stress-induced damage. The presence of a threshold is not obvious (<0.01%).
Fig. 10. Shear modulus changes vs shear stress during the Iosipescu shear test.
Fig. 11. Evolution of residual shear deformations e12p vs a damage parameter 1 ÿ (G12/G012).
3.2. Microscope observations of matrix damage 3.2.1. As-received samples. Two families of preexisting cracks were identi®ed in the as-received specimens (Fig. 12): (i) intra yarn microcracks located in the transverse tows, in the matrix, or in the ®bermatrix interface, parallel to the ®ber direction [referred to as a-cracks: Fig. 12(a)]; and (ii) inter yarns microcracks located between two perpendicular tows [referred to as b-cracks; Fig. 12(b)]. The microcracks formed during composite processing at high temperatures and the subsequent cooling down cycles. They result from thermally-induced stresses that arise because of a thermal expansion mismatch between ®bers and matrix (for the a-cracks involving ®ber-matrix debonding) and also between perpendicular tows (for the b-cracks). However, reliable data on the thermal expansion coecients of the CVI±pyrocarbon matrix and the C ®bers are not available yet. Estimates of the density of pre-existing microcracks were derived from the number of micro-
SIRON and LAMON: CARBON/CARBON COMPOSITE
6637
Fig. 12. Micrographs showing the pre-existing intra-yarn microcracks (a) and the inter-yarn microcracks (b) detected in the as-received C/C specimens.
cracks detected in the specimen area inspected with the microscope. The average densities are respectively 3.40 2 1.30 mmÿ2 for the a-cracks, and 0.45 2 0.27 mmÿ2 for the b-cracks. 3.2.2. Specimens tested under on-axis tension. The microcracks and cracks that have been detected in the specimens during the tests may be grouped into three families on the basis of their location, whatever the loading direction (1 or 2): (i) a-microcracks (Fig. 13) parallel and perpendicular to the loading direction. They are dictated by the high anisotropy of pyrocarbon (i.e. the orientation of the graphitic layers) and by the presence of pores. Besides, the pre-existing a-micro-
cracks did not generally propagate during the tensile tests. (ii) b-cracks (Figs 13 and 14) parallel and perpendicular to the loading direction. (iii) c-cracks (Figs 13 and 14) consisting of matrix cracks that extend ®rst between a longitudinal tow and either a transverse tow or the intertow matrix, and then across a longitudinal tow. They are then preferentially oriented perpendicular to the loading direction. However, cracks parallel to the loading direction as well as cracks that deviated at 908, were also observed. Depending on the applied load, the c-cracks were either shorter than two tow
Fig. 13. Schematic diagram summarizing the three families of cracks detected during the tensile and the Iosipescu shear tests.
6638
SIRON and LAMON: CARBON/CARBON COMPOSITE
Fig. 14. Optical micrographs showing the b- and c-cracks created in the matrix during the tensile tests.
diameters (c1-cracks), or as long as two or three tow diameters (c2-cracks). Table 2 shows that the a-, b- and c-cracks formed in chronological order. The a-cracks initiated ®rst
at a 0.05% applied deformation, whatever the loading direction. This deformation coincides with the onset of acoustic emission. Then, the b-cracks, and ®nally the c-cracks were observed at higher defor-
Table 2. Matrix cracking stresses and strains determined during on-axis tensile tests Direction of loading 1 2
strain stress strain stress
a-crack
b-crack
c1-crack
c2-crack
0.05±0.1% 16±30 MPa 0.05±0.1% 10±20 MPa
0.1±0.2% 30±60 MPa 0.05±0.15% 10±30 MPa
0.2±0.3% 60±90 MPa 0.1±0.2% 20±40 MPa
0.3±0.35% 90±100 MPa 0.2±0.3% 40±50/55 MPa
SIRON and LAMON: CARBON/CARBON COMPOSITE Table 3. Mean crack densities and sizes measured after on-axis tensile tests Crack family
Mean density (mmÿ2)
Mean length (mm)
1.00 0.50 0.55 0.25 0.15 0.05
0.14 0.20 0.20 0.35 0.35 1.00
a (perpendicular)* a (parallel)* b c1 (perpendicular)* c1 (parallel)* c2 (perpendicular)* *To loading direction
mations (Table 2). The matrix cracking stresses and deformations were generally smaller for those specimens loaded parallel to the direction 2, suggesting an eect of the crimped tows. Saturation of matrix cracking was not identi®ed. The crack densities increased steadily with the applied load. The three families of cracks were distributed homogeneously. Table 3 summarizes the crack characteristics obtained from the examination of the surface and the interior of the specimens. Further details are provided in [16]. The crack densities are generally <1 crack/mm2. They must be regarded as small when comparing with the crack densities observed in ceramic matrix composites including the SiC/SiC composites [17]. Data on the crack lengths indicate that the a- and b-cracks are localized failures. The c-cracks are the longest and include one up to two or three yarn diameters. 3.2.3. Specimens tested under shear. The cracks created during the Iosipescu shear tests pertain to the three previously described families of a-, b- and c-matrix cracks (Fig. 13). No additional damage modes nor crack types were observed. Table 4 summarizes the matrix cracking stresses and deformations. Data in parentheses refer to the extension of the pre-existing cracks since contrary to the on-axis tensile tests, the pre-existing a- and b-cracks propagated. ``Mixed'' designates the cracks which were de¯ected to a perpendicular direction. As previously, the three families of cracks may be regarded as independent. The a- and b-cracks initiated under 20±25 MPa whatever their orientation. The c-cracks initiated under 25 MPa (when parallel to the loading direction) and around 35±40 MPa (when perpendicular to the loading direction). As previously, the matrix cracks were homogeneously
6639
Table 5. Mean crack densities (mmÿ2) measured after a maximum applied stress of 45 MPa under Iosipescu shear tests Crack orientation* Perpendicular Parallel Mixed
a-crack
b-crack
c-crack
0.53 20.40 1.43 20.44 0.30 20.07
0.252 0.13 0.592 0.42 Ð
0.092 0.06 0.282 0.17 0.192 0.14
*To loading direction
distributed, and the notch tips did not appear to be preferential sites of cracking. The densities of cracks that initiated in the surface of specimens during the Iosipescu tests (Table 5), are larger than those determined during the tensile tests (Table 3). Besides, comparable densities were generally obtained for both the cracks perpendicular or parallel to the tow direction at stresses <15 MPa. The parallel cracks were preponderant at stresses >15 MPa. Comparable data were determined for the cracks detected in the interior of specimens [16]. The lengths of the a-, b- and c1-cracks are close to those measured in the tensile specimens. They are, respectively, 0.16 2 0.05 mm, 0.19 2 0.09 mm and 0.37 2 0.08 mm. The length of the c2-cracks was dicult to determine because of crack deviation. A mean length of about 1 mm was estimated. 4. DISCUSSION
The stress±strain behavior was found to exhibit several features such as a non-linearity even at very low loads, residual deformations at zero load and a signi®cant anisotropy. Let us examine the contribution of various factors including the pre-existing and the stress-induced damage, the presence of crimped tows and the matrix anisotropy. 4.1. In¯uence of damage Figure 15 shows the in¯uence of stress-induced matrix cracks on the stiness. Under tension, the acracks do not in¯uence signi®cantly the stiness whereas a stiness loss can be noticed when the bcracks appear whatever the loading direction. The c-cracks which also involve a localized intertow delamination contribute to the stiness degradation. However, their transverse extension across the matrix in longitudinal tows may not be expected to
Table 4. Matrix cracking stresses and strains determined during the Iosipescu shear tests (data in parentheses refer to the extension of the pre-existing cracks) Crack orientation* Perpendicular Parallel Mixed *To loading direction
a-crack
b-crack
c-crack
(10±20) 25±45 MPa (0.1±0.2) 0.3±1.2% (10±20) 25±45 MPa (0.1±0.2) 0.2±1.2% (0±15) 25±45 MPa (0±0.14) 0.3±1.2%
(0±20) 20±45 MPa (0±0.2) 0.2±1.2% (0±20) 20±45 MPa (0±0.2) 0.2±1.2% (30±40 MPa) (0.6±0.87%)
(20±35) 40±45 MPa (0.2±0.6) 0.9±1.2% (10±20) 25±45 MPa (0.1±0.2) 0.3±1.2% (15±35) 35±45 MPa (0.14±0.6) 0.6±1.2%
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SIRON and LAMON: CARBON/CARBON COMPOSITE
Fig. 15. Stiness (a) and shear modulus (b) changes as a function of the failure mechanisms. n, normal to the axis of loading; p, parallel to the axis of loading; m, mixed (i.e. parallel and normal to the axis of loading); dv, development of a damage mode; pr, extension of a pre-existing damage mode.
aect tremendously the composite stiness, because of the limited load carrying capacity of the pyrocarbon matrix, which exhibits a low longitudinal elastic modulus (Em130 GPa) relative to that of ®bers (Ef1200 GPa). Therefore, transverse cracking of the pyrocarbon matrix does not increase the load on the ®bers nor the associated deformations. Furthermore, poor load transfers must be expected through the ®ber/matrix interfaces. As a matter of fact, Fig. 16 shows that ®bers have been pulled in. Under shear load, a shear modulus loss [Fig. 15(b)] coincides with onset of matrix cracking. Most of the stress-induced damage during the shear tests results from the extension of the pre-existing cracks at very low loads. Therefore, it is quite dicult to
determine the respective in¯uence of each family of stress-induced crack on the shear modulus. The stress-induced matrix cracks are also responsible for the residual strains at a zero load (Figs 7 and 11). Signi®cant crack opening displacements of a few microns were measured at zero load during the tensile tests. Furthermore, Figs 7 and 11 show that the residual strains are proportional to the damage-induced modulus loss. The residual strains are attributed to the non-closure of the cracks, favored by the presence of rough sliding surfaces (Fig. 14). A contribution of the pre-existing cracks prior to the onset of the stress-induced cracking must also be logically expected under low tensile loads, since
Fig. 16. Micrograph showing ®bers that have been pulled-in during the tensile test (direction of loading 2).
SIRON and LAMON: CARBON/CARBON COMPOSITE
these pre-existing cracks are similar to the a- and bstress induced cracks. During the tensile tests, an initial growth of the pre-existing cracks was not detected, whereas during the shear tests, these cracks propagated under very low loads. The contribution of the pre-existing cracks is indicated by the initial non-linear domain of the tensile stress± strain curves (Fig. 5) and the associated residual deformations (Fig. 7). The residual deformations can be attributed, as previously, to the non-closure of cracks as a result of roughness of the sliding crack surfaces (Fig. 12). Under shear, the threshold is not obvious because of the growth of the pre-existing cracks under very low loads (Fig. 11). 4.2. In¯uence of the crimped tows The non-linear stress±strain behavior is essentially dictated by the delamination cracks which appeared to allow easier expansion of the specimens (Figs 5 and 9). Under tension, the expansion was signi®cantly larger along the direction 2 that involves the crimped tows, relative to that along the direction 1 that involves nearly straight tows. The in¯uence of crimped bundles on delamination and deformations has been examined by several authors [3, 4, 9, 10, 18±21]. Under the applied tensile loads, the crimped ®bers, that take most of the load because of their high modulus relative to that of the matrix, try to stretch [9]. If there is a delamination, stretching of the crimped bundles is easier. This phenomenon is indicated by the comparison of Poisson's ratios which re¯ects much larger lateral contractions when the load is applied parallel to direction 2. The consequence of ®ber stretching in the presence of delaminations is that the material extends more easily, as shown by the signi®cant non-linearity exhibited by the stress± strain curves when the load is applied parallel to direction 2. Fiber stretching leads to a crimp angle decrease. Permanent crimp angle decreases have been measured by microscopy and image analysis in a similar three-directional C/C composite [22] and by scanning electron microscopy in two-dimensional C/SiC composite [23]. As the ®bers try to stretch, forces are applied to the layers below the crimp [9, 21]. If there is a delamination, the crimped ®bers act as a string which tries to straighten and stretch, and the rest of the layers behave like a beam which bends [9]. The stretching tow slides against the adjacent perpendicular tow, along the delamination. Roughness of the crack surfaces prevents complete reverse sliding during unloading, leading to permanent deformations, and the non-linearity of the stress±strain curve observed at very low loads prior to the onset of stress-induced cracking. The permanent elongation of the crimped bundles causes a decrease in the crimp angle, leading to the
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slight modulus increase observed at very low loads under tension, prior to onset of stress-induced cracking. Similar eect has been reported in the literature [21±23]. It is worth mentioning that this initial stiening eect does not result from an overestimation of the modulus by the slope of the loading±unloading hysteresis loops, since the loop width was very small under a low load. Instead it is related to the permanent local elongation of the misaligned (along direction 1) or the crimped (direction 2) tows. It seems logical to consider that the presence of a similar stress-induced damage under tension and under shear results essentially from the contribution of the crimped tows to the tensile damage, that caused the same type of cracks (delamination) as the in-plane shear loading is expected to do. However, certain experimental results suggest that the stress-state under shear may be aected either by the crimped tows or by the ®ber architecture. The rather large crack opening displacements (between 2 and 25 mm) under shear loading may indicate a mode I of opening instead of mode II. Furthermore the c-cracks that extended perpendicular to the adjacent ®bers, suggest the presence of normal tensile stresses parallel to the ®ber axis. Therefore, one may conjecture that the shear load is locally decomposed along the tows into tensile and compressive loads. This conjecture seems to be consistent with the results of the ®nite element study of the in¯uence of ®ber bundle crimps on the local stress-state by Anand et al., [4]. Normal tensile and compressive stresses of almost twice the magnitude of the shear stress in the vicinity of the crimped bundles were predicted. This stress-state suggested that the local failure was initiated by a tensile separation of the crimp boundaries. Anand and coworkers examined interlaminar shear loading conditions. Determination of the through thickness stress-state, and particularly at the crimp boundaries, is quite dicult under an in-plane shear loading. Other experimental data suggest that the in¯uence of the crimped tows is not obvious. Thus, it was found that the density of cracks under shear is larger than that obtained under tension, and that, the stress±strain behavior under shear did not exhibit a dependence on the loading direction. Determination of the in¯uence of the ®ber architecture and the crimped tows on the stress-state under shear loading conditions will require further investigations. 4.3. In¯uence of matrix anisotropy The high anisotropy of the pyrocarbon matrix probably favors matrix cracking in the C/C CVIcomposites. Most of the pre-existing (a- and bcracks) and the stress-induced cracks (a-, b-cracks and a portion of the c-cracks) are parallel to the graphite layers (i.e. parallel to the axis of the adjacent ®bers).
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SIRON and LAMON: CARBON/CARBON COMPOSITE
Some aspects of the anisotropic nature of the mechanical failure of pyrolitic carbon have been reported in the literature [24]. In addition to interlaminar cleavage, the authors mention two distinct types of layer fracture surfaces, corresponding to crack propagation parallel to or perpendicular to the layers [24]. Layer rupture parallel to the layers is characterized by a low toughness and a low strength when compared with layer rupture characteristics. Furthermore, a great anisotropy has been pointed out for the strength [24]. A limited number of cracks perpendicular to the graphitic layers (i.e. the transverse cracks perpendicular to the adjacent ®bers axis) was observed. Such cracks are involved in the c-cracks which appeared at rather high stresses. This may be related to the above mentioned larger layer rupture characteristics when compared with those pertinent to rupture parallel to the layers. It may also be related to the modulus mismatch between the matrix and the ®ber. As a result, the stress for transverse cracking must be very large. It is worth mentioning that such families of a- and b-cracks are not observed in the 2D woven composites with a sti and isotropic SiC matrix (SiC/SiC and C/SiC composites). Instead, matrix cracking involves essentially transverse cracks perpendicular to the adjacent ®bers, that aect the composite stiness because the load carrying capacity of the matrix is high (Em/Ef12).
(ii) inter-tow cracks (b-cracks) and (iii) partly interand partly intra-tow matrix cracks extending across a longitudinal tow (c-cracks). It was shown that the inter-tow cracks exert a preponderant in¯uence on the mechanical behavior and are responsible for the non-linearity of the stress±strain curves and the presence of permanent deformations. Such delamination cracks are favored by the presence of crimped tows (under tensile loading conditions) and the great anisotropy of pyrocarbon. Similarities between the matrix cracks observed under tension and under shear were evidenced. Furthermore, the in¯uence of the ®ber architecture on the stress-state under shear loading conditions was indicated by several experimental data. However, this feature will require further analysis. This work provides data on the mechanical behavior in relation with damage and microstructure for a three-directional CVI±C/C composite. These data, as well as the measured mechanical properties and the observed failure mechanisms will be useful for developing models for predicting the stress±strain behavior of such composites. In a ®rst step, phenomenological models based upon damage mechanics are preferred, due to the lack of data on constituent properties and particularly on the pyrocarbon matrix. Such a model, based on the results reported in this paper, will appear in a forthcoming paper. Sound micromechanics-based models require sound constituent properties, which are not available.
5. CONCLUSIONS
The mechanical behavior of a three-directional C/ C composite was investigated under tensile and shear loads. Microscope inspection of the surface of test specimens under load and of the interior of a few specimens allowed identi®cation of the matrix damage caused by loading. Then, relationships between features of the stress±strain behavior and the matrix cracks that were detected were proposed. The stress±strain curves consist of a single nonlinear domain from very low loads close to zero. No initial linear domain was observed, as a result of the presence of pre-existing cracks. The mechanical behavior under tension and the related main properties (including Young's modulus and fracture strength) exhibited a great anisotropy, which was attributed to the presence of crimped tows along the warp direction whereas the tows were nearly straight along the weft direction. Various eects including initial stiening, the initial non-linearity of the stress±strain curve, and the initial permanent deformations were attributed to the presence of pre-existing cracks. Three families of stress-induced cracks were detected in the matrix, whatever the loading direction, and whatever the loading conditions (tension or shear): (i) microcracks essentially in the transverse tows (a-cracks),
AcknowledgementsÐThis work has been supported by the University of Bordeaux (M.R.E. Grant) and SocieÂte EuropeÂenne de Propulsion (SEP, Le Haillan, France). The authors are indebted to SEP for the supply of the materials, to B. Humez and E. Inghels for valuable discussions.
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