Crack development in a carbon-carbon laminate

992

Letters

National InstituteforResources and Environment

Tsukuba, Ibaraki JAPAN

to the Editor

M. KOBAYASHI E. ISHIKAWA Y. TODA

7. 8. 9.

REFERENCES I.

2. 3. 4. 5.

6.

M.M. Dubinin, Izv. Akad, Nauk SSSR Ser. Khim. 78, 17 (19781. M.M. Dubinin, I&. Akad, Nauk SSSR Ser. Khim. 78, 529 (1978). M.M. Dubinin and T.S. Jakubov, fzv. Akad, Nauk SSSR Ser. K&m. 77, 2428 (1977). T.S. Jakubov, B.P. Bering, M.M. Dubinin and V.V. Serpinsky, Izv. Ak%d, Nauk SSSR Ser. Khim. 77. 463 (1977X M. Jaroniec anh A.W. Marczewski, J. Colloid Inferface Sci., 101, 280 (1984).

Crack development

10. 11. 12.

M. Jaroniec and J. Choma, “Chemistry and Phvsics of Carbon”. Vol. 22 (Ed. P.A. Thrower). II Dekker, New York (1989), p. ‘197. Y. Toda and S. Toyoda, Carbon, 10, 231 (1972). H. Marsh and T. Siemieniewska, Fuel, 46, 441 (1967). D.M. Young and A.D. Crowell, “Physical Adsorption of Gases”, Butterworths, London (1962), Chapters 3 and 4. “Chemistry and Physics of M.M. Dubinin, Carbon”, Vol. 2 (Ed. P.L. Walker, Jr.), Dekker, New York (1966), p. 51. M. Kobayashi, E. lshikawa and Y. Toda, Curhun, 29, 677 (1991). W. Braker and A. Mossman, Matteson Gas Data Book, p. 23 1, Sixth Edition.

in a carbon-carbon

laminate

(Received 2 February 1993; accepted in revisedform 15 April 1993) Key Words - Carbon/carbon

Recent experiments have identified three types of microcracks within a carbon-carbon laminate : straight cracks, angle cracks and S-cracks. These cracks play an important role in the formation of tensile damage and the development of delamination. The geometry and significance of these cracks is explained here, along with their role in tensile failure m~h~isms. The material studied in this work was an eightharness satin (8HS) carbon-carbon laminate. The laminate was reinforced with T-300 fibers which were heat stabilized at 25OO’C for several hours prior to composite fabrication. The yarn counts were 9.45 yarn/cm (24 yarn/inch) for the warp direction, and 9.05 varn/cm (23 yarn/inch) for the fill direction. The average crimp angle (see Ref. [l] and [2]) in the warp direction was 8.4 degrees. with a standard deviation of 3.5 degrees. Th\ average crimp angle in the fill direction was 10.0 degrees, with a standard deviation of 3.2 degrees. The-average ply count was 38.5 ply/cm (97.7 ulv/in). and the densitv of the material (mass divided bv a ,. geometric volume) was 1.60 g/cc. Thd void content ih the composite, based upon microphotography and computerized feature analysis, varied from 7-9%. The matrix in the carbon-carbon was deposited entirely by a CVD process. This as-received CVD material had a maximum processing temperature not exceeding 990°C. Two sets of samples were prepared from the as-received material. The first set of samples was heat treated to 1480°C, while the second set was heat-treated to 235O’C. Figure la shows the microstructure of one of the carbon-carbon samples which was heat-treated to 235O’C. Features of the crack pattern in this photograph are sketched and named in Figure lb. Although this material is not a cross-ply laminate, the transverse yarns do tend to bIend together at a microstructurai level and can be thought of as a ‘*transverse ply”. Both “straight ,I

composites;

crack; delamination

cracks” and “angle cracks” can be seen in the transverse ply. The straight cracks run at right angles to the longitudinal yarns, and are similar to the transverse matrix cracks seen in graphite-epoxy composites under tensile load [3]. These straight cracks were visible in all the samples of the carbon-carbon composite, including the as-received CVD material. However, the samples in this study were not subjected to a mechanical load, and so the straight cracks must have been caused by processing stresses. These processing stresses can be classified either as thermal stresses or as matrix shrinkage stresses. Thermal stresses arise because of differences in the thermal expansion coefficients of the yarns in the transverse and longitudinal directions. This causes the cross-sections of the yarns to be in tension when they cool down from elevated temperatures. A shear lag phenomenon [3] is thought to be responsible for the pattern of straight cracks seen in Figure la. It is also possible for shrinkage stresses to occur in carboncarbon composites during processing, as a result of either matrix pyrolysis or changes in the crystallinity of the constituents ]I]. In this study, shrinkage stresses due to matrix pyrolysis were avoided because the entire matrix was formed by CVD, while changes in the fiber crystaltinity were minimized because of the initial heat stabilization of the T-300 fibers. Therefore shrinkage stresses are thought to be of secondary importance, compared to the thermal stresses. Figure la also shows a group of cracks which are designated as angle cracks. These angle cracks appear to be similar to the partial angle cracks and curved cracks reported by Groves et al. [4] for cross-ply graphite-epoxy laminates subjected to tensile loads. However, the angle cracks observed in the carboncarbon were generated by thermal stresses during processing, and not by mechanical loading. The samples heat-treated at 2350°C showed the best examples of

Letters

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to the Editor

longitudinal yarn

faint microTracks

I

transverse straight -crack

straight crack-

angle

8

\ longitudinal

yarn Figure 1. (a) Crack pattern in a fill specimen heated to 235iYC; (b) Key to micrograph and definition of crack angle 8 angle crack development (as in Figure la). Fewer angle cracks were observed in the samples heated to 148O”C, and none were observed in the as-received CVD material. Angle cracks were never observed without straight cracks also being present. Presumably the angle cracks form at a later stage than the straight cracks, and at higher stress levels. Figure la shows a case where three angle cracks are visible in the region between the two straight cracks. Two fine microcracks are present in the fourth location where an angle crack might have formed. The number of angle cracks in such a region (bounded by straight cracks) was observed to vary from one to four. When these angle cracks reach the interface with a fiber bundle, they sometimes turn and become partial delanlination cracks (see Figure la). Angle cracks were observed most frequently at locations where the thickness of the transverse ply (i.e. transverse yarns) was high e.g. at the cross-over points

in the satin weave. Figure la was taken at such a location. Groves et al. ]4] also observed the same behavior in graphite-epoxy laminates, where the number of angle cracks (or curved cracks) increased with the percentage content of 90 degree plies in the laminate. However, in the present case it was rare to see the angle cracks coalesce into the “curved cracks” reported by Groves et al. Also, the suggestion [4] that the angle cracks are formed by a ductile fracture process seems unlikely here. The matrix in the carbon-carbon laminate is brittle, and does not exhibit ductile fracture. Figure lb defines the acute angle 8 which the cracks make with the ply axis. Figure 2 shows the experimental distribution of angle cracks versus 8. This distribution was measured for samples heated to 2350°C. It can be seen that the highest number of angle cracks occurs within the range of 55-60 degrees. This agrees quite well with measurements by Groves et al.,

Letters

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the Editor

(a)

50 00

warp.

___

L

./_ /

00

25

3.0 35 40

45

Crack

Angle

Figure 2. Histogram (2350°C treatment).

50

55

60 65

70, 75

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fill

;__-_-_-_-:

_

I

-‘-y-----

I/“-. “d 1000

. -----

_

__~_____-.___...._ -2000 Strain

6 (degrees)

.

Is;-

-_“-..___.--._.

~~ 50

/ ____

..:p_cl’-

- ------i

< 5 0

:

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(microstrain)

of number of angle cracks vs. O

who observed the highest number of angle cracks at an orientation of 60-70 degrees (Groves et al. actually report the orientation of the principal stress, which has a value equal to [90 - O] degrees). This agreement between the two studies is surprising, because the types of materials and loadings are quite different. Evidently the microstructural stress fields must be similar in both cases. Attempts by the current authors to predict the orientation of these angle cracks have not yet been successful, Groves et al. report a finite-element analysis which appears to support their own experimental data. However, this agreement can only be obtained by inco~orating a resin-rich layer at the ply interfaces, and then by adjusting the properties of this layer until the analysis matches the experimental data. A compliant matrix layer of this type is plausible in graphite-epoxy composites, but seems inappropriate in a carbon-carbon composite where the matrix is brittle. An analysis by the current authors shows that there are very large interlaminar shear stresses at the fiber bundle interfaces, and that these stresses should initiate large-scale delamination. Further work is needed to explain why the angle crack mechanism is preferred, and why catastrophic delamination does not take place. Figure 3a, 3b show the tensile stress-strain curves for the carbon-carbon composite. The curves were obtained from specimens having a dogbone shape, and with the following dimensions : total length 15.24 cm (6.0 in.), length of gage section 7.62 cm (3.0 in.), width of gage section 1.90 cm (0.75 in.), and specimen thickness 0.838 cm (0.330 in.). Strain measurements were recorded using an extensomer. It is seen in Figure 3 that the stress-strain curves become nonlinear at high loads and show large increases in strain. This flattening of the stress-strain curve is caused by delamination. Gross ply separation was clearly visible during this part of the test. The onset of delamination, where nonlinear behavior begins, could usually be identified by an audible cracking sound coming from the specimens. Delamination began at the free edges of the specimens, as verified by direct mic~meter readings of the specimen thickness. The most noticeable difference between the samples heated to 1480°C and 2350°C was the strain to failure. A much large failure strain was observed in the samples heat treated at the lower temperature. Small increases in the delamination stress and ultimate stress were also observed. Note that the warp stress-strain curve shown in Figure 3b (235O’C treatment) is incomplete because some of the extensometer data were

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1000

2000 Strain

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4000

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(microstrain)

Figure 3. Tensile stress-strain curves for heat-treated at (a) 1480°C and (b) UfWC

SpeCimenS

lost (due to slippage). The ultimate stress in the warp direction for this material was 266 MPa (38.6 ksi). The rem~ning stress-strain curves are complete. A surface replication method was chosen to visualize the damage in the tensile specimens during loading. Cellulose tape* wetted with acetone was applied to the polished edges of some test specimens, and impressions of the crack patterns were recorded. This technique is quite difficult to apply to carboncarbons because of their high surface porosity. However, with experience it is possible to identify which image features are associated with the replication method (e.g. bubbles), and which are due to the crack pattern in the composite. Figures 4a-4c show the progression of tensile damage in a fill specimen which had been heat treated to 148O’C. Tensile loading is applied along the horizontal axis of the photographs. The large darkened areas and bubbles in the photographs are defects due to the replication method. At a stress of 159 MPa (23 ksi) in Figure 4a a few loading cracks are visible in the middle transverse ply (i.e. transverse yarns). These loading cracks run at an angle to the ply direction, and often initiate from the angle cracks previously described. At a higher stress level of 200 wa (29 ksi) in Figure 4b the pattern of loading cracks becomes more regular and welldefined. It is seen that the cracks take on an S-shape, and so they are referred to here as “S-cracks”. Figure 5 shows a magnified view of a set of S-cracks, taken at a different location on the specimen. It is seen that the Scracks extend across the transverse ply, and then become partial delamination cracks at the yam interface. At high stresses (Figure 4c) the S-cracks ioin * Catalog No. 12030, L&d Research Industries, Burlington, VT.

Letters

995

to the Editor

Figure 5. Magnified view of S-cracks in a fill specimen (1480°C treatment) at 200 MPa (surface replica).

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developed first along the straight sections of the longitudinal yams (rather than at the cross-over points in the weave). Figure 6 is a sketch of S-cracks in neighbo~ng plies. It was observed that S-crack patterns in neighboring plies can have opposite orientations. It was also noticed that some plies do not contain any S-crack patterns. The explanation for this behavior is indicated in Figure 6, which shows that the longitudinal plies are shearing with respect to their neighbors. Thus the “cross-ply” yarns are subjected to shear stresses at their interfaces, and also to tensile stresses from the applied load. The carbon matrix within the yarns exhibits brittle behavior, and so cracks develop at right angles to the line of principal tensile stress. Apparently the local shear stresses have a much larger magnitude than the applied tensile load, and so these shear stresses strongly influence the crack orientations. Figure 6 shows how the S-cracks open up under an applied shear load. The relative movement of neighbo~ng plies can be in either direction, and it is observed expe~menta~ly that the Scrack patterns show two orientations (see Figure 6). When no S-cracks are visible in a ply, it is presumed that there is no relative shearing at that location. The S-crack pattern varies at different locations in the laminate, and is caused by local differences in the stiffnesses of the longitudinal yarns. These stiffness variations are probably the result of variations in the crimp angles of the yarns. longitudinal

sheared element with tensile forces

Figure 6. Sketch of S-cracks in neighboring plies, and sheared element showing development of S-cracks.

996

Letters to the Editor

It was noticed that the S-crack pattern initiates from the angle crack pattern which is already present in the composite. Those angle cracks which have the right orientation to favor the S-crack pattern will open up, while the others will close. The role of the straight cracks in promoting the development of the S-cracks is less obvious. It appears, as shown in Figure 5, that at lower loads the straight cracks link with pairs of angle cracks to form the “S-crack” pattern. However, we have also observed instances where the straight cracks appear to close up completely, and the S-cracks run at an inclined angle t’joining a pair of angle cracks). The present work shows that the initial crack pattern in a carbon-carbon laminate does influence the progression of tensile damage. The occurrence of Scracks is an indication that at high loads the bonding between plies is breaking down, and the plies are undergoing relative shear deformation (as well as tensile deformations). These relative shear deformations are believed to be caused by stiffness variations in the yarns, which are in turn a function of the crimp angle distribution in the laminate. The local correlation of crimp angles between neighboring plies is thought to be especially relevant here. Future work will address a number of unresolved micromechanics issues, such as the mechanism of angle crack formation and the reduction in failure strain due to high temperature heat ~eatments.

Thermal

diffusivitykonductivity

Acknowledgements

- The authors wish to thank the Phillips Laboratory (AFMC), Edwards AFB, for their support of this work under contract FO4611-88-C-0020. Special thanks is also extended to Mr. Dana Exon of HITCO for his help in preparing the carbon-carbon composites in this study. UDRI, Bldg. 8424 Phillips Laboratory Edwards AFB. CA 93523

P.B. POLLOCK B.J. HINDS R.J. TEDERS C.G. KOCHER

REFERENCES J. Jortner, Carbon G.E. Griesheim. “Notch Strength’ Carbon-Carbon”, 944 (1993). J.E. Masters and

30, 153 (1992). P.B. Pollock and S.C. Yen. and Fracture Behavior of 21c American Ceramic Journal

76,

K.L. Reifsnider, In Damage in Composite Materials, ASTM Publication 775, p. 40-

62 (1980). S.E. Groves, C.E. Harris, A.L. Highsmith, D.H. Allen and R.G. Norvell. Exe. Mechanics 27. 73 (March 1987). ’ ‘

of a compact

of C70 fullerene

(Received 22 March 1993: accepted in revised form 5 May 1993) Key Words - C70 Fullerene, thermal diffusivity, thermal conductivity, specific heat

The thermal conductivity of C&uand C70 single crystals was measured by Yu, et al.[l], and Tea, et al.[2], over the temperature range of about 30 K to about 290 K. At this latter temperature the value for the thermal conductivity of the C&-tsingle crystal was found to be 0.4 W/m.K. The C7u single crystal near 300 K exhibited [2] values of thermal conductivity from approx. 0.7 to 2.0 W/m.K, depending on thermal history. This variation was attributed to the presence of solvents within the C7u single crystal. For this reason the data for the C&u and the C7o cannot be compared directly. Because of the irregular shape of the single crystal of C70 without solvents, Tea et al.[Z], could only obtain numerical data for the thermal conductance, also precluding a direct comparison between the thermal conductivity data for the C&oand C7u single crystals. More recently Withers et al. [3], determined the thermal conductivitv of a compact of Cm, with a density of 1.58 g/cc fapprox. 94.5% of theoretical), from room temoerature to about 600°C. The room temperature value was found to be about 0.21 W/m.K, aboui half the value obtained by Y u et al. [ l] for the Cm single crystal. The difference in these values for the single crystal and compact was attributed to the presence of structural defects and pores in the compact, which lead to a lowering of the thermal conductivity by presenting a barrier to heat flow [4] and/or enhancing phonon scattering[S]. Withers, et al.[3], also found a value of about 0.12 W/m.K at room temperature for the thermal conductivity of a Cm-C70 compact with density of 1.56 -----‘--~

glee con~ning about 92% of Cm and 6% of C70, the remainder consistinn of fullerenes with higher molecular weight. Comparea to the C60 cornpacT, the thermal conductivity value for the Cm-C70 compact was found to be lower by a factor of about two, which at that time was associated with the analogous effect found in metallic or dielectric solid-solutions [3,4]. We wish to report the thermal conductivity values for a compact of C7u molecules and compare these values with those obtained previously for the Cm single crystal and the Cc0 and C@-C70 compacts [1,3]. The C7u was obtained from a mixture of fullerenes made by vaporizing graphite electrodes in an atmosphere of helium atib&i 100 Torr, as described by Ktitschmer f61 et al. The resultine oroduct consisted of about 92% of Cm and 6% of C7u:with the reminder consisting of higher molecular weights of fullerenes. From this mixture the C7o. with a purity of 99.9+%. was obtained by liquid chromatography using alumina as the stationary chase and hexane as the liquid phase. The C7n powder obtained in this manner w& pressed in a steer mold at room temperature to a pressure of approx. 30 MPa into a 12 mm diameter disc with a thickness of about 1 mm. Following cold-pressing, the sample was heated to 300°C at 10-6 Torr to drive off the solvent. The resulting compact had a density of approx. 1.52 g/cc. The thermal conductivitv was determined as follows. The thermal diffusivity”was measured by the flash technique [7,8], using a Nd-glass laser as the flash source, with an estimated experimental error of about c-2