Mechanical tests of a three-dimensionally-reinforced carbon-carbon composite material

MECHANICAL TESTS OF A THREE-DIMENSIONALLYREINFORCED CARBON-CARBON COMPOSITE MATERIAL JOHN L. PERRY Aeronutrcnic

Division, Philco-Ford

Corporation,

Newport Beach. California.

U.S.%

and

DONALDF. ADAMS? Department of Mechanical Engineering. University of Wyoming. Laramie. Wyoming. U.S.h. (Receioed 21 September 1970 hbstrart-A tightly woven block consisting of high modulus graphite fiber yarns in a three-dimensional array ha\ repeatedly impregnated with a series of low viscosity polymer resins, each reimpregnation being followed by a 1540°C carbonization process. After thirteen reimpregnations and carbonizations. the densified block was graphitized at 2650°C. Strength and modulus properties were measured after the seventh and thirteenth carbonizations. and after graphitization, by cutting test specimens from the block. These properties are compared with microphotographs of cross sections of the block taken at various process stages. The influence of the anisotropic thermal expansion characteristics of the composite on the variations in mechanical properties i5 dixu\\ed.

1. INTRODUCTION

Philco-Ford using the shorter duration, lower carbonization temperature process with multiple graphitization cycles. The decision to utilize a 3-D preform block which possessed the same structural configuration a:; that previously processed enabled a direct correlation and comparison to be made between the two densification processes.

Reinforced carbon-carbon materials are rapidly becoming candidates for numerous engineering applications. Unlike polycrystalline bulk graphite, fiber-reinforced carboncarbon composites are responsive to a large array of material, geometric orientation, and processing parameters. Bulk graphite is intrinsically somewhat anisotropic due to its graphite plane alignment. Three-dimensionally woven (3-D) fiber-reinforced carbon-carbon composites more closely approach the orthotropic properties required for service in high temperature environments. The more uniform multi-axial strength of 3-D carbon-carbon coupled with its favorable ablation, thermal shock resistance. and chemical inertness characteristics has given rise to increasing interest in more fully assessing the potential of this material. In an earlier study [ I-31, the behavior of unidirectionallyreinforced carbon-carbon composites was studied. It was concluded, on the basis of both empirical and analytical considerations, that a 3-D carbon-carbon composite might possess superior properties if processed at higher carbonization temperatures than used in the unidirectional composite study, using longer duration thermal processing cycles also. In addition to the longer duration, higher temperature carbonization sequences, it was expected that the carbon-carbon properties could be upgraded by culminating the process with a single, long duration graphitizdtion cycle, rather than the three intermediate graphitization cycles previously used[l/. To effectively evaluate this “extended process method”, a 3-D preform (woven block) was utilized, consisting of Thornel 7% graphite fiber yarns (Union Carbide Corporation[4]) in a weave geometry identical to a block which had been previously processed by ;AIw

Consultant,

Philco-Ford

2. %lATERIAI. DESCRIPTION The 3-D preform block was fabricated for Philco-Ford by Fiber Materials, Incorporated of Graniteville. Massachusetts. It had overall dimensions of 16.0 x 8.6 x 22.1 cm (6.3 x 3.4 x 8.7 in.) in the X. Y and Z directions. respectively. The preform consisted of Thornel 7%. 2 ply graphite yarn spaced on 0.41 mm (0.016in.) centers in the X and Y directions and Thornel 75s. 4 ply graphite yarn spaced in a 0.76 mm (0.030 in.) center to center square array in the Z direction. Although such a preform block is typically called a woven preform. this is somewhat of a misnomer since the individual yarn bundles actually remain straight in the orthogonal pattern. The Thornel 7% graphite fiber is a commercially available fiber produced by Union Carbide Corporation. It is a very high modulus fiber by current standards, having the following typical properties[4]: Tensile strength-2.38 GN/m’ (345,000 lb/in’): Tensile modulus-517 GN/m’ (75 x IO”lb/in’): Density-l .80 g/cm’: Fiber diameter-6 km. The block consisted of approximately 62 X-Y plane layers per inch in the Z direction, a fiber volume of 41.4%. and hence a bulk density of 0.75 g/cm’. The fiber fractions in the X, Y, and 2 directions were 0.24, 0.24. and 0.52, respectively. The volume fraction of fiber per unit cell in the X. Y. and Z directions was 0.13. 0.13. and 0.14, respectively.

Corporation. 61

62

J. L. PERRYand D. F. 3. PROCESSING TECHNIQUE

The method of densification was a modified version of the standard Philco-Ford densification process [l]. It has been demonstrated that density increases are highly predictable using the standard process. Thus, the differences in rate of densification achieved by modifying the process could be readily identified. Densification involved the use of Monsanto’s SC1008 phenolic for the precursor resin matrix. Quaker Oat’s P-3 furfuryl alcohol resin modified by the addition of Dow Chemical’s DEN-438 epoxy novolac (ENF), as well as unmodified P-3 resin, was used for subsequent densifications. The reduction of porosity due to density increases necessitated the use of the progressively lower viscosity ENF and P-3 resin systems for maximum resin penetration of the preform. The properties of these various resin systems are presented in Table 1. Table 1. Propertiesof organicresins utilized

Upon receipt of the woven block, radiographic prints were made of all three axes of the block to assure that the woven preform possessed no obvious anomalies. To remove any moisture from the preform, it was subjected to a temperature of 107°C for 24 hr in an air circulating oven. No weight loss was detected as a result of this drying procedure. The various steps in the densification process were as follows. 3.1 Initial precursor resin impregnation Multidimensionally woven structures are extremely susceptible to deformation when in the unimpregnated condition. To reduce the possibility of deformation during initial impregnation, the block was cradled in a rigid, expanded metal platform, suspended above the container of densifying resin. This initial impregnation cycle is referred to as the “drip-dry” process since it entails subjecting the block to an impregnation using a low solids solution of SC1008 phenolic resin, then suspending and rotating it at 82°C to allow the bulk of the resin to flow out of the block. The purpose of this process is: (1) to promote better filament wetting by the low viscosity resin; (2) to rigidize the woven structure; and (3) to produce a high porosity

ADAMS

structure that will not restrict subsequent resin penetfations. This densification was achieved by subjecting the woven preform to repetitions of vacuum/pressure impregnation cycles using a 30% solids solution of preheated Monsanto SC1008 phenolic resin diluted in methyl ethyl ketone. The block was suspended in a preheated (SO’C) autoclave by a cable which was operable from the outside. Below the block was placed a container of the preheated (66°C) SC1008 phenolic resin. The autoclave was evacuated to 711 mm-Hg (28 in.-Hg) for 30 min to remove air from the block. The block was then lowered into the evacuated resin and the vacuum maintained for 30 min. After releasing the vacuum, 0.69 MN/m’ (lOOpsi) air pressure was applied. This pressure was maintained for 30 min. Repetitions of the vacuum/pressure cycles were performed until the preform was impregnated sufficiently to self-submerge into the container of resin. The impregnated preform was then removed from the autoclave and suspended by the impregnation cradle in a room temperature air circulating oven. The oven temperature was increased to 82°C. After 30 min at 82°C the block was rotated 180”and subjected to an additional 30 min at 82°C. SC1008 phenolic resin becomes less viscous at that temperature and flows out of the preform. This out-flow of resin is desirable since it forms passageways for the escape of volatiles during curing and for more efficient subsequent impregnations. The resin was then cured by elevating the oven temperature to 177°C for 16hr. Bulk density, i.e. resin pick-up, was measured after this cure. The cured preform was then carbonized in a controlled thermal cycle to a maximum temperature of 760°C as described subsequently. 3.2 Precursor resin densification impregnations To achieve maximum resin impregnation after the initial impregnation process, B-staged SC1008 resin was used. In this sequence, a large volume of resin was impregnated into the highly porous block. Briefly, this process is a method of impregnation by applying repetitions of vacuum and pressure cycles to the block while submerged in the preheated, B-staged resin. A B-staged resin refers to the process of partially curing (polymerizing) the resin to increase its viscosity. The SC1008 resin was B-staged for 24 hr at 82°C. The solids content of the resin was increased to approximately 70% by this process. The block was preheated to 100°Cand then suspended above the container of B-staged resin by 2.54cm wide glass straps in the preheated (50°C) autoclave. The block was immersed in the resin and after several repetitions of vacuum and pressure cycles, the heat was increased to reduce the solvent content of the resin to approximately 20%. This reduction of the solvent content is important because it allows a greater volume of resin to remain in the block during cure, and it also reduces the stresses in the block created by volatilization during the curing operation. The resin container was then removed from the autoclave and placed into an air circulating oven at 82°C. The impregnated block, while remaining in the resin

63

Mechanicaltests of a three-dimensionally-reinforced carbon-carboncompositematerial

densification progression. The degree of graphitization achieved is dependent upon time at temperature as well as temperature. The final product achieved in the prrviou\ study was not considered to be a well-~aphitized structure. In the present study, a single, long duration 2650°C graphitization cycle was used as the final matrix conditioning operation, to promote a more ordered 3.3 Initial resin carbonization The 3-D block was carbonized in two steps. Step 1 crystalline structure. The block was graphitized in the entailed preconditioning of the resin matrix to 760°C in a same induction furnace used for the 1540°Cphase ‘,)fthe silica sand sealed, nitrogen atmosphere retort heated by a c~bonization cycle. For comparative purpose:< the cam-controlled oven. The cam-controlled oven has an previously used and the current extended graphiti:~ation operating limit of only 9Oo”C,so carbonization to 1540°C cycles are given in Fig. 2. (Step 2) was performed in a manually operated high temperature induction furnace. This two-step carbonization cycle is plotted in Fig. 1, 2500 with the previously used[ I] single step, low temperature cycle included for comparison.

container, was heated until the resin was advanced su~ciently to not flow during cure. Since the resin was advanced to the no-flow condition, bagging of the block prior to curing was unnecessary. Cure was accomplished by heating the block at 93°C for 30min, then heating at 120°Cfor 30 min, and tinally, heating at 177°Cfor 60 min.

/600 -~ Standard Em----Extended 0

carbonization cycle carbonization cycle

[I]

12w

,,q ,i /

/

i ‘,

/ I To RT at 132 hi

;

96

112

:

I

o-’

0

16

32

48

64

80

Time, hr Time.

Fig. 3.4

1. Carbonization

hr

Fig. 1. Extended

cycles.

Subsequent impregnation and carbonization

graphitization cycle compared used cycle [I],

to preiiously

cycles

The remaining impregnation and cure cycles were similar to the initial resin impregnation cycle except that subsequent redensification cycles required the use of Quaker Oat’s furfuryl alcohol resin modified by the addition of Dow Chemical’s DEN-438 epoxy novolac (ENF), as well as unmodi~ed P-3 resin. The reduction of porosity, due to density increases, necessitated the use of progressively lower viscosity ENF and P-3 resin systems for adequate resin penetration. ENF and P-3 resin systems are of a low enough viscosity at room temperature to not require heating during the impregnation cycles. impregnation of the ENF and P-3 resins was performed by applying repetitions of 760mm Hg vacuum and 0.69MN/m~ (lOOpsi) pressure cycles to the block while being submerged in the resin. The cycle were culminated by applying a 34.5 MN/m’ (SO00psi) pressure to the resin for 2 hr. Both the ENF and the P-3 resin systems were cured in an air circulating oven at 106°Cfor 16hr. The low viscosity ENF and P-3 resin systems required containment during cure to prevent out-flowing. The containment was provided by a silicone rubber box, cast to the approximate dimensions of the block. After each redensification process, the resin was carbonized to 1540°Cusing the same two-step cycle as in the initial carbonization (Fig. 1). The block was subjected to a total of 13 redensificationlcarbonization cycles. 3.5 Graphitization In the previous study [ 11,three graphitization cycies of 4 hr at 2650°C were used at intermediate points of the

4. PROCESSING

RESULTS

The 3-D preform block, prior to initial matrix impregnation, had a buik density of 0.75 g/cm’, corresponding to a fiber volume fraction of 0.414. After subjecting the block to 13 repetitions of the extended reimpregnation/carbonization cycles and a final, singular, graphitization cycte, the bulk density was increased to 1.72g/cm’. After graphitization, the specific gravity was determined by the liquid displacement method. The displacement liquid used was 2,2,4_trimethylpentanc. selected because of its excellent wetting ability. The measured specific gravity was 1.84.The difference in bulk density value vs specific gravity (1.72-I .84) is attributed to the presence of open porosity adjacent to the Z yarns after graphitization, as discussed later. The densification progression of the 3-D block is ,hown in Fig. 3. As indicated, in the early processing stages density increased at an encouraging rate. After about the eighth cycle, however, the rate of density increase per densification cycle was significantly reduced. The increase in density from 1.67 to 1.72g/cm’ after graphitization was a result of volumetric shrinkage; a?: was anticipated, very little weight loss (
64

J. L. PERRY and D. F. ADAMS

appeared to be loosely contained in the matrix due to gaps around their circumference, with the largest gaps being around the Z yarns. After process stage C3 (Fig. .5),the gaps around the Y yarns appeared to be almost completely filled with matrix material, whereas gaps around the Z yarns were only partially filled. It appears that subsequent to the C3 cycle, the mjaor portion of the resin penetration was parallel and adjacent to the Z yarns. After cycle CS (Fig. 6), the matrix density and quality was 080 A075 070’ &v

1 CI

c,

c*

c3 c4

t5

1 1 I I C6 c7 C8 c-9 Cl0 CII Cl2 Cl3

I G

Process stage

Fig. 3. 3-D carbon-carbon

composite stage.

bulk density vs process

Photomicrographs were taken of representative samples of the block at densification (carbonization) process steps Cl, C3, C5, C7 and C13, and after graphitization. Twenty power magnification using oblique (45”), white (a) light was used to study samples at the Cl and C3 conditions. Forty power magnification with direct (O”), white light was used to study the remaining samples. The different lighting techniques were empirically arrived at as providing the best illumination of the irregular surfaces of Y-Yom the matrix for the particular specimens being observed. In End the following discussions, no distinction is implied between the X and the Y directions. The block was fabricated symmetrically in the X and Y directions and remained symmetrical in properties and in visual observations in these directions. For convenience, the photographs are identified by X-Y and X-Z plane views, and (b) reference is made only to Y and Z direction yarns. Due to Fig. 5. Cross sections of 3-D carbon-carbon composite after third symmetry, however, any comments directed toward the Y recarbonization (densification), 20X, oblique lighting. (a) X-Y plane; (b) X-Z plane. yarns apply identically to the X yarns. After process stage Cl (Fig. 4), the Y and Z yarns

Y -

Yarn

Z-Yarn End

End

(a)

Y-Yarn End

(b) Fig. 4. Cross sections of 3-D carbon-carbon composite after first recarbonization (densification), 20X, oblique lighting. (a) X-Y . II \ .I _ 1 plane; to) A-L plane.

(b)

Fig. 6. Cross sections of 3-D carbon-carbon composite after fifth recarbonization (densification), 40X, direct lighting. (a) X-Y plane; (b) X-Z plane.

Mechanical tests of a three-dimensionally-reinforced carbon-carbon composite materlal adequate to respond to 6 p diamond polishing compound, thus permitting the higher magnification (40X) photographs. At the C5 condition, the matrix appeared to have random regions of carbon which responded differently to

the metallographic polishing technique. Also, the large dark areas around the 2 yarns, seen in Fig. 6(a), show the relative size of the densification resin passageways. Figure 7 shows photomicrographs taken after the 7th carbonization cycle. These photographs reveal larger areas of hard carbon which have developed with correspondingly smaller Z yarn circumferential gaps. It is seen that the hard carbon is in the form of ribbon-like bands partially circumventing the Z yarns. The carbon bands are in close proximity to, but not in intimate contact with, the Z yarn surfaces. Sample C7 was machined on the X-Z plane in such a manner as to reveal length-wise segments of 2 yarns. The circumferential gap can be seen to traverse the full length of the exposed Z yarn, creating a continuous passageway for resin penetration (see arrow in Fig. 7bl. As shown previously in Fig. 3, the density progression rate decreased quite rapidly after about the 8th carbonization cycle. This decrease is attributed to the reduction in gap size. Figure 8 shows the matrix formation after the 13th (final) carbonization cycle. In Fig. 8(a), the circumferential gaps around the Z bundles are shown to be extremely small, and are even non-existent in some areas. By comparing Fig. 7(b) to Fig. 8(b), it is evident that the resin passageways in the latter material (the Cl3 condition) have become almost completely filled with carbonized resin. Additional densification cycles beyond Cl3 would undoubtedly have resulted in very little additional density increase. After the final 1540°C carbonization cycle (C13), the block was subjected to the 2650°C graphitization cycle. Note in Fig. 3 that the density increased from 1.67 to 1.72g/cm’ after graphitization. As mentioned previously,

Z-Yarn

z-Yarn

End

(bl

Fig. 8. Cross sections of 3-D carbon-carbon composite after thirteenth (final) recarbonization (densification), 40X, direct lighting. (a) X-Y

plane; tb) X-Z

-

plane.

Z -Yarn

Eld

End

(bi

Fig. 9. Cross sections of 3-D carbon-carbon composite after graphitization, 40X. direct lighting. (a) X-Y plane; (hi X-Z plane.

Resin

Possogeway

Y-Yarn

End

Fig. 7. Cross sections of 3-D carbon-carbon composite after seventh re-carbonization (densification), 40X. direct lighting. (a) X-Y plane:(b) X-Z plane.

the increase was primarily a result of volumetric shrinkage of the overall block dimensions. Figure 9 reflects the resulting fiber-matrix interface after graphitization. Large gaps formed around the 2 yarns, but not around the Y yarns. The gaps around the Z yarn:; were comparable in size to, or even larger than, those observed at cycle C7 (Fig. 7). Note in Fig. 9(b) than many of the Y yarn ends exhibited diametrical separations. The direction of the separation and the absence of a circumferential gap suggests that tensile stresses were induced in the Y yarn bundles in the tZ directions. As previously described. the block used in this study was made of Thornel 755 yarn, using two ends of T 7%

66

J. I... PERRYand D. F.

yarn in the X and Y directions on 0.41 mm (0.016in.) centers, and four ends of T 75s yarn on 0.76mm {O.O3Oin.)centers in the 2 direction. The resuIting average diameters of the two-end and the four-end yarns were approximately 0.36 and 0.51 mm (0.014 and 0.020 in.), respectively. If a transverse, radial expansion of 10.1x 10-6cm/cmi”C is assumed for these yarn bundles, as was experimentally measured for a graphitized, unidirectional T 7SSlSC-1008 carbon-carbon composite in Ref. [I], then the decrease in diameter of these yarn bundIes would be about 0.94 x 10s3cm and 1.35x lob3cm (0.37 x lo-’ and 0.53 x 10m3 in.), respectively, as the temperature was decreased from the 2650°C temperature. The large gaps shown in Fig. 9 were measured to be of the same order. This suggests the ~ssibility that the compressive forces generated by the radial expansion of the yarn bundles during graphitization caused a permanent deformation of the surrounding matrix material. (The coefficient of thermal expansion of the graphitized matrix material is somewhat less than that of the fiber/matrix yarn bundle, having been measured as 3.5 x W6cm/cm~C in Ref. [l].) Figure 10 includes 200X photomicro~aphs taken of the Z fiber bundle to matrix interfaces after the final carbonization (C13) and after the graphitization processes. Note in Fig. 10(a) the existence of closed pores, the intimacy of the Z yarns and surrounding matrix, and the generally unconsolidated appearance of the matrix. By subjecting the structure to the 2650°C ~aphitization temperature, the matrix underwent considerable change, as shown in Fig. 10(b). The compressive forces applied by the radial expansion of the 2 yarns apparently collapsed the spherical, closed pores, producing a much more tightly compacted matrix, and creating the circumferential gaps. Note the replication of the expanded fiber bundle in the matrix in Fig. IO(b).

ADAMS

Figure 1I includes 200X photomicrographs taken in the Y direction after Cl3 and after ~aphitization. Cireumferential gaps were not formed. instead, di~etrical cracks were formed within the individual yarn bundles, as shown in Fig. 11(b). Examination of these photomicrographs suggests that adherence of the matrix to the Y yarn surfaces was sufficient to prevent the formation of ribbon-like gaps during cooldown. Note that the unreinforced matrix surrounding the bundles would be expected to be stronger than the matrix within the yarn bundles, because of the stress concentrations created by the individual fibers within the yarn bundles. Diametrical separations within yarn bundles were not common only to the graphitization cycle. Separations were visible in the Y yarn bundles as early in the processing as C? (see Fig. 7b). In Fig. 1I(b), two distinct separations are indicated. One was formed and subsequently healed by redensification, while the other appears to be a new separation, formed during the final processing.

(bf

Fig. 11.Cross sectionsof 3-D c~~n~~~n ~orn~s~~enormal to Y-direction,200X.direct lighting. (a) View of matrix after thirteenthre-den&cation; (b)viewof matrixaftergraphitization. 5. MATERIALS TESTING

Test specimens were machined from the 3-D block at three stages in the processing: (1) after the seventh (midway) densification cycle, (2) after the thirteenth (final) densification cycle; and (3) after the graphitization cycle. Three tensile and three short beam flexure specimens in both the Y and 2 directions were machined from the block at each of these three processing stages. Small coupons for use in deter~ning the coefficients of thermal expansion were also removed after the ~ap~tization process.

Fig. 10. Cross sectionsof 3-D carbon-carbon composite normal to Z-direction, 200X, direct lighting. (a) View of matrix after thirteenth re-densification; (b) view of matrix after graphitization.

Tensile tests The tensile specimen configuration is shown in Fig. 12, As will be discussed, the strength properties of the 3-D carbon-carbon composite were highly anisotropic. It was found to be necessary to reduce the width of the Z-direction specimens from the original 1.91cm (0.75in.) width to 0.84 cm (0.33 in.) to prevent a shear failure in the

Mechanical tests of a three-dimensionally-reinforced

Flexure

I”,1 2.54

4

/.27/+

1’

’ , C-f.12

‘\

R

I

@Eprbond @

/22/952

Alumrnom

and

Rolled

Steel

Scram

Clofh

Drrection

Tens)/0

@Z

Directton

TenslIe

tests

3-D composite along the three orthogonal planes. Short beam specimens each 4.06 cm (1.60 in.) long, 1.27 cm (0.50 in.) wide, and 0.64 cm (0.25 in.) thick were selected for this purpose, the Z-axis fibers being oriented in the length, width, and thickness directions of the beam, respectively, to obtain the three shear stress loading conditions desired. The specimens were supported on a span of 3.56cm (1.40 in.). and loaded through a single load point at mid span. The span length to specimen thickness ratit,, WI\ thus 5.6 to I, which was selected as being low enough to induce a shear failure in the composite material. However, the 3-D composite actually failed in a typical flexure mode, i.e. either a tensile or a compressive fracture or some combination of both. Thus. a shear strength was not determined. Rather, these tests were used to determine the flexural strength and modulus of the material in the three orthogonal directions. Thermal

(Reduced

67

Adhesive

Spacer

@V

Fig.

Cold

104 Gloss

composite material

It was desired to determine the shear strengths of the

i +-084c~ I

@O 476 f3/16 inch1 Ground

carbon-carbon

Width)

I?. Tensile specimen geometry (dimensions in cm)

end tab regions prior to tensile failure in the gage length. This specimen modification is also indicated in Fig. 12. Electrical resistance (foil) strain gages were bonded to opposite sides of the specimen in the gage length region; to measure the axial strain, and also to detect any bending effects due to loading misalignments. The specimens were tested in an lnstron Universal testing machine at a crosshead travel rate of I.3 mm/min (0.05 in/min). .A double set of universal joints were used to load the specimens through the holes in the end tabs.

expansions

The coefficient of thermal expansion of the 3-D carbon-carbon composite was determined after the graphitization process. It was measured in both the Y and Z directions, using 0.64~064~ 3.81 cm (0.25 x 0.35 >: 1.5 in.) coupons heated to a maximum temperature of approximately 1000°C at a rate of 4°C per min. Measurements were made using a Model HTV Leitz Diiatometer. 6. TEST RESULTS Results of all of the uniaxial tensile tests are presented in Table 2, along with the average values of each set of three tests and the coefficient of variation. Certain of the individual tests produced values somewhat higher or lower than the average, but since only three data points were available. all values were retained in computing the averages.

Table 2. Tensile properties of 3-D carbon-carbon

composite

68

J. L. PERRY and D. F. ADAMS Table 3. Flexural properties of 3-D carbon-carbon composite

Strength

No.

Length Of Beam

StrengCh

Modulus

lo3 IO6 ml/m2lb/in2 cwm2 1b1i”2

muIn

Modulus

lb;:2 GN/m2lbcn2

1

197

28.5

24. I

3.5

261

37.8

15.8

2.3

109

15.8

24.8

2

205

29.1

14.4

2. I

250

16.2

14.4

2.1

105

15.2

21.3

3.1

3

185

26.9

25.5

3.7

243

35.2

IS. 1

2.2

100

14.5

19.3

2.8

19.6

28.4

21.3

3. L

251

36.4

15.1

2.2

105

15.2

2L.8

7.9

1.2

4.9

0.9

7.6

1.3

0.56

0. 1

3.1

0.6

Average tandard vcviation ---.-

SfrengLh

IO3 IO6 Eiwm2lb/ln2 cwrn2 lb/in2

specimen

Z-Yarn lrientatio”

Modulus

3.6

3.2 0.4

2.3

-

Width Beam

af

1

35

5. 1

15.8

2.3

87.5

12.7

21.3

3.1

2

63

9.2

11.0

1.6

79.9

11.6

19.3

2.8

3

47

6. 8

6.8

1.0

83.4

12.1

11.9

2.6

48

7.0

11.2

1.6

83.6

12.1

19.5

2.8

13

2.1

3.7

0.6

3.10

0.6

1.39

0.2

.verage tandxd

"="iatio" n__

Depth Beam

of

1

32

4.6

4. L

0.6

51.3

11.8

16.5

2.4

93.7

13.6

15.8

2

44

6.4

4.8

0.7

106

15.4

19.3

2.8

77.2

11.2

17.2

2.5

3

30

4.3

5.5

0.8

75.8

11.0

16.5

2.4

95.1

13.8

17.9

2.6

x5

5. I

4.8

0.7

87.7

12.7

17.4

2.5

88.6

12.9

16.9

2.5

1. 1

P.56

0. 1

13.2

2.3

1. 31

0.2

8.1

1.4

0.9

0.2

Average Standard Deviatlo”

6.4

Table 3 is a similar tabulation of the flexural test data. As previously noted, these specimens consistently failed in a flexural mode, viz. in tension at the tensile surface. However, the fracture plane did not always occur at midspan. Also, the graphitized specimens having Z-yarn bundles oriented along the length of the beam exhibited bundle movement at fracture, as indicated in Fig. 13. This is a view of one of the machined ends of a specimen, and shows protruding fiber bundles at the compression side and retracted fiber bundles at the tension side. This yarn bundle movement did not occur in the other two Z-yarn orientation, graphitized test specimen configurations. The implications of this will be discussed in subsequent paragraphs. The thermal expansion response of the 3-D composite after graphitization, in both the direction of the Z-bundles and in a transverse direction (the Y-direction), is indicated in Fig. 14. As can be seen, the response was the same in both directions, exhibiting a negative response up to about 725”C, and a positive value at higher temperatures. In general the tensile strengths are consistent with the results obtained from similar 3-D carbon-carbon composites[5], with the exception of the Z-orientation, Protruding

Fiber

Bundles

\

Tensile

Retracted

Side

Fiber

Bundles

Fig. 13. View of machined end of graphitized 3-D carbon-carbon composite flexure specimen after failure.

2.3

k

0 0

100

200

200

400

300 600

400 800

500 1000

600

700 1’200

800 1400

900 1600

1000 -c 1800°F

Temperature, “C and “F Fig. 14. Thermal expansion (room temperature to 982°C)of 3-D carbon-carboncomposite aftergraphitization.

graphitized material tests. Figure 9(a) illustrates the existence of circumferential gaps around the Z yarn bundles. The effect of the gaps is illustrated in Fig. 15. Of the 99 fiber bundles in the gage length cross section of this particular specimen, 68 have pulled out through the entire length of the tab region. Obviously, the tensile strength values given in Table 2 are, in reality, only a measure of the combined contributions of the yarn bundles that did fail, the X and Y yarn to matrix bond transverse strengths, and the frictional force between the surrounding matrix and the Z yarn bundles that pulled out. Fiber pullout of this magnitude was achieved only in the three Z-direction, graphitized specimens. Failures typical of the C7 and Cl3 carbonized structure in the Z-direction are shown in Fig. 16. The Z yarn bundles protruding from the failed specimens provided material for tensile tests of the individual yarn bundles. Yarn bundles were removed from the specimen shown in Fig. 15. This specimen contained 99 Z yarn bundles in the gage length total cross-sectional

Mechanical

tests of a three-dimensionally-reinforced

”

(b)

Fig. 15. Typical failure of graphitized Z-direction tensile specimens showing yarn bundle pullouts. (a) Side view of failed specimen:(b) rotated end view.

carbon-carbon

composite

69

material

individual Z yarn bundles was I.12 GN/m’ (162.ksi). based upon a 202 N (45.5 lb) load and an average Z yarn bundle diameter of 0.483 mm (0.019 in.). Although sufficient material was not available for additional densification steps beyond the graphitization, it is likely that further processing would have increased the Z-direction tensile strength of the material to at lea,st the theoretical value of 296 MN/m’. As previously described, the fiber volume fractions in the three orthogonal directions of the 3-D block tested were approximately equal. Since tensile strength is s. fiber dominant property, it would thus have been expected that the Y-direction tensile strength would be about equal to the Z-direction tensile strength. However. the highest average tensile strength in the Y-direction (60.7 MN/m’ after graphitization) was only about one-third as high as the highest Z-direction tensile strength (192 MN/m’ after the 13th densification). Figure I7 is a photomicrograph of a Y-Z plane CJfthe 3-D block after the 13th densification and prior to testing. Clearly a reduction in tensile strength would be expected due to the nature of the broken Y direction yarns !,hown in the figure. Note in Fig. 17 also the cyclic nature of the broken Y direction yarns, as indicated by the broken line. It is not likely that yarn damage such as this could have occurred during the vacuum/pressure impregnations, or from pressures generated during curing, carbonization, or graphitization. Rather, this damage was probably the result of Z yarn placement during the initial preform weaving operation. Although Fig. I7 illustrates one of the worst cases of X and Y yarn breakage observed, varying degrees of damage were noted throughout all of the X-Z and Y-Z surfaces. In all instances damage was found to have occurred to only the X and Y yarns: no Z yarn damage wa$ observed. Damage more typical of that throughout the structure is shown in Figs. 4(b) and 6(b). The excessively low Y-direction tensile strengths and elongations reported for the 13th densification (Table 3) are thought to be the result of a particularly high concentration of broken yarns within the region of the block from which these particular test specimens, were taken. After the third carbonization cycle, visual inspection revealed that a separation had occurred on the outer Y-Yarn $

Path

X-Yarn

Fig. 16. ‘Typical

failure of Z-direction tensile tested after C7 and (213.

strength

End

specimens

area of 67.7 mm’ (0.105 in’). The average tensile breaking force of the three Z yarn bundles tested was 202 N (45.5 lb). Based on these values, a theoretical tensile strength of the 3-D block in the Z direction would be 296 MN/m* (42.9 ksi). This value can be considered to be a minimum because the Z yarn bundles which were tested may have been degraded while being pulled out of the original test specimen. The average tensile strength of the

Fig. 17. Photographs densification

of a Y-Z plane showing broken

cross section Y yarns. 20X

atter

13th

70

J. L. PERRY and D. F. ADAMS

surface of the block. Radiographic inspection in the three axes further revealed two primary (0.76 mm wide) and several secondary (-0.25mm wide) separations in the X-Z plane with some separations of lesser magnitude in the Y-Z plane. Radiographic inspection was originally planned to be performed at only the as-woven and the specimen removal sequences in the process cycling. As a result, it cannot be determined at what precise stage in the process cycling the separation occurred. However, the block was radiographically inspected after both the 760°C and the 1540°C steps of all subsequent carbonization cycles. The two primary separations increased in width in small incremental steps. The dimensions of the two separations appeared to stabilize at a maximum of 1.52mm after the seventh carbonization cycle. Secondary separations expanded by no measurable amount. Before machining test specimens from the block at the C7 condition, a 6mm thick cross-sectional specimen was excised from the center of the block to examine the extent of the separations. The separations were localized; large unaffected areas provided ample material for test specimen removal. The existence of the separations did not prohibit the successful completion of this particular study; however, large scale hardware applications would not be possible if such separations in 3-D materials were commonplace. Thus, an additional investigation was undertaken to attempt to isolate the causes of the separations [5]. 7. CONCLUSIONSAND RECOMMENDATIONS The use of an extended carbonization and graphitization process such as presented here appears to have a potential for yielding 3-D structures with satisfactory mechanical properties and low, symmetrical thermal expansions in the X, Y and Z directions. This conclusion is based on the following observations: (1) Thermal expansion of the 3-D block was approximately equal in the X, Y and Z directions when measured to the maximum test temperature of 982°C. (2) The Z direction tensile strength was shown to have a theoretical value of approximately 296 MN/m*. Circumferential gaps were formed around the Z fiber bundles due to radial expansion of the yarns during the 2650°C graphitization cycle. Additional densifications should enhance the fiber to matrix adhesion, thus allowing the fuller utilization of the strength of the Z fiber bundles. (3) The particular Z direction yarn bundle quantities used here in combination with particular center-to-center spacings appeared to yield a preform with broken and disoriented X and Y fiber bundles. Photomicrographs revealed that X and Y direction fibers were disoriented and damaged. The yarn paths were shown to have a cyclic nature, suggestive of damage caused by neighboring Z direction yarn, or bobbin, movement during weaving. The excessively low strengths reported for the Y direction tensile tests are attributed to the discontinuous, and sometimes broken, Y fiber bundles. It is suggested that the properties of the 3-D carbon-

carbon material would be improved by using a modified version of the extended process cycle used in this study, in conjunction with a woven preform geometry that did not inherently contain damaged fibers. Recommendations to upgrade the properties of the 3-D Thornel 75s block include the following: (1) Additional efforts should be directed toward developing analytical methods of selecting preform geometries that would more fully utilize the full strength potential of the fibers. (2) The use of the longer duration, higher temperature carbonization cycle is recommended. Benefits gained from the extended cycle were, (a) more rapid density increases in the initial few process cycles, (b) more complete volatilization prior to graphitization. (3) Future organic resin impregnation pyrolysis processes should incorporate several low solids solution impregnations as the initial preliminary densification steps. The low solids impregnations would decrease volatilization pressures generated during cure. (4) The singular, final graphitization process step is recommended in future densification processes. The singular graphitization cycle did not promote the formation of circumferential gaps around the X and Y fiber bundles. If a woven preform geometry is selected that possesses Z direction fiber bundle quantities, for example, similar to the X and Y fiber bundle quantities used in this program, circumferential gaps should not form and full fiber strength utilization should be attained. On the other hand, if circumferential gaps are formed after graphitization, re-densification may be necessary to re-establish the fiber to matrix relationship. (5) Recognizing the very practical problem of impregnating a 3-D weave block, and the corresponding necessity of maintaining a low viscosity impregnant, it is recommended that emphasis should continue on the search for a higher carbon yield precursor matrix. Polymers having carbon yields in excess of 90% have been synthesized[6], but these systems are typically not sufficiently processible for the present application as of yet. With the present need for a low viscosity, high carbon

yield precursor matrix in mind, work should be encouraged in their continual development. Acknowledgements-This investigation was conducted under the sponsorship of the REVMAT Program of the United States Naval Ordnance Laboratory, White Oak, Silver Spring, Maryland, under the technical direction of C. R. Rowe and F. .I. Koubek. The authors wish to also acknowledge the assistance of their colleagues, J. L. Kirkhart of Aeronutronic and A. K. Miller of the University of Wyoming. REFERENCES 1. 2. 3. 4.

Perry J. L. and Adams D. F., J. Mater. Sci. 9, 1764 (1974). Adams D. F., J. Camp. Mater. 8, 320 (1974). Adams D. F., Mater. Sci. and Engng 17, 139 (1975). Technical Information Bulletin, Thomel 75s Graphite Yarn. Union Carbide Corporation, New York (1974). 5. Adams D. F., Mater. Sci. and Engng 23, 55 W76). 6. Neuse E. W., Mater. Sci. and Engng 11, 121 (1973).