lable at ScienceDirect

Solid State Sciences 12 (2010) 617–623

Contents lists avai

Solid State Sciences

journal homepage: www.elsevier .com/locate/ssscie

Microstructure and tribological properties of 3D needle-punched C/C–SiCbrake composites

Peng Xiao, Zhuan Li*, Xiang XiongState Key Laboratory of Powder Metallurgy, Central South University, Changsha Hunan 410083, China

a r t i c l e i n f o

Article history:Received 30 September 2009Received in revised form29 December 2009Accepted 7 January 2010Available online 18 January 2010

Keywords:Ceramic matrixFibres and filamentsVapor depositionWearMicrostructure

* Corresponding author. Tel./fax: þ86 731 88830131E-mail address: li_zhuan@yahoo.com.cn (Z. Li).

1293-2558/$ – see front matter � 2010 Elsevier Massdoi:10.1016/j.solidstatesciences.2010.01.014

a b s t r a c t

Carbon fibre reinforced carbon and silicon carbide dual matrix composites (C/C–SiC) show excellenttribological properties and are promising candidates for advanced friction materials. A pressure infil-tration/carbonization combined with liquid silicon infiltration was developed for fabricating C/C–SiCcomposites. The carbon fabric preform was fabricated with the three-dimensional needling method. Inthe pressure infiltration process, the carbon fibre reinforced plastic was prepared by infiltration of thefabric preform with the furan resin. Then the carbon fibre reinforced plastic was carbonized which waspyrolysed to form a porous carbon/carbon composites. Finally, the porous carbon/carbon was infiltratedwith molten silicon to obtain C/C–SiC composites. The composites exhibit excellent friction behavior,including a good stability of brake, and the average dynamic m is 0.38 and static m is 0.50, in combinationwith the linear wear rate of about 5.6 mm cycle�1. Moreover, the friction surface was covered with frictionfilm which is about 10 mm in thickness. These results show that the C/C–SiC brake composites arepromising candidates for advanced brake and clutch systems.

� 2010 Elsevier Masson SAS. All rights reserved.

1. Introduction

Carbon fibre reinforced carbon and silicon carbide dual matrixcomposites, developed for the thermal protection of spacecraft atthe German Aerospace Center, Stuttgart, have received consider-able attention for their high potential in brake systems [1]. Asadvanced brake discs, C/C–SiC composites have been developedsince the end of 20 century following grey cast iron, powdermetallurgy and carbon/carbon (C/C). The grey cast iron is cheap inproduction and shows balanced friction behavior. However, thedensity of the iron is comparatively high (about 7.6 g/cm3), andthe mechanical and tribological performance is limited at hightemperature [2,3]. Powder metallurgy brakes have the advantagesof the maturity in material development and low cost, while theyalso have the disadvantages of poor performance and pronenessto corrosion at high temperature [4,5]. C/C composites aredeveloped to overcome the disadvantages of the grey cast ironand powder metallurgy, and are the current brake materials fordisks and pads of aircraft and racing cars. The advantages of C/Care low density (about 1/5 of iron density), high specific heatcapacity and small coefficient of thermal expansion and so on.

.

on SAS. All rights reserved.

However, the C/C brakes suffer from insufficient stability of fric-tion coefficient caused by humidity and temperature. Moreover,C/C composites are prone to oxidation beyond the temperature of400 �C and their cost is high, which prevents their use as brakecomposites in passenger cars and trains or emergency brakes oflifts and cranes [6–9].

Ceramic matrix composites, based on reinforcements of carbonfibre and matrices of silicon carbide (such as C/C–SiC), exhibitsuperior tribological properties in comparison with the abovebrake materials. In combination with their low density (about 2.0g/cm3), high thermal shock resistance, longer service life, espe-cially lower sensibility to surroundings and temperature fora silicon carbide share of at least 20% in mass, the C/C–SiC arepromising candidates for advanced friction materials in future[10,11]. Several institutes and industries now exist to investigate C/C–SiC composites for their use as friction materials for brake padsand disks [7,12–14]. For example, researchers of Stuttgart Univer-sity and German Aerospace Center have applied C/C–SiCcomposites to friction domain since the Middle of 1990s, and havedeveloped C/C–SiC brake lining applied in 911 Turbo of Porscheand Audi A8 automobiles.

At present, the main preparation methods of C/C–SiC compositesinclude: (1) a gas phase route, also referred to as chemical vaporinfiltration (CVI), (2) a ceramic route, i.e. a technique combining theimpregnation of the reinforcement with a slurry and a sintering step

Fig. 1. Schematic representation of manufacturing process for C/C–SiC composites.

P. Xiao et al. / Solid State Sciences 12 (2010) 617–623618

at high temperature and high pressure, also referred to high pres-sure-sinter process (HP-Sinter), as well as (3) a liquid phase routeincluding the polymer impregnation/pyrolysis (PIP), and liquidsilicon infiltration (LSI) also called (reactive) melt infiltration (RMI orMI) processes. Each of the former routes displays advantages anddrawbacks. CVI has been developed as an extension of the well-established chemical vapor deposition (CVD) and has already been incommercial production. The primary advantage of this method isuniform coating of tailored compositions for multi-layer at relativelylow temperatures (900–400 �C). Moreover, the CVI technique, similarto the CVD, is expensive, corrosive, toxic, explosive and has lowdeposition rates [15,16]. The HP-Sinter has a short processing timeand is inexpensive in developing at practice, and has been used mosteffectively for glass and glass-ceramics matrix systems. The problemof this process is degradation of fibre and damage resulting frommechanical contact and limitation to produce only one- or two-dimensional fibre reinforced materials. The PIP with great composi-tion homogeneity is easy for infiltration of forming. The main

Fig. 2. Thermogravimetric analysis of applied furan resin.

drawbacks of this technique are high shrinkage and low yield ofpolymer during pyrolysis [17]. Liquid silicon infiltration was firstlydeveloped at General Electric Company and German AerospaceCenter. It is based on a polymer pyrolysis and subsequently followedby a process that the liquid silicon reacts with a small amount of thecarbon and forms SiC as the ceramic matrix. The method showssignificant advantages: the process is conducted only once and noadditional densification step is necessary, the lowest manufacturingcost of all available C/C–SiC manufacturing processes and so on.However, the obstacles most likely to be encountered in applying this

Fig. 3. Schematic diagram of liquid silicon infiltration experiment.

Fig. 4. Schematic diagram of dynamometer. (1) inertial wheel, (2) bearing, (3) clutch,(4) rotor holder, (5) rotor, (6) stator (C/C–SiC), (7) pressing cylinder, (8) strap, (9) motor,and (10) lathebed.

Fig. 6. XRD pattern of C/C–SiC composite.

P. Xiao et al. / Solid State Sciences 12 (2010) 617–623 619

method to ceramic matrix are: fibre damaging and chemical reac-tions at infiltration temperature [18,19].

The aim of current work is to develop an improved techniqueto reduce the processing costs of the C/C–SiC brake composites.Pressure infiltration/carbonization combined with liquid siliconinfiltration is proposed for fabricating C/C–SiC composites. Thestructure characteristics, tribological performance and wearmechanisms of the composites were investigated.

2. Experimental procedure

2.1. Preparation of fabrication

PAN-based carbon fibre (T700) employed, each yarn contained12,000 filaments, was supplied by the Toray, Japan. Furan resin wasemployed for infiltration (FA-2, Duwei Chemical Co., LTD, Chang-shu, China). The diameters and purity of the silicon powder used inthis study were 30–50 mm and 99.0%, respectively, according toinformation supplied from the manufacturer of Da Zelin-silicon Co.,LTD, Beijing, China.

Fig. 5. Schematic diagram of testing specimen.

2.2. Fabrication of the composites

The processing of investigating C/C–SiC composites consisted offour steps as shown in Fig. 1. The first step involved preparation offabric preform by the three-dimensional needling technique, start-ing with repeatedly overlapping the layers of 0� non-woven fibrecloth, short-cut fibre web, and 90� non-woven fibre cloth withneedle-punching step by step. The density of the preform was about0.60 g cm�3. The second step involved production of carbon fibrereinforced plastic (CFRP) green preform by infiltration of the fabricpreform with the furan resin at 65 �C and a pressure of 1.6 MPafollowed by subsequent curing at 200 �C under pressure of 1.6 MPa.

Fig. 7. Phase compositions of C/C–SiC composite.

Fig. 8. DSC-TG curves of C/C–SiC composite.

Fig. 9. Microstructure of C/C–SiC composite for (a) Non-woven cloth layer, (b) Short-cut web layer, (c) Cross-section of the preform, (d) different component (e) SiC particles.

P. Xiao et al. / Solid State Sciences 12 (2010) 617–623620

The CFRP green preform was then pyrolysed in the temperaturerange between 200 �C and 850 �C under nitrogen atmosphere andthe CFRP was transformed to porous carbon/carbon composites.This step leads to the conversion of the furan resin to resin carbonas parts of carbon matrices. Above 650 �C only very small losses inmass happened as the temperature rose (as shown in Fig. 2).Consequently the resin was slowly heated between 200 �C and 650�C during pyrolysis to allow volatile decomposition products todiffuse through the matrix without disrupting the integrity of the

Fig. 10. Macroscopic friction surfaces for (a) stato

composites. The carbon yield of the precursor was about 65 wt%and accompanied by a high shrinkage of the resin of about 20% inlength. The density of the porous carbon/carbon composites wasabout 1.45 g/cm3.

The last step was the liquid silicon infiltration (infiltration of themolten silicon into the obtained porous carbon/carbon), which wasused to prepare the C/C–SiC composites (as shown in Fig. 3) byheating the porous C/C–SiC composites in vacuum at the range of1420–1700 �C for 1 h. When the melting point of silicon (1414 �C)

r disk (C/C–SiC composite) and (b) rotor disk.

Fig. 11. Relationship of friction coefficients and braking number of C/C–SiC composite.

P. Xiao et al. / Solid State Sciences 12 (2010) 617–623 621

passed, silicon rapidly infiltrated the carbon/carbon preform andreacted with both fibre and matrix to form silicon carbide. Thedensity and open porosity of the C/C–SiC composites were about2.37 g/cm3 and 2%, respectively.

2.3. Testing methods

The density and the open porosity of the samples weremeasured with Archimedes’ method. The tribological properties ofC/C–SiC brake materials were tested on an MM-1000 disk-on-disktype laboratory scale dynamometer (Fig. 4). The kinetic energyabsorbed by brake was supplied by the inertia wheels, which weredriven by a DC motor. The tested specimens acted as stator (Fig. 5),and rotor was 30CrMoSiVA alloy. The end surfaces of tested spec-imens were smoothed with the grinder, and the true friction areabetween the stator and rotor should running-in reach up to 80%before test. When the inertial wheel, which rotated with the rotorsimultaneously, was accelerated to a certain rotational velocity, thebrake was achieved through the friction between the rotor andstator. The given brake speed, brake pressure, and inertia were 25m/s, 1.0 MPa, and 0.1 kg m2, respectively. Rotating velocity, brakemoment, and brake time were recorded by computer. The frictioncoefficient can be calculated from the equation

M ¼ mðr1 þ r2Þp=2 (1)

where M is moment (N m), m friction coefficient, P brake pressure(N), r1 inner radius of friction surface (m), and r2 outer radius offriction surface (m).

The above test was repeated ten times and the average values ofthe measurements were expressed in the results.

Fig. 12. Morphology and wireframe of fr

2.4. Microstructure observation and analysis

The phase identification was performed by means of X-raydiffraction (XRD) (model Rigaku D/max 2550PC, Japan) using CuKa1 radiation with Si as an internal standard. The differentialscanning calorimetry-thermogravimetry (DSC-TG) curves of the C/C–SiC composites were determined using a detector (ModelST449C, Netzsch). Heating rate at 10 �C/min, and O2 gas at 50 mL/min were used in the DSC-TG experiment. The phase compositions,morphology of worn surfaces and wear debris were measured withscanning electron microscopy with energy dispersive X-ray spec-troscopy (SEM/EDX) (model JSM-6360LV, Japan) and three-dimensional digital video microscope (model KH-7700, Questar,America).

3. Results and discussion

3.1. Phases and DSC-TG analysises

The XRD pattern of C/C–SiC composites is shown in Fig. 6. Thephase analysis reveals the presence of residual silicon, carbon andsilicon carbide of C/C–SiC composites. The gravimetric analysis wasemployed to determine the content of carbon, residual silicon, andSiC in the composites. By etching the C/C–SiC specimen withaqueous solution of 90 vol.% nitric acid and 10 vol.% hydrofluoric atroom temperature for 48 h, it was possible to remove the residualsilicon, whereas the content of carbon fibre and resin carbon matrixwas measured by burning it off at 700 �C for 10 h in air. So thecontent of each component can be calculated. The gravimetricanalysis results are shown in Fig. 7, which reveals that the residualSi is few but SiC produced is much because of the high activity ofthe amorphous carbon.

Fig. 8 shows the results of DSC-TG analysis of the C/C–SiCcomposites, in which the composite has an obvious endothermicregion at 600–1300 �C. The former part is mainly the oxidation ofcarbon fibre and resin carbon, and the latter part is correspondingto the oxidation of residual Si. The TG curve decreases to about 52%sharply at 600–1300 �C. The melting point of Si crystal is 1414 �C,whereas the starting oxidation temperature of the residual Si in C/C–SiC composites under the melting point. This is because theoxidation of the near carbon lead to the local temperature higherthe Si melting point.

3.2. Microstructure characterization

The microstructures of C/C–SiC composites are shown in Fig. 9.C/C–SiC are composed of non-woven fibre cloth (Fig. 9(a)), short-cut fibre web (Fig. 9(b)) and needle fibre (Fig. 9(c)). In this kind ofstructure, non-woven cloth parallel to the friction surface benefitedto the formation of friction film to improve the tribological

iction film on C/C–SiC worn surface.

Fig. 13. SEM micrograph of C/C–SiC composite friction surface.

P. Xiao et al. / Solid State Sciences 12 (2010) 617–623622

properties of the materials. Meanwhile, the needle fibres not onlyenhance the inter-laminar shear strength, but also increase thethermal conductivity and hence the coefficient of friction stabilityof the materials.

For the porous carbon/carbon materials with high densitymicrostructure, molten silicon can only infiltrate the macroscopicholes of the carbon/carbon preform because of capillary and reactwith the meeting carbon (including resin carbon and partial carbonfibre) forming SiC during the liquid silicon infiltration. So the SiCusually is formed among the fibre bundles in the non-woven web,inside the short-cut fabric and around the needing fibre. Althoughthe fibre bundles show also holes, silicon can only partly infiltratethem because some are isolated; and the other holes are so smallthat they are closed by the formed SiC before silicon enters further.This results in a microstructure of carbon/carbon bundles sur-rounded by SiC (Fig. 9(d)).

The growth of SiC leads to the formation of a continuous poly-crystalline SiC layer, and further growth of this layer is attributed tothe diffusion of the reactive species through the silicon carbidecoating. Whereas the climbing velocity of liquid silicon on carbon ismore than the diffusion speed of the reactive species through theSiC coating for the moist angle between the liquid silicon andcarbon is 0�–37� [20]. Fig. 9(e) shows the SiC morphology of thebrake materials after removing the residual silicon. These SiC grainswith sizes between 5 and 20 mm, show facetted surfaces and aresometimes very big, and some of the grains have grown aroundothers.

Fig. 14. SEM micrographs of C/C–SiC composite worn debris for (a) worn grain cov

3.3. Frictional characteristics

3.3.1. Brake behaviors of the C/C–SiC compositeBrake is a mutual affection between two relative motional worn

surfaces, so the tribological characteristics are usually determinedby composites microstructure and worn surfaces under certainbrake condition. After brake tests for 10 cycles, the average dynamicm and static m were 0.38 and 0.50, respectively. The linear wear rateof C/C–SiC composite and corresponding alloy disk were about 5.6mm cycle�1 and 3.7 mm cycle�1. The disk samples showed perfectfriction surfaces, and the typical friction surfaces of alloy rotor diskand stator disk (C/C–SiC composite) are shown in Fig. 10. The wearsurface of the alloy rotor disk (Fig. 10(a)) is relatively smooth anddisplayed ‘‘green spot’’ in the direction of sliding. The real contactarea (effective friction area) of the C/C–SiC composite (Fig. 10(b))approaches the nominal area, and the shining regions are frictionfilms.

Considering the cyclic using and the friction stability of C/C–SiCcomposites disks, the influence of brake number on the coefficientof friction (COF) is analyzed. Fig. 11 shows that the COF valuefluctuates with a small range of about 0.04 versus the brakenumber, which suggests a good stability of brake against fadingversus several brake stops. The COF is determined by severalfactors, including the mechanical deformation of the contactjunction or wear debris, the groove effect of wear debris, theadhesion and abrasion between the contact surfaces [21]. Thefluctuation of COF is related with the roughness, hardness,temperature, sliding velocity and wear mechanisms on frictionsurfaces.

3.3.2. Morphologies of friction surface and wear debrisThe surface microstructure and wear debris on wear-induced

surfaces of C/C–SiC disk was investigated by a three-dimensionaldigital video microscope and the scanning electron microscopeanalysis. After the brake tests, the surface of C/C–SiC brake disk iscovered by friction film with micro-cracks and ‘‘green spot’’(Fig. 12(a)). The three-dimension wireframe of the friction film isshown in Fig. 12(b). It is obvious that there are two stacking frictionfilms, and the thickness of the above is about 10 mm.

Tribological contact is mainly determined by solid–solid inter-actions between wear debris on the mated surfaces of C/C–SiCsample and alloy disk, depending on the initial roughness and theapplied load, resulting from relative movement of large scale ofadhesion conjunction and forming stacking friction film as shownin A region of Fig. 13. Along with the brake process, a great deal ofbrake energy is translated into friction heat and leads thetemperature of the friction surface to rising continually. The highestinstantaneous temperature of the secondary friction surface is 459

ered on the friction surface and (b) thin flakes of material in the worn debris.

P. Xiao et al. / Solid State Sciences 12 (2010) 617–623 623

�C detected by thermocouple. However, various components of C/C–SiC with different coefficient of thermal expansion, thus it causesthe heat expansion rate between friction surface and secondaryfriction surface different and leads to thermal stresses, as well asthe different zones in the same friction surface. The thermal andmechanical stresses (constant load) is continuous between thecontact wear debris, which shows localized fatigue phenomenarelates to the displacement of wear areas and micro-cracks infriction film (as shown in Fig. 13B region). This friction filmproduces peeling and flakes (Fig. 13C and D region) during thebrake repeatedly with the relative hard matching alloy disk.

In order to understand the processes of brake materials wear, itis necessary to study the brake wear debris besides friction film.There are wear debris, including grains (Fig. 14(a)) and flakematerials (Fig. 14(b)), on the friction surface. XRD shows the debrisare not only composed of carbon, residual silicon, and SiC, but alsocomposed of SiO2, Fe3C and FeC. The results show that the oxidationand material transfer happened on the friction surface duringbrake. The temperature of local ‘‘hot spot’’ of the friction surfacecould overtop 1000 �C, while the antioxidation of carbon is poorover 400 �C as well as silicon carbide over 800 �C, which suggestedthe C/C–SiC composites would be oxidized unceasingly for thefriction heat and form ‘‘green spot’’ in friction film (as shown inFig. 12(a)).

As has been discussed above, it is obvious that the main wearmechanism of the C/C–SiC brake composites includes adhesion andoxidation abrasion.

4. Conclusions

Pressure infiltration/carbonization combined with liquid siliconinfiltration processes was developed for fabricating carbon fibrereinforced carbon and silicon carbide dual matrix composites inorder to reduce processing costs and improve their applicationarea. The microstructure and the tribological behavior of thecomposites were investigated, which exhibit excellent frictionbehavior, including a good stability of brake, the average dynamic m

is 0.38 and static m is 0.50, in combination with the linear wear rateof about 5.6 mm cycle�1. The wear is significantly determined bycyclic mechanical and thermal stresses, which results in the mainwear mechanism of the C/C–SiC brake composites adhesion and

oxidation abrasion. Moreover, the friction surfaces are covered withfriction film which is about 10 mm in thickness.

Acknowledgements

The authors gratefully acknowledge the financial support forthis study provided by the National Hi-Tech Research DevelopmentProgram of China (2006AA03Z560) and the Graduate Degree ThesisInnovation Foundation of Central South University (Contract No.2008yb019).

References

[1] W. Krenkel, in: Proceedings ISATA conference on supercars, Aachen, Germany,1994, pp. 316–322.

[2] G. Cueva, A. Sinatora, W.L. Guesser, A.P. Tschiptsch, Wear 255 (11) (2003)1256–1260.

[3] M.H. Cho, S.J. Kim, R.H. Basch, J.W. Fash, H. Jang, Tribology International 36 (7)(2003) 537–545.

[4] B. Breuer, U. Dausend, USA, SAE International, 2003, pp. 37–47.[5] P.P. Yao, H.C. Sheng, X. Xiong, B.Y. Huang, Transactions of Nonferrous Metals

Society of China 17 (1) (2007) 99–103.[6] C. Blanco, J. Bermejo, H. Marsh, R. Menendez, Chemical and physical properties

of carbon as related to brake performance. Wear 213 (1–2) (1997) 1–12.[7] W. Torsten, B. Gordon, Materials and Design 18 (1) (1997) 11–15.[8] J.D. Chen, J.H. Chern Lin, C.P. Ju, Wear 193 (1) (1996) 38–47.[9] B.K. Yen, Wear 192 (2) (1996) 208–215.

[10] W. Krenkel, B. Heidenreich, R. Renz, Advanced Engineering Materials 4 (8)(2002) 427–436.

[11] S. Vaidyaraman, M. Purdy, T. Walker, S. Horst. In: Fourth internationalconference on high temperature ceramic matrix composites (HT-CMC4)proceedings, Munich, Germany, 2001, pp. 802–808.

[12] R. Gadow, A. Kienzle, in: Proc 6th int symp on ceramic mat and componentsfor engines, Arita, Japan, 1997, pp. 412–418.

[13] Z. Stadler, K. Krnel, T. Kosmac, Journal of the European Ceramic Society 27 (2–3) (2007) 1411–1417.

[14] Z. Li, P. Xiao, X. Xiong, Chinese Journal of Nonferrous Metals 18 (5) (2008)1071–1075.

[15] R. Naslain, CVI composites. in: R. Warren (Ed.), Ceramic Matrix Composites,vol. 8. Chapman and Hall, , London, 1992, pp. 199–243.

[16] Y.D. Xu, L.F. Cheng, L.T. Zhang, Carbon 37 (8) (1999) 1179–1187.[17] R. Gadow, Ceramic Engineering and Society Proceeding 21 (3) (2000) 15–29.[18] W. Krenkel, Ceramic Engineering and Science Proceeding 22 (3) (2001) 443–

454.[19] W.B. Hillig, American Ceramic Society Bulletin 73 (4) (1994) 56–62.[20] F.H. Gern, R. Kochendorfer, Composites Part A 28A (1997) 355–364.[21] S.Z. Wen, Tribology Theory. Tsinghua University Book Concern, Beijing, 1991,

pp. 398–470.