characterization of the microstructure of three-dimensional-needled carbon/silicon carbide...

Characterization of the Microstructure of Three-Dimensional-Needled Carbon/Silicon Carbide Composites

Fang Xu

National Laboratory of Solid State Microstructures, Department of Materials Science and Engineering,Nanjing University, Nanjing 210093, People’s Republic of China

National Key Laboratory of Thermostructure Composite Materials, Northwestern PolytechnicalUniversity, Xi’an 710072, People’s Republic of China

Yongdong Xu


Xiangkang Meng,* Yanfeng Chen, Haiming Lu, and Yi Zhang

National Laboratory of Solid State Microstructures, Department of Materials Science and Engineering,Nanjing University, Nanjing 210093, People’s Republic of China

Litong Zhang, Laifei Cheng, and Shangwu Fan


Int. J. Appl. Ceram. Technol., 7 [6] 821–829 (2010)DOI:10.1111/j.1744-7402.2009.02394.x

Ceramic Product Development and Commercialization

With great sadness we learned about the news that Professor Y. D. Xu, the main contributor to this article, had passed away suddenly. We will not forget Professor Xu’s open minded kindness

and widespread knowledge and wish to offer our sincere condolences to his family. We devote this article to his memory.

*[email protected]

r 2009 The American Ceramic Society

mailto:[email protected]

Three-dimensional-needled, carbon-fiber-reinforced silicon carbide matrix composites (C/SiC) were prepared by a chem-ical vapor infiltration and reactive melt infiltration method. It was found that two kinds of SiC existed in the C/SiC composites,that is, micro-b-SiC grains within the range of 5–15mm and nano-b-SiC grains with a size of about 100 nm. The interface ofC/SiC and the distribution of SiC showed evidence for the reaction mechanism of the reactive melt infiltration process.

Introduction

Carbon-fiber-reinforced silicon carbide matrix(C/SiC) composites are newly developed as brake ma-terials. Compared with conventional metallic materials,C/SiC composites have obvious advantages, such as ex-cellent thermal and mechanical properties in high tem-peratures and much lower weight with long life.1,2

Moreover, C/SiC composites also show other excellentproperties in comparison with C/C brake materials,such as excellent wear properties and lower sensitivityto surroundings and oxidation, that is, it is not so easyfor C/SiC to absorb moisture and oxidize. C/SiC com-posites also have a high and stable friction coefficient,especially the static friction coefficient.3–8 Therefore, C/SiC composites have become the new generation brakematerials for aircraft application.

In the early 1990s, DLR started the study of C/SiCbrake materials, which was followed by many institutesand companies.2,5–9 C/SiC brake materials with two-di-mensional (2D) fabrics, short-fiber-pressed preforms, andthree-dimensional (3D)-needled preforms have been in-vestigated extensively.2,3,5–9 Because of the poor thermalconductivities perpendicular to the friction surface for the2D fabric C/SiC composites, the surface temperature canexceed 10001C. The friction coefficient decreases and thewear rate increases with increasing temperature, both ofwhich are bad for brake performance. Short-fiber-pressedcomposites exhibit strong anisotropy in thermal conduc-tivities because the arrangement of fibers is anisotropic,which degrades the stability of brake performance. Thus,the applications of 2D fabric and short-fiber-pressed C/SiC composites are limited to relatively low performance.Recently, attention has been directed to 3D-needled C/SiC composites in order to improve the mechanical andthermal properties. Fan et al.9 has prepared 3D-needledC/SiC composites using a chemical vapor infiltration(CVI) and reactive melt infiltration (RMI) method,and the composite exhibits high interlaminar shearstrength, high and steady friction coefficient, and lowwear rate.

The present work was devoted to not only the in-vestigation of the microstructure of 3D-needled C/SiCcomposites and its influence on frictional and mechan-

ical properties, but also the investigation of the reactionmechanism.

Experimental Procedures

With CVI and RMI methods, C/SiC compositeswere prepared in three steps. First, the preform (Sup-plied by Nanjing Institute of Glass Fiber, Nanjing,China) of carbon fiber (T300, 3 K, 7 mm, each fiber,Nippon Toray, Tokyo, Japan) was prepared by a 3Dneedling method, starting with the repeated stacking of01 nonwoven cloth layers, short-cut web layers, and 901nonwoven cloth layers. Then, the needling method wasapplied to join different lamina by introducing carbonfiber bundles perpendicular to the lamina direction. Thetotal fiber volume fraction was 40%, the density of thepreform was 0.55 g/cm3, and the frequency of needlingwas about 35–40 needling/cm2. Second, the preformwas infiltrated with pyrolytic carbon (PyC) by CVI,which resulted in the formation of a C/C composite.The temperature and time for CVI were about 800–10001C and 300–700 h, respectively. Third, the RMIprocess, which involved the infiltration of C/C withmolten silicon, was performed. RMI was conducted at1420–17001C for 1–3 h.

Gravimetric analysis was used to determine thecontent of C, Si, and SiC in the composites. Si wasremoved by dissolving the composite in a mixture ofhydrofluoric and nitric acid (HNO3:HF 5 4:1) at 401Cfor 48 h. The content of C was measured by burning itoff at 7001C for 20 h in air, so the content of residualSiC could be calculated. At the same time, the SiC skel-eton could be used for further structural investigation.

The microstructure was observed by scanning elec-tron microscopy (SEM) (JSM 6700, Hitachi, Tokyo,Japan), mainly using secondary electron imaging at15 kV. Optical micrographs of polished specimenswere taken in an Olympus OLS3100 (Olympus, Tokyo,Japan), and the thermal conductivities were measuredusing a thermal properties analyzer Hot Disc (TPS1500, ThermTest, Houston, TX).

For transmission electron microscopy (TEM) andhigh-resolution transmission electron microscopy

822 International Journal of Applied Ceramic Technology—Xu, et al. Vol. 7, No. 6, 2010

(HRTEM) analysis, sections of C/SiC specimens werecut into a thickness of 1 mm using a low-speed diamondsaw. By cross-section machining, the holders of a3 mm disk were obtained. Alumina cross-section hold-ers were used to stabilize the disks during the prepara-tion process. The disks were polished on one side usingdiamond cloth. Then, the samples were turned and me-chanically ground to a thickness of more or less 20 mm.Thinning was performed on an ion polishing system(PIPS, model 691, Ar, Gatan, Pleasanton, CA ) at anion energy of 3.5 keV and an incident angle of 41. Theinvestigations were performed on a FEI Tecnai F20S-TWIN operating at 200 kV(FEI, Hillsboro, OR).This instrument has a point resolution of 0.24 nm.

Results and Discussion

Macrostructure of C/SiC Composite

The final composite consisted of 65 wt% C, 27 wt%b-SiC, and 8 wt% Si.9 Figure 1 shows macrographs ofthe final 3D C/SiC composite. The composite architec-ture had repeated stacking of 01 nonwoven web layers,short-cut fabric layers, and 901 nonwoven web layers.The layers were attached to each other by needle-

punched fiber bundles, oriented perpendicular to the fi-ber layers. After the RMI process, which includes theinfiltration of molten Si and the reaction of Si and C toform SiC, the large pores in porous C/C were filled withSiC and excess Si. However, some small pores betweenthe fibers within the bundles remained because they wereisolated or they were so small that they were closed bySiC formation before the additional Si could infiltrate.

Optical micrographs of the C/SiC composite areshown in Fig. 2, where the 3D-needled C/SiC compos-ites were asymmetrical materials with an obvious sub-segment structure (sub-C/C composites and SiCislands). In the region where fibers are sparse, meltedSi infiltrated into local C/C composites and SiC formed(the light gray region in Fig. 2). The SiC zone seemed todevelop as islands, which were separated from eachother by excess silicon (the white region in Fig. 2),rather than as a dense layer. Microcracks existed aroundthe SiC and Si areas. During cooling after RMI, Si ex-hibited a high-volume expansion (8%),6 which, togetherwith the thermal expansion mismatches among thephases in the composite, resulted in local stresses andthe occurrence of microcracks. Thus, it was necessary tocontrol the cooling rate in order to avoid damage to thecomposites. In the region of the fiber bundles, fairly

Fig. 1. Macrographs of a three-dimensional (3D) C/SiC composite: (a) a Short-cut fabric layer; (b) a nonwoven web layer; (c) a cross-sectionview; and (d) a three-dimensional view. 143 mm� 100 mm (150� 150 DPI).

www.ceramics.org/ACT 3D-Needled Carbon/Silicon Carbide Composites 823

dense sub-C/C composites formed. The optical picturerevealed that most of the carbon fibers, especially thoseinside nonwoven cloth, did not react with the moltensilicon during the infiltration process, locally retainingthe structure of C/C composites, as shown by the la-beled C/C in Fig. 2. Hence, the reinforcement effect offibers was maintained, ensuring the toughness of the fi-nal products.

In order to observe the distribution of SiC, SEMmicrographs of porous C/C composites and SiC skele-tons were taken and are shown in Fig. 3. The density ofthe C/C composite is 1.4–1.6 g/cm3 and the final den-sity of the C/SiC composite is 2.0–2.2 g/cm3; the openporosities of the C/C composite and C/SiC compositewere 16% and 5%, respectively. From Figs. 1a and 3a, itcould be seen that the short-cut fabric layer was ar-ranged randomly and there were many large pores,which afforded channels for the infiltration of meltedSi. Therefore, SiC formed uniformly inside the short-

cut fabric layers (Fig. 3b). As for the nonwoven web,Figs. 1b and 3 show that melted Si infiltrated into poresbetween fiber bundles or pores around needling fibers asa result of the dense arrangement of fibers. The strengthparallel to the friction surface can be guaranteed becausethe web arrangements alternate at different angles. Nee-dling fiber bundles perpendicular to the layers have fixedthe long fibers in the nonwoven web and made gaschannels perpendicular to the friction surface, throughwhich the reactant precursor entered the composite dur-ing CVI and the melted Si could infiltrate during RMI.Near the needling fibers, SiC was produced from PyCaround needling fibers perpendicular to the friction sur-face in the final product, all of which could improve theinterlaminar shear strength.10

As for standard 2D C/SiC and short-fiber-pressedC/SiC, the unstable friction coefficient mainly origi-nates from the higher surface temperatures, which aredue to the low transverse thermal conductivity.5 Theproduct of pn, proportional to the performance densityP/A, is responsible for the value of the friction coeffi-cient11

P=A ¼ mpn ð1Þ

where P, p, A, n, and m denote braking performance,pressure, friction area, velocity, and the friction coeffi-cient in units of W, MPa, m2, and m/s, respectively.

According to Eq. (1), higher values of n and p in-crease the performance density P/A and result in highersurface temperatures, as the brake process is a conver-sion of kinetic energy to heat energy. For a certain ma-terial system, P/A has a critical rating value; up to thecritical level of P/A, the coefficient of friction remainsstable. Exceeding the critical value, the C/SiC tribo-sys-tem cannot sustain higher braking performance and thematerial will overheat, locally forming so-called hot

Fig. 3. Scanning electron micrographs of a cross-section: (a) a porous C/C composite before RMI; (b) the SiC skeleton left after removing Cand residual Si from the C/SiC composite. 151 mm� 53 mm (150� 150 DPI).

Fig. 2. Optical micrograph of a C/SiC composite.109 mm� 79 mm (150� 150 DPI).


spots. As a result, the friction coefficient decreases whilethe surface temperature increases further.2 The improve-ment of thermal conductivity is effective for increasingthe critical value of the performance density and thenreducing the temperature on the friction surface.

In C/SiC composites, two methods can be used toachieve a higher transverse thermal conductivity. One isincreasing the ceramic content and the other is increas-ing the angle between the fibers and the friction sur-face.5 In our 3D-needled C/SiC composites, the contentof SiC increases because of the addition of short-cutfabric, and the angle between the fibers and the frictionsurface reaches 901C because the needling fibers areperpendicular to the friction surface. Thus, the obtained3D-needled C/SiC composites have higher transversethermal conductivity than those of short-fiber-rein-forced and 2D-fiber-reinforced C/SiC composites (aslisted in Table I).

Microstructure of C/SiC Composite

Further investigations were conducted to study themorphology of the SiC skeleton. As shown in Fig. 4,there are two different types of SiC, that is, micro scaledcoarse grains and nano scaled fine ones. Figure 4apresents coarse SiC grains with a size of 5–15 mm, whichcan be called micro-SiC. These grains show facettedsurfaces and have a high density of stacking faults, asindicated by the stripe pattern. Furthermore, it seemsthat some fine grains had grown around the coarse ones.Figure 4b shows both fine and coarse SiC grains, inwhich the fine SiC can also be divided into two groups:one is about 300 nm and the other is o100 nm. Both ofthem can be considered as nano-SiC.

Figure 5 shows that micro-SiC is located mainlyinside the short-cut fabric (Fig. 5a) and around the nee-dling fibers (Fig. 5b), where large pores exist before in-

filtration. However, nano-SiC was present between thefiber bundle/bundle in the nonwoven web (Fig. 5c),where smaller pores exist before infiltration. Specifically,micro-SiC usually formed where Si could infiltrate suffi-ciently.

According to tribology, every friction surface is notsmooth, but consists of a number of particulates. Ap-parently, the size and properties of these particulatessuch as hardness and elastic modulus are essential toform a continuous and stable friction film, which thenaffects the friction performance of the materials.11 Thebrake performance of the C/SiC composite can be rea-sonably explained by its composition.9,12 It is wellknown that the crystal structures of both the carbon fi-bers and the PyC matrix are hexagonal.11 The sub-C/C

Fig. 4. Micrographs of SiC grains in the C/SiC composite: (a) large SiC grains; and (b) small SiC grains. 148� 52 mm (150� 150 DPI).

Table I. Comparison of Transverse ThermalConductivities Among Short-Fiber-Reinforced C/SiCComposites, Two-Dimensional Fiber-Reinforced, and

Three-Dimensional-Needled C/SiC Composites

Thermal conductivity(W/mK)

C/SiC (short fiber)5

Schunk FU2952 14SGL carbon Sigrasic 40DaimlerChrysler C-brake 24Brembo CCM 20MS production Sicom 7DLR Silca SF 25–30

C/SiC (2D reinforcement)5

Schunk CF226/2 P77 20DLR Silca XS 15.3DLR Silca XG 18.9

C/SiC (3D needled) 71–83


composites easily form film-type debris after beingpeeled off. However, Si and SiC are brittle phases andhave high hardness, which will result in a high frictioncoefficient because of the furrow effect. Note that SiCcan be directly imbedded into the softer PyC matrixwhen the grains are coarse, leading to a serious furroweffect and very high wear rates, simultaneously. Toomany hard particulates are harmful to the formation of acontinuous friction film. However, the addition ofnano-SiC may solve this contradiction. Micro-SiC im-proves the friction coefficient through the furrow effect

while nano-SiC should easily slip when surrounded byan amorphous carbon film and thus help to form con-tinuous friction films, which make the friction coeffi-cient stable and decrease the wear rate.12

Reaction Mechanism of Microstructure Evolution

Figure 6 shows an EDX line scan over the two in-terfaces of a nano-SiC region. The line scan was per-formed from region 1 to region 2 (Fig. 6a), where onlythe signals of Si and C could be detected and the signal

Fig. 5. SiC grains in a C/SiC composite after etching of Si and burnout of carbon: (a) SiC within the short-cut fabric layers; (b) SiC aroundthe needling fibers; and (c) SiC between the fiber bundles. 146� 102 mm (150� 150 DPI).

Fig. 6. EDX line scan over the two interfaces of a nano-SiC region: (a) STEM image of the area where the line scan was conducted; and (b)EDX line scan over the interface of C/nano-SiC and the interface of nano-SiC/micro-SiC. 142� 68 mm (150� 150 DPI).


of silicon remained zero in the carbon region. Then, anano-SiC layer with a thickness of about 500 nm wasobserved. At the interface of C/nano-SiC, the carbonsignal decreased clearly with a slight slope over a widthof 100 nm, whereas the signal of silicon increased in thesame region. The slope of the signal was observed be-cause the electron beam and the interface were notaligned parallel to each other. The intensities of the Cand the Si signals remained unchanged when scanningover the two different types of SiC regions. Namely, thechemical compositions of these two areas appear to besimilar.

The C/SiC interface was shown more clearly in Fig.7a. As discussed above, there was also a nano-SiC layerwith a thickness of 100 nm between the micro-SiC andC. The micro-SiC had a number of stacking faults, asindicated by the stripe pattern (Fig. 7b). Combinedwith the results shown in Fig. 5, it can be concluded thatmicro-SiC grains exist at the SiC/Si interface while thenano-SiC layer forms at the SiC/C interface.

Moreover, an interesting phenomenon was ob-served in the SiC skeleton (Fig. 8). As the carbon fiberin region A and the PyC in regions C and E had beenremoved, nano-SiC in region B between the carbonfiber (region A) and the PyC (region C) was left asa hollow tube, which could not grow into coarse SiCbecause no further Si could infiltrate into the spacebetween the carbon fiber and the PyC. However, therewas nano-SiC near region C and coarse SiC might beobserved near region E. Because micro-SiC could befound between the nano-SiC layer and Si in the region

where Si was present, it could be concluded that thesemicro-SiC grains could not newly nucleate from thereaction with carbon, but grow on or nucleate fromthe nano-SiC grains by the dissolution/recrystallizationprocesses.

Investigations of the C/nano-SiC interface weremade using HRTEM. The micrograph of partly siliconi-zed PyC matrix is shown in the left part of Fig. 9 and thecrystal lattice of SiC can be observed in the right part. Nointerphase was detected at the C/SiC interface. However,on the left side of the interface, some regions that appearto be crystalline SiC can be detected (the area betweentwo white lines in Fig. 9), which indicates that SiC maygrow inside PyC in the following way: Si diffuses from

Fig. 7. Transmission electron (TEM) micrographs of a C/SiC composite: (a) the interface between nano-SiC/micro-SiC; and (b) TEMmicrographs of nano-SiC. 140� 70 mm (150� 150 DPI).

Fig. 8. SiC grains around a C fiber. 73� 51 mm(150� 150 DPI).


SiC to PyC on the left and reacts with PyC. Then newnano-SiC was generated near the C/SiC interface.

Figure 10 shows TEM and HRTEM micrographsof the SiC/Si interface. As shown in Fig. 10a, the SiC/Siinterface is obvious, and growth with different orienta-tions can be observed in Fig. 10b. No diffuse compo-sitional transition was observed between Si and SiC.Note that the HRTEM micrographs of SiC and Sicould not be taken in the same area because the TEMsample was uneven after Ar1 ion milling.

It has been reported that SiC can be formed by adissolution–deposition mechanism during the RMIprocess, although there was little direct evidence.13,14

Our present study supports this assertion. Figure 7 in-dicates the diffusion of Si to C. Namely, after the exo-thermic reaction between Si and C at the beginning ofinfiltration, Si has to diffuse through the previously de-veloped SiC nanocrystals to form new nano-SiC grainsat the C/SiC interface. Because of the nonequilibriumconditions of the reaction, the SiC grains contain a highdegree of stacking faults to minimize the energy of thegrown pattern (Fig. 7). As shown in Figs. 4 and 5, whensufficient Si enters an area, nano-SiC can grow into mi-cro-SiC regions by dissolution–deposition with the driv-ing force provided by the reduction of stacking faults(Figs. 5 and 6).

Summary

(1) In 3D-needled C/SiC composites prepared byCVI and RMI, b-SiC zones form inside the short-cutfabric, around the needling fibers, and among the fiberbundles in the nonwoven web. C/C segments remaininside the fiber bundles and residual Si is located amongSiC grains. Moreover, microcracks form near SiC and Sizones. (2) In C/SiC composites there are two differentkinds of SiC, both of them having a number of stackingfaults. One type was micro-b-SiC grains with a size of5–15 mm at the interface of SiC/Si. The other type wasnano-b-SiC that formed as a layer at the smooth inter-face between SiC/C. (3) The distribution of SiC and theC/SiC interface show evidence for the reaction mecha-nism of the RMI process.

Fig. 10. Interface of SiC/Si: (a) a transmission electron micrograph; and (b) a high-resolution transmission electron micrograph of Si–SiC,marked by the frame in (a). 134� 65 mm (150� 150 DPI).

Fig. 9. A high-resolution transmission electron micrograph of theC/SiC interface. 75� 60 mm (150� 150 DPI).


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characterization of the microstructure of three-dimensional-needled carbon/silicon carbide...

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