Tensile Strength and Fracture Toughness of C-C

Available online at www.sciencedirect.com

www.elsevier.com/locate/compositesa

Composites: Part A 39 (2008) 231–242

Tensile strength and fracture toughness of C/C and metalinfiltrated composites Si–C/C and Cu–C/C

Wen-Xue Wang *, Yoshihiro Takao, Terutake Matsubara

Research Institute for Applied Mechanics, Kyushu University, Kasuga-koen 6-1, Kasuga, Fukuoka 816-8580, Japan

Received 6 June 2007; received in revised form 28 September 2007; accepted 3 November 2007

Abstract

In this paper, tensile strength and mode I fracture toughness of a low cost carbon–carbon composite (C/C) and liquid metal infiltratedcomposites Si–C/C and Cu–C/C are investigated based on tensile testing and compact tension testing. Si–C/C and Cu–C/C are manu-factured by infiltrating liquid silicon and copper into the low cost C/C, respectively, to make the densification of porous C/C, and toimprove the oxidation resistance and the frictional and wearing properties of C/C for various potential applications under high temper-ature and frictional environments. Experimental results and analysis of microstructures of three composites reveal that the infiltration ofliquid metal into C/C significantly influences the tensile strength and fracture toughness of the composite. The tensile strength and frac-ture toughness of Si–C/C are much lower than original C/C. On the other hand, infiltration of copper into C/C significantly improves thestiffness and strength of Cu–C/C.� 2007 Elsevier Ltd. All rights reserved.

Keywords: A. Carbon–carbon composites; B. Strength; B. Fracture toughness; E. Liquid metal infiltration

1. Introduction

Carbon–carbon composite (C/C) is an advanced heatresistant material and has been widely used in structuresof space vehicles, nuclear reactors, aircraft brake, and rac-ing car brake due to its light weight, high strength andtoughness at very high temperature [1,2]. Although C/Chas many advantages, it still has some drawbacks. One ofthe drawbacks is related to its long manufacturing process.In general, two methods of chemical vapor deposition(CVD) and organic carbonization are utilized to manufac-ture C/C and both of them require long time and repeatedprocesses for the densification of the C/C. Consequently,the long manufacturing process leads to quite high produc-tion cost, which limits its usage to some special industrialfields as mentioned above. In addition, the oxidation ofC/C at high temperature of above 600 �C is also a problem.

1359-835X/$ - see front matter � 2007 Elsevier Ltd. All rights reserved.

doi:10.1016/j.compositesa.2007.11.004

* Corresponding author. Tel.: +81 92 583 7757; fax: +81 92 583 7760.E-mail address: [email protected] (W.-X. Wang).

To improve the oxidation resistance, reduce the productioncost of conventional C/C, and extend the application ofC/C to general industrial structures working at high tem-perature, a low cost C/C manufactured by a preformedyarn method [3] and liquid metal infiltrated C/C compos-ites such as Si–C/C and Cu–C/C were developed [4–6].These low cost composites provide new candidates of heatresistant materials for space vehicles, aircraft brake, alsobrakes of automobile and train cars, and other structuresworking at high temperature. Hatta and his colleagues[7–9] conducted a series of researches on the fracture andfatigue of the low cost C/C. Matsubara et al. [10] investi-gated the effects of contact pressure and reciprocationnumber on the coefficient of friction of C/C, Si–C/C andCu–C/C. Si–C/C appeared a larger friction coefficient thanC/C and Cu–C/C, and Cu–C/C shown the best wearingresistance among these three composites. These resultsare certainly interesting for engineers to design brake discor other wearing elements. Moriyama et al. [11] investi-gated the strength of C/C, Si–C/C and Cu–C/C under ten-sile cycle loading and found that the cyclic strength of these

mailto:[email protected]

Ai: (Pi,δi,ai)

Ai+1: (Pi+1,δi+1,ai+δa)ad, P

Ai: (Pi,δi,ai)

Ai+1: (Pi+1,δi+1,ai+δa)ad, P

232 W.-X. Wang et al. / Composites: Part A 39 (2008) 231–242

materials at 106 cycles was above 70% of each ultimate ten-sile strength. On the other hand, the tensile strength andfracture toughness of these composites are also two impor-tant material properties to the design of brakes. Crackingfrom the corner of slot area of a disk is one of the majordamages. However, the study on the effects of the infiltra-tion of metals into C/C on the tensile strength and fracturetoughness of Si–C/C and Cu–C/C is not sufficient up tonow.

In the present paper, analysis of microstructures, tensiletesting and compact tension (CT) testing were conducted toinvestigate the tensile properties and fracture toughness oflow cost C/C and metal infiltrated composites Si–C/C andCu–C/C. Attention was focused on the effects of metal

6

L

20

(a) Specimen

40

1.0

0.3

r

(b) Specimen for

L=140 for C/C and Si-C/C, L

Fig. 1. Geometries of specimens for tensil

Table 1Material properties of C/C, Si–C/C and Cu–C/C composites

Material Density(g/cm3)

Porosity(%)

Coefficient offriction(P = 10 MPa)a

Interlaminar shearstrength (MPa)

C/C 1.7 20 0.25 4 ± 0.3Cu–C/C 2.68 9 0.23 14.1 ± 1.2Si–C/C 2.05 5 0.22 23.3 ± 1.1

a The coefficient of friction varies with contact pressure P.

infiltration on the microstructures, tensile properties andfracture toughness of these composites. Microstructureswere analyzed using optical microscope, scanning electronmicroscope (SEM) and electron probe microanalyzer(EPMA).

0

4Al

for tensile test

50

=10

48113a

(mm)

compact tension test

=130 for Cu-C/C; size in mm

e testing and compact tension testing.

O δi δi+1

LoDisplacement, δ

δ(i+1)rδirO δi δi+1

LoDisplacement, δ

δ(i+1)rδir

Fig. 2. Determination of critical energy release rate from load–displace-ment curve.

W.-X. Wang et al. / Composites: Part A 39 (2008) 231–242 233

2. Experimental

2.1. Materials and specimens

The low cost C/C used in this study was produced byAcross Co. Ltd. (Saitama, Japan) based on the preformedyarn method, and the liquid metal infiltrated compositesSi–C/C and Cu–C/C used in this study were manufacturedby NGK insulators, Ltd. (Nagoya, Japan) based on the lowcost C/C. These composite laminates consisted of 0� and

Fig. 3. Microstructures of C/C, Si–

90� layers alternatively and the thickness of each layer isabout 0.4 mm. Preformed yarn was a bundle of carbonfibers together with carbon binders wrapping by a polymertube. The preformed yarns with fillers were then woveninto a prepreg-like sheet using polymer threads, thesesheets were stacked into a designed lamination pattern,and stacked layers were finally sintered using a hot pressmachine. This manufacturing process of C/C is quite sim-pler than the manufacturing process of conventional C/Ccomposites and it leads to a low cost C/C. However, the

C/C and Cu–C/C composites.


low cost C/C usually contains about 20% pores. To reducethe porosity and to improve the oxidation resistance, fric-tional properties, and wearing resistance, liquid silicon orcopper was infiltrated into the pores of the low cost C/C

Fig. 4. SEM images of impregnated Si and Cu,

laminate to form Si–C/C or Cu–C/C composite laminate[6]. The infiltration of Si into C/C is performed under theconditions of about 1600 �C, 1 hPa low pressure, and Argas atmosphere for about 1 h, while the infiltration of Cu

and EPMA mapping of carbon of Si–C/C.


into C/C is performed under the conditions of about1200 �C, 10 MPa pressure, and Ar gas atmosphere forabout 20 min. Details of these infiltration processes canbe found in Ref. [6]. It is noted that a potential applicationof Cu–C/C composite to the brakes of trucks and train carsis specially expected since the working temperature of thesebrakes may be lower than the melting point of Cu. Table 1gives several material properties of these composites. It isseen that the porosity was reduced to 5% for Si–C/C and9% for Cu–C/C. On the other hand, the density increasedby about 20% for Si–C/C and 57% for Cu–C/C. The inter-laminar shear strength (ILSS) was obtained from doublenotched compressive shear testing [11].

Specimens used in the present experiments were cutfrom the laminates of C/C, Si–C/C and Cu–C/C manufac-tured by above mentioned methods. Geometries of speci-

Fig. 5. In situ CCD camera video images of three com

mens for mechanical tests are described in Fig. 1.Specimens for tensile testing were rectangular, as shownin Fig. 1a. The length of tensile specimens of C/C andSi–C/C was 140 mm and the length of tensile specimensof Cu–C/C was 130 mm due to the limitation of originallaminate of Cu–C/C. The gauge length of all specimensof three composites was the same as 60 mm. Aluminumtabs of 1 mm thickness were bonded to the two ends ofspecimens to protect the grip regions. A strain gauge waslocated at the center of gauge length of each specimen tomeasure the strain. Fig. 1b shows the geometry of speci-mens for compact tension testing. The thickness of all spec-imens of three composites was 4 mm. The width ofspecimens is denoted by W and was 40 mm which is equalto the distance from the loading line to left edge. The lengthof pre-crack is denoted by a which is equal to the distance

posites at tensile failure, and fractured specimens.

0

50

100

150

200

250

300

0 0.1 0.2 0.3 0.4 0.5 0.6

Strain (%)

Str

ess

(MP

a)

0

50

100

150

200

250

300

350

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7

Strain (%)

Str

ess

(MP

a)

0

10

20

30

40

50

60

0 0.05 0.1 0.15 0.2

Strain (%)

Str

ess

(MP

a)

C/C

Cu-C/C

Si-C/C

Fig. 6. Typical stress–strain curves of C/C, Cu–C/C and Si–C/Ccomposites.


from loading line to crack tip. Pre-crack was made firstlyusing a cutting machine with a 0.3 mm thick blade and thenusing a razor blade to obtain a relatively sharp crack tip.

2.2. Testing procedures

Analysis of microstructures of three composites was per-formed using optical microscope, scanning electron micro-scope (SEM) and electron probe microanalyzer (EPMA)before mechanical tests to investigate the effects of metalinfiltration on the original C/C.

Tensile testing was performed using a MTS 810 materialtesting system at room temperature. Six specimens weretested for C/C and five specimens were tested for each ofSi–C/C and Cu–C/C. Loading rate was 0.5 mm/min. Strainwas obtained from the strain gauge located at the center ofeach specimen. One side edge of each specimen was pol-ished. Optical observation of side edges of specimens wascarried out before and after testing to identify the damageevolution. CCD video camera system was also used toobserve the damage evolution during loading.

Compact tension testing for three composites was con-ducted also using MTS 810 material testing system at roomtemperature to evaluate the mode I fracture toughness.Loading was applied through a pin-and-shackle systemand loading rate was also 0.5 mm/min. It is well knownthat damages, such as fiber breaking, matrix cracking, fiberpull-out, debonding, and delamination, might occurmixedly around the crack tip when a crack grows in C/Ccomposites. Thus it is difficult to clearly define the cracklength during testing. In the present work, an effectivecrack length approach [12] using the compliance calibra-tion curve of each composite was adopted to determinethe effective crack length. Therefore, the present compacttension testing contained two procedures. Firstly mono-tonic tensile loading procedure was performed for CT spec-imens with various pre-crack lengths to obtain thecompliance calibration curve of each composite. Then arepeated loading/unloading procedure was applied to eachCT specimen under displacement control to obtain theload–displacement curves. Seven CT specimens were testedfor C/C and five CT specimens were tested for each of Si–C/C and Cu–C/C. The length of the pre-crack was about4 mm for Si–C/C specimens, 6 mm for Cu–C/C specimens,but 20 mm for C/C specimens because the loading regionfailed before the crack growth for smaller pre-cracks. Thisfact proved that the low cost C/C is a very tough materialcomparing other two composites. A theory of linear elasticfracture mechanics applied to composites [13] was adoptedto approximately evaluate the fracture toughness of presentthree composites, although the crack growth processinvolving with multi-damages was quite complicated.Fig. 2 schematically describes the determination of criticalenergy release rate from the load–displacement curve. Theshaded area represents the energy contributed to the crackgrowth, whereas the white area is considered to be con-sumed by other damages and hysteretic energy loss during

a cyclic loading. The critical energy release rate is expressedby

GC ¼P iCDCi

2BDai; DCi ¼ Ciþ1 � Ci; Ci ¼

di � dir

P iCð1Þ

where Ci represents the compliance at loading Pi during acyclic loading, PiC denotes the critical load, Dai denotesthe effective crack extension, and di and dir represent thedisplacement and residual displacement, respectively. Effec-tive crack length ai was determined from the compliance Ci

and compliance calibration curves as mentioned above.


3. Experimental results and discussion

3.1. Microstructures

Optical microscope images presented in Fig. 3 show thepolished side surfaces of C/C, Cu–C/C, and Si–C/C lami-nates. Global images through the laminate thickness aregiven in Fig. 3a and enlarged local images are given inFig. 3b. Dark spots and crack-like strips represent poresand voids. Many pores and voids are observed in the caseof C/C laminate, whereas a few pores and voids can be seenin the cases of Cu–C/C and Si–C/C. These facts reveal thatmetal infiltration indeed is an effective method for the den-sification of porous C/C composite. On the other hand, it isworth noting that new cracks appeared between every twooriginal cracks after infiltration process in the case of Si–C/C, though the original cracks were infiltrated with silicon.This kind of new cracks can not be observed in the case ofCu–C/C. One of the reasons for the difference of microstruc-ture between Si–C/C and Cu–C/C is due to the influence ofinfiltration process as mentioned in the above section.Besides, it is noted that pores and voids equally existed in0� layers of laminates of three composites, though the voidsin 90� layers seem to be more obvious than those in 0� layers

Fig. 7. Side views of damage evolutio

in Fig. 3. To see the detail morphology of infiltrated metals,microscopic analysis of Si–C/C and Cu–C/C was conductedusing SEM and EPMA. SEM images of infiltrated siliconand copper are shown in Fig. 4a and EPMA mapping of car-bon of Si–C/C is shown in Fig. 4b. It is seen that a siliconcarbide (SiC) interlayer exists between silicon and surround-ing C/C due to the reaction between silicon and carbon athigh temperature during the infiltration process. This factis proved by the EPMA mapping of carbon of Si–C/C, asshown in the left one of Fig. 4b. In contrast, there is noany reacted interlayer between copper and carbon in Cu–C/C. The formation of SiC interlayer between silicon andsurrounding C/C leads to a high ILSS as given in Table 1since the infiltrated silicon formed some networks in C/Cfrom the view point of three-dimension and the SiC inter-layer induced a strong combination of carbon and silicon.However, high ILSS is not always an advantage for com-posites and it may lead to low strength and fracture tough-ness of composites, which can be realized from followingexperimental results. Furthermore, as shown in Fig. 4c,some of carbon fibers near to the infiltrated Si seem to beeroded by the chemical reaction between silicon and carbonfiber. This degradation of carbon fibers may reduce themechanical properties of carbon fibers.

n in C/C and Cu–C/C specimens.

Fig. 8. Side views of damage evolution in Si–C/C specimens.

279.4±15

54.1±10 62.5±14 57.8±1055.2±7

238.7±19

0

50

100

150

200

250

300

350

C/C Cu-C/C Si-C/C

Ave

rage

d va

lues

of t

ensi

le s

tren

gth

and

stiff

ness

strength (MPa)

stiffness (GPa)

Fig. 9. Averaged stiffness and strength of C/C, Cu–C/C and Si–C/C composites.


Si-C/C

200

400

600

800

1000

1200

1400

Load

P (

N)

C/C

00 321

200

400

600

800

1000

1200

1400

Displacement (mm)

Load

P (

N)


3.2. Tensile testing results

Typical in situ video images of specimens at tensile fail-ure are presented in Fig. 5a, and fractured specimens areshown in Fig. 5b. Three composites appeared quite differ-ent failure behaviors. Fracture of 90� layers and large scaledelamination are observed in the case of C/C specimens,relatively small scale delamination together with fractureof 90� layers is observed in the case of Cu–C/C specimens,whereas only fracture of 90� layers can be observed in thecase of Si–C/C specimens before the final fracture of 0� lay-ers. In other words, no obvious delamination can beobserved from the Si–C/C specimens. In addition, plasticdeformation due to the effect of infiltrated copper can beobserved from the image of Cu–C/C specimen. These fea-tures also appear in the typical stress–strain curves of spec-imens of three composites, as shown in Fig. 6. Stress–straincurves of specimens of C/C and Cu–C/C are linear orslightly nonlinear until their critical loads, respectively.After critical load and before final fracture, the decreaseof stress in the C/C specimen is induced by the large delam-ination occurring and partially breaking of 0� layers,whereas a slight decrease of stress in the Cu–C/C specimenreveals a pseudo-plastic behavior since Cu is ductile andthe interface between Cu and carbon has very weak bond-ing. On the other hand, the stress–strain curve of Si–C/Cspecimen becomes obviously nonlinear as the load increas-ing, and the final facture occurs almost at critical load. Thecritical strain of Si–C/C is also much smaller than those ofC/C and Cu–C/C specimens. The nonlinear behavior isconsidered to be induced by the cracking in the regionsaround pores and voids during the loading, which is con-firmed by the optical microscope observation of side edgesof specimens before and after testing, as shown in the fol-lowing figures. Elliptic and circular circles marked in Figs.7 and 8 denote the new damages occurred during tensiletesting. Typical images of C/C and Cu–C/C specimens

0.0E+00

5.0E-07

1.0E-06

1.5E-06

2.0E-06

0 0.2 0.4 0.6 0.8a/W

Com

plia

nce

(m/N

)

C/C

Si-C/C

Cu-C/C

Fig. 10. Compliance calibration curves of C/C, Si–C/C and Cu–C/Cspecimens.

before (left) and after (right) tensile testing are presentedin Fig. 7. A delamination and a coalition of two crackscan be seen in C/C specimen, and a debonding betweeninfiltrated copper and surrounding C/C is observed inCu–C/C specimen. Fig. 8 presents typical images of Si–C/C specimens before (left) and after (right) tensile testing.Comparing with the specimens of C/C and Cu–C/C, manycracks are observed. These cracks certainly lead to theincrease of strain of the specimen during tensile testing,which finally results in a nonlinear stress–strain curve forSi–C/C specimen. The lower and right one in Fig. 8 also

00 0.5 1 1.5 2

Displacement (mm)

Cu-C/C

0

200

400

600

800

1000

1200

1400

1600

1800

0 0.5 1 1.5 2

Displacement (mm)

Load

P (

N)

Fig. 11. Typical loading–unloading curves of C/C, Si–C/C and Cu–C/Cspecimens.


shows the partial fractured location of a Si–C/C specimen.It is noted that the images before testing were taken atselected locations of each specimen. However, it was diffi-cult to predict the final fracture location of the specimenbefore testing. Hence, only very few images before testingcould catch the fracture location. In addition to the differ-ence of failure behaviors of three composites, the tensilestrength of Si–C/C is also very different from those of C/C and Cu–C/C. Averaged values and scatter ranges of ten-sile strength and tensile stiffness of three composites aregiven in Fig. 9. Cu–C/C composite shows the highest ten-

Fig. 12. Typical fractured specimens of C

sile strength and stiffness and Si–C/C shows the lowest ten-sile strength. Comparing with C/C, infiltration of copperinto C/C leads to 17% increase in tensile strength and15% increase in tensile stiffness, whereas infiltration of sili-con results in very low tensile strength and 7% increase intensile stiffness. The stiffness and strength of Cu–C/C arealmost in consistence with the prediction using mixed rule.However, the mixed rule is not available for Si–C/C. Refer-ring to the microstructure analysis it is considered that themajor reasons for low strength and slight improvement instiffness in the case of Si–C/C are the occurrence of new

/C, Si–C/C and Cu–C/C composites.

0

2

4

6

8

10

12

14

16

18

20

0 0.2 0.4 0.6 0.8

Normalized equivalent crack length a/W

Crit

ical

ene

rgy

rele

ase

rate

(kJ/

m2 )

C/C

Cu-C/CSi-C/C

1

Fig. 13. Critical energy release rate of C/C, Si–C/C and Cu–C/Ccomposites.


cracks, degradation of carbon fibers, and high interfacebonding between Si and carbon due to the formation ofSiC during infiltration process. Cracks and degradationof carbon fibers reduce the stiffness and strength of lami-nate. On the other hand, high ILSS restrains the occurrenceof delamination which plays an important role in the relax-ation of stress concentration in the 0� layer ahead the cracktip. Consequently the high stress level leads to the fractureof 0� layers of Si–C/C specimens at a low loading level.Thus restraining the crack occurrence, fiber degradationand the formulation of SiC interlayer during infiltrationprocess is necessary to obtain high strength Si–C/Ccomposite.

3.3. Compact tension testing results

Compliance calibration curves of three compositeobtained from monotonic tensile testing of CT specimenswith various pre-crack lengths are depicted in Fig. 10.Cu–C/C shows smaller compliance than other two com-posites. The difference of compliance between C/C andSi–C/C is not clear. These results are considered to be rea-sonable referring to the results of tensile stiffness of threecomposites. Typical load–displacement curves of threecomposites are given in Fig. 11. It is noted that each spec-imen for each composite has different pre-crack length asdescribed in the previous section. C/C and Cu–C/C speci-mens show similar crack growth behavior, whereas Si–C/C specimen shows a different behavior. Observation ofthe typical fractured specimens of three composites, asshown in Fig. 12, can help us to understand the difference.Large scale delamination occurred as the crack grows inthe case of C/C and it is difficult to see a clear crack. Rel-ative small scale delamination occurred in the case of Cu–C/C and crack growth path is relatively clear. In contrast,no delamination can be seen from the Si–C/C specimenand the crack growth path is clear. According to theseload–displacement curves and compliance calibrationcurves, critical energy release rate due to each crack incre-ment was determined using Eq. (1). Crack growth resis-tance curves (R-curves) for three composites are depictedin Fig. 13. In the early stage of crack growth in Si–C/Cand Cu–C/C, no or only little damages existed aroundthe crack tip when the main crack grown, thus the criticalenergy release rate shows relatively small values. After thisstage, fiber bridging in Si–C/C specimen caused theincrease of critical energy release rate. In contrast, fiberbridging, delamination, and plastic deformation of infil-trated copper significantly enhanced the energy release rateof Cu–C/C. Moreover, silicon is brittle but copper isductile, and the SiC formed from chemical reactionbetween silicon and matrix or carbon fiber causes a stronginterfacial bonding which restrains the interlaminardelamination and interfacial debonding. On the contrary,copper is inert to carbon and the infiltration temperatureis also lower than that for silicon, a weak interface bondingleads to a pseudo-plastic fracture behavior for Cu–C/C.

Consequently, Cu–C/C shows much higher crack growthresistance than Si–C/C. On the other hand, original lowcost C/C showed the highest fracture toughness amongthree composites. This high fracture toughness of C/C iscaused by interlaminar delaminations and interfacial deb-onding between fiber and matrix. Large scatter range inthe case of C/C is supposed to be caused mainly by thescatter of delamination size according to the images offractured specimens. Therefore, similar to the tensile test-ing results, restraining the crack occurrence, fiber degrada-tion and the formulation of SiC during the infiltrationprocess may be the key point to obtain tough Si–C/Ccomposite.

4. Conclusions

Low cost C/C and liquid metal infiltrated low cost C/Ccomposites Cu–C/C and Si–C/C are investigated. Analysisof microstructure, tensile testing, and compact tensiontesting are conducted, which leads to the followingconclusions:

(1) Infiltration of metals into C/C composite is an effec-tive method for the densification of porous C/Ccomposites.

(2) Infiltration of copper into the low cost C/C enhancesthe tensile strength by about 17% and tensile stiffnessby about 15%. The mode I fracture toughness of Cu–C/C composite is three times of that of Si–C/C.

(3) Infiltration of silicon into the low cost C/C causes therapid decrease of tensile strength and mode I fracturetoughness.

(4) Optimally controlling the microstructure of metalinfiltrated C/C composites is important to obtaindesired densification of composites. In particular,restraining the crack occurrence and controlling thereaction of silicon/carbon matrix and silicon/carbonfiber during the infiltrating process are necessary forSi–C/C to obtain high strength and high fracturetoughness.


Acknowledgements

Testing materials and partial financial support from theNGK (NGK Insulators, Ltd. Japan) are greatly acknowl-edged. Authors are grateful to Mr. K. IIhoshi and K. Ishi-matsu for their efforts in mechanical experiments.

References

[1] Taylor R. Carbon matrix composites. In: Kelly A, Zweben C, editorsin-chief, Warren R, volume editor. Comprehensive composite mate-rials, vol. 4. Amsterdam: Elsevier; 2000.

[2] Savage G. Carbon–carbon composites. Chapman and Hall; 1993.[3] Chang T, Nakagawa T, Okura A. Studies on a new manufacturing

process of carbon fiber reinforced carbon matrix (C/C) composites.Report of institute of industrial science, The University of Tokyo,vol. 35(8); 1991.

[4] Krenkel W, Gern F. Microstructure and characteristics of CMCmanufactured via the liquid phase route. In: Proceedings of ICCM-9,Madrid, Spain; 1993; vol. II. p. 173–81.

[5] Krenkel W, Fabig J. Tailoring of microstructure in C/C–SICcomposites. In: Proceedings of ICCM-10, Whistler, Canada, vol.IV; 1995. p. 601–9.

[6] Japanese Patent Number, 2000-081062; 2000 [in Japanese].[7] Hatta H, Suzuki K, Somiya S, Sawada Y. Strength improvement by

densification of carbon–carbon composites. Carbon 2001;39:83–90.[8] Denk L, Hatta H, Misawa A, Somiya S. Shear fracture of C/C

composites with variable stacking sequence. Carbon 2001;39:1505–13.

[9] Goto K, Hatta H, Katsu D, Machida T. Tensile fatigue of alaminated carbon–carbon composite at room temperature. Carbon2003;41:1249–55.

[10] Matsubara T, Takao Y, Wang WX, Hanzawa S. Effects of contactpressure and reciprocation number on the coefficients of friction ofcarbon/carbon composites with metal infiltration. In: Proceeding ofthe second Asian–Australasian conference on composite materials,Kyongju, Korea; 2002. p. 701–6.

[11] Moriyama M, Takao Y, Wang WX, Matsubara T. Strength of metalinfiltrated C/C under cyclic loading. Compos Sci Technol2006;66:3070–82.

[12] Kostopoulos V, Markopoulos YP, Pappas YZ, Peteves SD. Fractureenergy measurements of 2-D carbon/carbon composites. J Eur CeramSoc 1998;18:60–79.

[13] Hashemi S, Kinloch AJ, Williams JG. The analysis of interlaminarfracture in uniaxial fiber–polymer composites. Proc Roy Soc Lond A1990;427:173–99.

Tensile Strength and Fracture Toughness of C-C

Documents