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Materials Science and Engineering A 507 (2009) 6–12

Contents lists available at ScienceDirect

Materials Science and Engineering A

journa l homepage: www.e lsev ier .com/ locate /msea

nfluence of the strain rate on the mechanical behavior of the 3Deedle-punched C/SiC composite

ulong Li ∗, Tao Suo, Mingshuang Liuchool of Aeronautics, Northwestern Polytechnical University, P.O. Box 118, Xi’an 710072, People’s Republic of China

r t i c l e i n f o

rticle history:eceived 4 May 2008eceived in revised form5 November 2008

a b s t r a c t

Compressive mechanical behavior of a 3D needle-punched C/SiC composite at room temperature wasinvestigated at strain rate ranging from 10−4 to 6.5 × 103 1/s using the electronic universal testing machineand the split Hopkinson pressure bar. The experimental results show the strain rate has slight influenceon the mechanical property of the material. The dynamic compression strength obeys the Weibull dis-

ccepted 7 January 2009

eywords:D needle-punched C/SiC compositesechanical propertiesigh strain rate

tribution with the Weibull parameter m calculated as 8.19. Scanning electron microscope (SEM) imagesof the fracture surface of specimens tested at various strain rates indicate that the failure pattern of thematerial also varies with the strain rate.

© 2009 Elsevier B.V. All rights reserved.

ynamic compression strengthamage angle

. Introduction

Carbon fiber-reinforced SiC-matrix composites (C/SiC) haveuch advantages as low density, high ratios of stiffness/weightnd strength/weight, and failure stress that sustains under highemperature conditions, all of which have enabled their extensivepplication in many engineering structures, especially in aeronau-ic and aerospace fields [1]. The C/SiC composites developed cane grouped into 2D composites and 3D ones. The interface sheartrength of 2D-C/SiC composites was found by many researchers toe relatively weak [2–6]. To overcome this disadvantage, 3D C/SiComposites were developed [7]. Among them, 3D woven C/SiC com-osites and 3D needle-punched C/SiC composites were the mostidely investigated ones. While the experimental results of the 3Doven C/SiC composites exhibit remarkable anisotropy [8–10], the

esearch results of the 3D needle-punched C/SiC composite demon-trate that its interface shear strength as well as the mechanicalrosion resistance could benefit from the special microstructure ofhe material [11,12].

Considerable interest has been focused on 3D needle-punched/SiC composite in recent years. Microstructure and flexural prop-

rties of a 3D C/SiC composite was examined by Xiao et al. [13],rom which the composite was found to be damaged layer byayer with good pseudo-plasticity and that pyrolytic carbon coat-ng could moderate the interface strength between C fibers and SiC

∗ Corresponding author. Tel.: +86 29 8849 4859; fax: +86 29 8849 4859.E-mail address: liyulong@nwpu.edu.cn (Y. Li).

921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2009.01.024

matrix. Recently, Fan et al. [14] and Xu et al. [11] analyzed its fric-tion performance and suggested that the grain-abrasion was themain wear mechanism of the material. The mechanical behavior of3D needle-punched C/SiC composite under quasi-static loading atroom temperature was also investigated by Wan [12], who reportedthat the bending strength vertical to layers was weaker than thestrength parallel to layers, while the compression strength verticalto layers almost equaled that parallel to the layers. However, to thebest of our knowledge, the dynamic compressive behavior of 3Dneedle-punched C/SiC composite has not yet been reported.

In this paper, the uniaxial compressive behavior of a 3D needle-punched C/SiC composite was investigated under both quasi-staticand dynamic loadings at room temperature. The failure pattern wasobserved via a scanning electron microscope (SEM).

2. Experimental procedure

2.1. Material preparation

The 3D needle-punched C/SiC composite tested in this studywas offered by the State Key Laboratory of Solidification Process-ing in Northwestern Polytechnical University, People’s Republicof China. The carbon fiber utilized was T300 carbon fiber fromthe Nippon Toray Corporation. The quasi 3D reinforced preform

was firstly prepared by three-dimensional braided method. Thenisothermal/isobaric CVI (ICVI) was employed to deposit a thin pyrol-ysis carbon layer on the surface of carbon fiber as the interfaciallayer with butane at 850 ◦C prior to densification. Finally, methyl-trichlorosilane was used for deposition of the SiC matrix to get the

Y. Li et al. / Materials Science and Engineering A 507 (2009) 6–12 7

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It should be noticed that the catastrophic brittle failure was notobserved for the specimens tested at different strain rates. Instead,despite of the decrease of true stress with true strain after the stressreaches its compression strength, the material still possesses a rel-

Fig. 1. The diagram illustration of

D-C/SiC composite with density of 1.90 g/cm3 and open porosity6.09 ◦C.

.2. Experiment procedure

The specimens were cut from a 3D needle-punched C/SiC com-osite plate to ensure that the loading direction is parallel to thehickness direction of the material. The specimens tested underuasi-static and dynamic loadings at room temperature were bothylindrical, with a diameter of 5 mm and a length of 4 mm. Bothhe quasi-static and dynamic experiments were performed at theoom temperature. The strain rates were controlled in the rangerom 1.0 × 10−4 to 6.5 × 103 1/s.

An electronic universal testing machine with a maximum loadapacity of 10 kN was used to perform the quasi-static experiments.he split Hopkinson pressure bar (SHPB), situated in Northwesternolytechnical University, was used for the high strain rate experi-ents. A steel pad was employed to restrain the maximum strainithin certain limits so that the fracture surface of the tested spec-

mens could be preserved. The diagram of the SHPB and the steelad was shown in Fig. 1. According to the one-dimensional elas-ic stress wave theory, the strain, stress and strain rate of testedpecimen can be calculated as

�S = E(

A

AS

)εT

εS = −2C0

lS

∫ t

0

εR d�

ε̇S = −2C0

lSεR

(1)

here εR and εT are, respectively, the transmitted and reflectedtrain pulses which can be measured by the strain gages stuck onhe input and output bars; C0, E, and A denote the longitudinallastic wave velocity, Young’s modulus and cross-sectional area ofhe loading bars respectively; ls and As are the length and cross-ectional area of the specimen respectively.

. Experimental results and discussion

.1. Experimental results

Fig. 2 shows the true stress vs. strain curves of the 3D needle-unched C/SiC composite at different strain rates. It can be observedhat both the compression strength and failure strain of the 3Deedle-punched C/SiC composite increase with applied strain rate,hile elastic modulus of the material is almost independent on

lit Hopkinson press bar systems.

strain rate. The compression strength of the material increasesfrom about 350 MPa at the strain rate of 10−4 1/s to 430 MPa atthe strain rate of 6.5 × 103 1/s with a relative increment of about23%. Although remarkable dispersity exists, a close inspection ofthe experimental results reveals that the compression strength ofthe material varies linearly with the logarithm of the strain rate (asshown in Fig. 3). Such phenomenon was also observed for 2D-C/SiCcomposite [2,15].

It is believed that microcracks are contributed to the strain ratesensitivity of the material. As we known, such defects as micro-cracks caused by the mismatch of thermal expansion coefficientsbetween the carbon fibers and the SiC matrix are unavoidable dur-ing the preparation of the composites. Under compression loading,the nucleation and extension of the microcracks will cause thefailure of the composites. It should be pointed out that both thenucleation and extension are processes of time dependence. Underquasi-static loading, the microcracks have enough time to nucleateand extend. While under dynamic loading, the applied load pulseis usually several hundred microseconds, which is insufficient forthe nucleation and extension of microcracks. As the result, highercompression strength is needed at high strain rates.

Fig. 2. The quasi-static and dynamic true stress vs. strain curves of the 3D needling-punched C/SiC composites.

8 Y. Li et al. / Materials Science and Engineering A 507 (2009) 6–12

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ture surface of the 3D needle-punched C/SiC composite tested atthe strain rates of 10−4 and 10−2 1/s are shown in Fig. 7(a) and(b) respectively. It can be seen that shear failure and delamina-tion feature the fracture of the material, which was also observed

ig. 3. Compression strength vs. strain rate curves of the 3D needling-punched C/SiComposites.

tively high load-bearing capacity. The failure strain of the materialested at different strain rates even exceeds 8%. It is reasonable toelieve that the fracture of the SiC matrix may cause a drop in the

oading capacity of the material after the peak stress; yet pulloutf the needle-punched carbon fibers demands additional energyhich in a way results in a higher failure strain of the 3D needle-unched C/SiC composite. Thus, catastrophic brittle failure doesot occur in Fig. 2, in place of which an obvious strain softening

s observed. While for 2D-C/SiC composites tested by Liu [15], theaterials fracture quickly after the stress reaches its compression

trength. That is to say, the needle-punched fibers may make a pos-tive contribution toward the toughness of the 3D needle-punched/SiC composite. As the results, the 3D needle-punched C/SiC com-osite have improved fracture toughness compared with 2D-C/SiComposite.

.2. Weibull distribution of the dynamic compression strength

As the material was prepared using the CVI process, defectss gas-holes are inevitable and the heterogeneous distributionf them may lead to the dispersity of the material strength.hus, remarkable dispersity of the compression strength of the 3Deedle-punched C/SiC composites tested at the same strain ratender dynamic loading conditions can be observed in Fig. 4. Ithould be explained that as steel pads were employed to restrainhe maximum strain within certain limits during experiments onome samples, the failure strain of three samples (4#, 5# and 7#)s only 0.04, while five of other samples have failure strains morehan 0.16.

To quantify the degree of dispersity of the dynamic compressiontrength of the material, the Weibull distribution was applied inhis study. In the case of 3D needle-punched C/SiC composites, thehree basic constituents, namely, fiber, interphase, and SiC matrixre essentially brittle and the cracking involves defect-inducedandom failures. It thus follows that the statistical distribution ofhe strength of the 3D needle-punched C/SiC composite can beescribed by Weibull equation as[ (

�)m

]

(�) = 1 − exp −

�0(2)

here F(�) is a probability function, � the stress applied, �0 thecale parameter or characteristic strength, and m is the Weibullodulus, the most important parameter which characterizes the

Fig. 4. The true stress vs. strain curves of the 3D needling-punched C/SiC compositesat the same strain rates under dynamic compression.

dispersity of the material. Larger value of m indicates low den-sity of defects and low dispersity of the compression strength.The maximum density of the Weibull distribution can be achievedwhen � equals �0. It can be seen from Fig. 5 that the dynamiccompression strength obeys the Weibull distribution. The calcu-lated Weibull modulus m and the characteristic strength �0 ofthe 3D needle-punched C/SiC composite are 8.19 and 426.4 MParespectively.

Because of the remarkable dispersity of the compressionstrength of the 3D composites, at least three tests were conductedfor each experimental condition. The stress–strain curve with thefailure strength closing to the trend line in Fig. 4 was picked tobe presented in Fig. 2 as a typical stress–strain curve of the 3Dcomposites at corresponding experimental condition.

3.3. Observation of failure surface

The fractures of the failure specimens were observed by using anoptical microscope and JSM6460 SEM. Optical images of the frac-

Fig. 5. The Weibull distributions of dynamic failure strengths of the 3D needling-punched C/SiC composites.

Y. Li et al. / Materials Science and Engineering A 507 (2009) 6–12 9

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Fig. 6. Fractures of the 3D needling-punched C/S

n 3D C/C composite tested at compression load by Xiong et al.16]. An increase of shear fracture angle of the failed specimensrom 35◦ at the strain rate of 10−4 1/s to about 45◦ at 10−2 1/s waslso observed in Fig. 6(a) and (b). Garland et al. [17] and Bi et al.18] proposed that the strain rate effect of the interface strength

ay contribute to the magnitude of shear fracture angle in fibrousomposites. In other words, high compression strength of the 3Deedle-punched C/SiC composite at high strain rates results in a

arge shear fracture angle. Fig. 6(c) and (d) shows the fracture sur-aces of the material tested at the strain rates of 1.1 × 103 and× 103 1/s respectively. Differing from the failure pattern underuasi-static loading condition, the failure mode under dynamicoading displays a split pattern. The same failure mode was alsobserved in 3D needle-punched C/C composite loaded at high strainates [19]. It is also interesting to find one end of the specimens ofD composite tested at high strain rate broke into pieces. The exacteason for such phenomenon is not clear at present time. Shearracture angle of the 3D needle-punched C/SiC composite are notefined either.

The SEM images of the ruptured surface of the specimens testedt strain rates of 10−4, 10−2, 1.1 × 103 and 3 × 103 1/s are shown inig. 7 respectively. It can be observed from Fig. 7(a) and (b) that aarge number of fibers were pulled out during deformation at strainates of 10−4 and 10−2 1/s, the length of which decreases with the

ncrease of strain rate. In contrast, the specimens tested at strainates of 1.1 × 103 and 3 × 103 1/s display smoother fracture surfacessee in Fig. 7(c) and (d)). This indicates the number of fragmentalbers in the 3D needle-punched C/SiC composite under dynamic

oading increases with the strain rate.

posites under static and dynamic compression.

3.4. Comparison with 2D-C/SiC composites

In this section, in order to display the advantages of the 3Dneedle-punched C/SiC composite tested in this study, the mechan-ical behavior of the material is also compared with that of 2D-C/SiCcomposites tested by Liu et al. [15]. The true stress vs. strain curvesof both two kinds of 2D composites with different densificationand the 3D composite tested in this study are shown in Fig. 8. Itcan be seen that at the same strain rate, the compression strengthof 3D needle-punched C/SiC composite is higher than that of lowdensification (LD) 2D-C/SiC composite but lower than that of highdensification (HD) counterpart. It is also interesting to find the3D needle-punched C/SiC composite shows larger failure strainthan both kinds of 2D-C/SiC composite. As has been proposed in[15], the relatively low interface strength of LD 2D-C/SiC compos-ite resulted in a weak load-bearing capacity. Thus an increase inthe strength of 2D-C/SiC composite necessitates an increase of itsintensity. But it should be pointed out that the increase of inten-sity can also lead to a reduction in inhere low toughness of thematerial [15]. By contrast, the carbon fibers needle-punched intothe laminates in thickness direction of 3D composite can lead tohigher interface shear strength. At the same time, the insertedcarbon fibers can also play a toughening effect on the 3D needle-punched C/SiC composites. As the result, the 3D needle-punched

C/SiC composite shows not only high compression strength but rel-atively high toughness even after the stress reaches its compressionstrength. Hence it can be concluded that the 3D needle-punchedC/SiC composite possess advantages of both LD and HD 2D-C/SiCcomposites.

10 Y. Li et al. / Materials Science and Engineering A 507 (2009) 6–12

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Fig. 7. SEM micrographs of the 3D needling-punched C/SiC co

Table 1 lists shear fracture angles of both kinds of 2D compositend 3D needle-punched C/SiC composite. It can be seen that theD composite shows larger shear fracture angle under quasi-static

oading than that of LD 2D-C/SiC composite but smaller than that ofD 2D-C/SiC composite. As mentioned above, higher compression

trength will lead to larger shear fracture angle. Thus, larger shearracture angle also indicates excellent compression behavior of theD composite which is consistent with above conclusion derivedirectly from Fig. 8.

ig. 8. Comparison of the 3D needling-punched C/SiC composites with low densitynd high density 2D-C/SiC composites under static and dynamic compression.

ites ruptured surface under static and dynamic compression.

SEM images of the ruptured surface of both kinds of 2D-C/SiCcomposite loaded under quasi-static and dynamic loading condi-tions are shown in Fig. 9 respectively. For LD 2D-C/SiC composite,in Fig. 9, pull out fibers can be observed while bundles of fibers werecut off for HD 2D-C/SiC composite. The smoother fracture surface ofHD composites under both quasi-static and dynamic compressionindicates more small cracked fibers. Matrix fragments generatedunder dynamic compression for both LD and HD composites. Whilefor 3D needle-punched C/SiC composite, the fracture surfaces arerelative rough and more fibers were pulled out under both quasi-static and dynamic loading (see in Fig. 7). This phenomenon can beattributed to needle-punched carbon fibers in thickness direction.

As proposed in Section 3.2, the Weibull modulus of the 3Dneedle-punched C/SiC composite is 8.19 while those of the LD andHD 2D-C/SiC composites are 8.36 and 5.27 respectively. It is obviousthat the Weibull modulus of the 3D needle-punched C/SiC compos-

ite is comparable with that of LD 2D-C/SiC composite and higherthan that of HD 2D-C/SiC composite. As a result, low dispersity isexpected for the compression strength of the 3D needle-punchedC/SiC composite.

Table 1Shear fracture angles of 2D-C/SiC composites and 3D needle-punched C/SiCcomposite.

Material Shear fracture angles

Quasi-static loading Dynamic loading

LD-2D composite 30◦ (ε̇ = 10−4 l/s) 36◦ (ε̇ = 2300 l/s)HD-2D composite 50◦ (ε̇ = 10−4 l/s) 55◦ (ε̇ = 850 l/s)3D composite 30◦ (ε̇ = 10−4 l/s), 45◦ (ε̇ = 10−2 l/s) –

Y. Li et al. / Materials Science and Engineering A 507 (2009) 6–12 11

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Fig. 9. SEM micrographs of the 2D-C/SiC composites wit

. Conclusions

In this paper, the uniaxial compressive behavior of 3D needle-unched C/SiC composite at room temperature was investigatedt strain rate ranging from 10−4 to 6.5 × 103 1/s using the elec-ronic universal testing machine and the split Hopkinson pressurear technique. Pseudoplasticity due to the fracture of SiC matrix

s observed for specimen tested at all strain rates. Although theompression strength of the material increases with the strain rate,he strain rate sensitivity of the material is considered to be low.he compression strength of the 3D needle-punched C/SiC compos-te obeys the Weibull distribution. The calculated Weibull modulus

is 8.19, which is comparable with that of the LD 2D-C/SiC com-osites. The needle-punched fibers enhance the toughness of theD needle-punched C/SiC composite. The failure pattern of the 3Deedle-punched C/SiC composite also varies with strain rate. Shear

ailure and delamination play prominent roles in the fracture pro-ess under quasi-static loading whereas a split pattern featureshe failure mode under dynamic loading. Rough fracture surface ofpecimens tested is observed at the strain rates of 10−4 and 10−2 1/shile relatively smooth fracture surface is observed for those tested

t high strain rate.

cknowledgements

This research was supported by the National Natural Scienceoundation of China (No. 90405016), Doctoral Fund of Ministry of

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erent densification at quasi-static and dynamic loading.

Education of China (No. 20070699044) and the 111 project (B07050)to the Northwestern Polytechnical University. The authors alsothank Professor Laifei Cheng for providing the 3D needle-punchedC/SiC composite.

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