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C-SiC composites containing Ti3SiC2
Wear 274–275 (2012) 188–195
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Friction and wear behaviors of C/C-SiC composites containing Ti3 SiC2 Xiaomeng Fan, Xiaowei Yin ∗ , Shanshan He, Litong Zhang, Laifei Cheng Northwestern Polytechnical University, National Key Laboratory of Thermostructure Composite Materials, Xi’an, Shaanxi 710072, PR China
a r t i c l e
i n f o
Article history: Received 1 August 2010 Received in revised form 22 August 2011 Accepted 23 August 2011 Available online 30 August 2011 Keywords: Ceramic–matrix composite Ti3 SiC2 MAX-phase Brake materials Wear testing
a b s t r a c t In the present paper, friction and wear behaviors of a carbon fiber reinforced carbon–silicon carbide–titanium silicon carbide (C-SiC–Ti3 SiC2 ) hybrid matrix composites fabricated by slurry infiltration and liquid silicon infiltration were studied for potential application as brake materials. The properties were compared with those of C/C-SiC composites. The composites containing Ti3 SiC2 had not only higher friction stability coefficient but also much higher wear resistance than C/C-SiC composites. At an initial braking speed of 28 m/s under 0.8 MPa pressure, the weight wear rate of the composites containing 5 vol% Ti3 SiC2 was 5.55 mg/cycle, which was only one-third of C/C-SiC composites. Self-lubricious filmlike debris was formed on the composites containing Ti3 SiC2 , leading to the improvement of friction and wear properties. The effect of braking speed and braking pressure on the tribological properties of modified composites were investigated. The average friction coefficient was significantly affected by braking speed and braking pressure, but the wear rate was less affected by braking pressure. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Carbon fiber reinforced carbon/silicon carbide binary matrix (C/C-SiC) composites fabricated by liquid silicon infiltration process (LSI) possess excellent friction properties and low sensitivity to service environments, making them promising brake materials, which is also named as C/SiC composites [1–3]. The LSI C/C-SiC composites comprise carbon fibers, C-SiC binary matrix and more than 8 vol% residual silicon mainly locating in the interbundle pores [2]. As a residual phase in brake materials, silicon in the matrix increases the wear and leads to the instability of friction behaviors [3,4]. This is supposed to be avoided by replacing silicon with alloys to give stable silicides or carbides [3]. The effect of graphitization on tribological properties of C/CSiC composites was studied, and it was found the stability of COF increased significantly [5]. Graphite filler was added into the C/CSiC brake materials, it was found that C/C-SiC composites with graphite filler was relatively insensitive to changes in braking speeds and displayed higher braking performance at high braking speeds due to the formation of the film-type wear debris [6]. Owing to its special nanolaminate structure and excellent properties, Ti3 SiC2 is extensively researched in recent years [7]. Ti3 SiC2 ceramics have been proved to be promising tribological materials [8–10]. The dependence of friction and wear properties of Ti3 SiC2 material on its grain size was investigated [8]. The wear resistance increases with increasing grain size for sliding and abrasive
∗ Corresponding author. Tel.: +86 29 88494947; fax: +86 29 88494620. E-mail address: [email protected] (X. Yin). 0043-1648/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.wear.2011.08.029
wear tests. Under dry conditions at room temperature, the friction coefficient of Ti3 SiC2 /Ti3 SiC2 pair was 1.16–1.43, but that of Ti3 SiC2 plate/diamond pin pair was 0.05–0.1, which was even lower than that of graphite plate/diamond pin pair [9]. The difference may depend on whether lubricious oxide tribofilms, TiCx Oy , can be formed or not [10]. For the Ti3 SiC2 ceramics, the oxide is important for its lubricating effect in the tribological tests. In our previous work [11], in order to eliminate residual free silicon in the LSI C/C-SiC composites, forming Ti3 SiC2 , a porous C/C preform fabricated by CVI was infiltrated with TiC particles, and then the preform containing TiC was infiltrated with liquid silicon, resulting in the formation of a carbon fiber reinforced CSiC–Ti3 SiC2 matrix composite. The effects of fabrication process on phase composition and microstructure of the composites were studied, and mechanical properties of the composites containing Ti3 SiC2 were compared with those of C/C and C/C-SiC composites, which indicated the introduction of Ti3 SiC2 was beneficial to improve the mechanical properties of C/C-SiC composites. The aim of the present study was to understand friction and wear behaviors of C/C-SiC composites containing Ti3 SiC2 , which was compared with those of C/C-SiC composites. 2. Experimental procedures 2.1. Sample preparation The fabrication procedure of C/C-SiC–Ti3 SiC2 composites included three steps. In the first step, porous C/C composite was prepared by chemical vapor infiltration (CVI) using 3D needled fiber felts as preform, and propylene was used as the precursor of
X. Fan et al. / Wear 274–275 (2012) 188–195 Table 1 Braking conditions for friction and wear performance tests. Test conditions
I
Inertia (kg m2 ) Braking pressure (MPa) Braking speeda (m/s) Braking timesb
0.235 0.3, 0.5, 0.8 5 10 30 30
a b
II
III
IV
V
VI
20 15
25 15
28 15
Braking speed: the linear velocity of outer edge of rotate disk. Braking times: the number of repeated test under the applied conditions.
carbon and Ar was used as carrying and diluting gas. In the second step, the CVI C/C composite was infiltrated by TiC slurry using the process parameter depicted in the previous work [11], which led to the formation of C/C-TiC composite. At last, the as-received C/C-TiC composites and C/C composites were infiltrated by liquid silicon at 1500 ◦ C in a vacuum furnace, and the C/C-SiC–Ti3 SiC2 composites and C/C-SiC composites were obtained. For ease of the comparison, the C/C-SiC composites and the C/C-SiC–Ti3 SiC2 composites were designated as sample A and sample B. 2.2. Characterizations of materials 2.2.1. Friction and wear properties Both samples A and B were machined into 76.5 mm in outer diameter, 52.6 mm in inner diameter and 12 mm in thickness as rotating disk, and into 90 mm in outer diameter, 55 mm in inner diameter and 12 mm in thickness as stationary disk, respectively. Friction and wear properties of the samples were performed on a disk-on-disk MM-1000 testing machine (Xi’an ShunTong Technical Research Institute, China), which was described previously in detail [2]. The braking conditions for friction and wear performance tests are shown in Table 1. The friction stability coefficient, S, during each braking test can be calculated from Eq. (1). S=
cp max
Table 2 Density and open porosity of samples A and B before and after infiltration with liquid silicon. Samples
15 30
(1)
where max is the maximum COF during each braking test, and cp is the average COF, which is automatically recorded by the machine. The temperature of the samples during braking was on-line measured using NiCr–NiSi thermal–electrical couples, which were inserted into a hole of 2 mm in diameter and 7 mm in depth of the stationary disk, and the distance between the hole and the friction surface was 2 mm. In every braking cycle, the braking speed of the brake disk fell from a certain speed to static, and the energy was transformed from kinetic energy into heat energy. The kinetic energy was determined by the braking speed, so that the kinetic energy increased with increasing braking speed under the same braking pressure. The friction resistance was determined by braking pressure, so that the friction force increased with the increase of braking pressure at the same braking speed. 2.2.2. Open porosity, density, phase composition, and microstructure The open porosity and the bulk density of the samples were measured by Archimedes method according to ASTM C-20 standard. The samples were crushed into power and analyzed by X-ray diffraction (XRD), via a computer-controlled diffractometer (X’Pert Pro, Philips, Netherlands) with Cu K␣ radiation at 40 kV and 100 mA. Data was digitally recorded in a continuous scanning mode in the angle (2) ranging from 8 to 80◦ with a scanning rate of 0.12◦ /s. Microstructure of the polished-surface and fracture-surface of the samples was characterized by Scanning Electron Microscope (SEM, S-2700, Hitachi, Japan) at 15 kV and 10 mA, and the element composition was conducted by Energy-Dispersive X-Ray Spectrometer (EDS).
189
A B
Before infiltration with silicon
After infiltration with silicon
(g/cm3 )
Po (%)
(g/cm3 )
Po (%)
1.48 1.67
24 20
1.90 2.18
11 8
3. Results and discussion 3.1. Microstructure and phase composition The effect of infiltration of TiC on open porosity and density of samples A and B is shown in Table 2. Density and open porosity of sample A before infiltration with liquid silicon were 1.48 g/cm3 and 24%, respectively, and those of sample B before infiltration with liquid silicon were 1.67 g/cm3 and 20%, respectively. The increase in density and the decrease in open porosity of sample B were owing to the filling of 4 vol% TiC in the C/C composites. TiC particles mainly distributed in the interbundle pores of the C/C composites. After liquid silicon infiltration, the porosity of sample B was slightly lower than that of sample A owing to the good wetting ability of liquid silicon on TiC particles, which led to a larger capillary force driving the infiltration of liquid silicon through the micron-size pores among TiC particles. Therefore, the existence of TiC particles facilitated the infiltration of liquid silicon into the porous C/C preform. After infiltration of silicon, the density of sample B, 2.18 g/cm3 , was higher than that of sample A, 1.90 g/cm3 . Both carbon matrix and TiC particles of sample B reacted with silicon melt at high temperature, leading to the formation of Ti3 SiC2 and SiC, which can be described as: 9TiC(s) + 8Si(l) + 2C(s) = 3Ti3 SiC2 (s) + 5SiC(s) G1773 = −338.7 kJ/mol
K
(2)
Known from the thermodynamic calculations, the formation of Ti3 SiC2 in the existence of carbon according to Eq. (2) is much more favorable than that without carbon. The XRD analysis confirms the formation of Ti3 SiC2 in sample B. As shown in Fig. 1, sample A was composed of carbon, SiC and residual silicon. As a comparison, Ti3 SiC2 was detected in sample B, instead of residual silicon. The contents of carbon, silicon, and SiC
Fig. 1. XRD patterns of samples A and B.
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Fig. 2. Backscattered electron (BSE) images of morphologies of (a) sample A and (b) sample B.
in the composites were determined by gravimetric analysis. Carbon phase was removed by burning it off at 700 ◦ C for 20 h in air, and silicon was removed by dissolving the samples in a mixture of hydrofluoric and nitric acid (HNO3 :HF = 4:1) at 25 ◦ C for 48 h. By analyzing the weight variation during oxidation and etching, the content of carbon, silicon and SiC can be determined. For sample B, the residual phases after removing carbon phase were Ti3 SiC2 and SiC. The content of Ti3 SiC2 was calculated from Eq. (2) by the content of TiC, which was measured after slurry infiltration. According to the above way, it was revealed that sample A was composed of 70 vol% carbon, 22 vol% SiC and 8 vol% silicon, and sample B was composed of 73 vol% carbon, 22 vol% SiC and 5 vol% Ti3 SiC2 . Fig. 2 shows the typical microstructure of samples A and B observed in backscattered electron (BSE) images. Both samples A and B were composed of the layers of 0◦ non-woven fiber cloth, short fiber web, 90◦ non-woven fiber cloth, which were repeatedly overlapped, and restricted by needle fibers. Confirmed by EDS analysis, SiC and silicon in sample A were mainly distributed in the short fiber web layers and the inter-bundle matrix (Fig. 2a), as well as SiC and Ti3 SiC2 in sample B, as shown in Fig. 2b. SEM images of fracture surface of sample B are presented in Fig. 3, which further revealed that Ti3 SiC2 and SiC distributed in the matrix among carbon fiber bundles (Fig. 3a). The high magnification SEM image shows the typical nanolaminated structure of Ti3 SiC2 (Fig. 3b). Though the content of Ti3 SiC2 was only 5 vol%, it distributed in the SiC matrix, and the volume ratio between Ti3 SiC2 and SiC was approximately 1:4. Therefore, the properties of the C/C-SiC composites can be altered by the formation of Ti3 SiC2 owing to the variation in the composition and microstructure of the ceramic matrix. The existence of Ti3 SiC2 had been proved to be beneficial to improving the mechanical properties of C/C-SiC composites, which
Fig. 3. SEM images of fracture surface of sample B, showing (a) the distribution of Ti3 SiC2 in the matrix, and (b) the nanolaminated structure of Ti3 SiC2 .
exhibited characteristics of delamination, deformation, kinking, layered fracture and intergranular cracking during bending [11]. 3.2. The improved friction and wear properties by the introduction of Ti3 SiC2 The variation curves of COF of samples A and B at different initial braking speeds under 0.8 MPa pressure are shown in Fig. 4, which had shown similar tendency. At lower braking speeds, both of samples A and B showed a persistent increase of the COF. When the braking speed exceeded 20 m/s, the COF curve exhibited a “saddle” shape, a sharp peak appeared at the start of the curve, a relatively smooth middle stage, and the COF reached a highest value at the end of braking, which decreased with increasing braking speed. The friction properties depend on the friction surface of the samples. As shown in Fig. 5, there existed more grooves on the surface of sample A after the braking test at 28 m/s under 0.8 MPa pressure. As a comparison, the friction surface of sample B was relatively smooth. The grooves had an important effect on the friction behaviors. During braking, the micro-peaks composed of brittle SiC and carbon were pressed to contact each other, first under pressing stress and then together with shear stress, which spalled off into large debris. SiC particles in the debris had high hardness, which resulted in the formation of plough grooves, leading to a high COF. When the asperities were destroyed and the debris was grinded and filled into the grooves, the COF decreased. With increasing initial braking speed, the temperature on the surface of the braking disks increased, as shown in Fig. 6. Since high temperature was produced, the degree of adhesion of the contact surfaces was enhanced, which led to a quick increase of COF at the end of braking. Ti3 SiC2 has a much lower hardness than SiC, so it can reduce the formation of plough grooves and be beneficial to filling the grooves.
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191
Fig. 4. Friction coefficient curves of (a) sample A and (b) sample B with different braking speeds under 0.8 MPa pressure.
The variation of cp with number of braking tests at 28 m/s is shown in Fig. 7. From the calculated data according to 15 times of braking, at 28 m/s under 0.8 MPa, the cp of sample A was changed from 0.33 to 0.26, but the cp of sample B was changed from 0.30 to 0.27, which had shown more stable than that of sample A. Known from the calculated results, the average S of samples A and B under 0.8 MPa pressure were 0.43 and 0.48, respectively, which confirmed that the COF of sample B was more stable than that of sample A. The above results indicated that the existence of residual silicon could cause the instability in the friction properties of C/C-SiC composite, which is consistent with the previous report [3]. For sample B, residual silicon was replaced by the in situ formed Ti3 SiC2 , and it can be found that the stability of the friction behavior was improved efficiently. The weight wear rates of samples A and B are shown in Fig. 8. Both samples A and B showed increasing tendency in weight wear rate with increasing initial braking speed. Under the same braking pressure 0.8 MPa, the weight wear rate of sample B were much lower than those of sample A at higher braking speeds, and the difference increased significantly with increasing initial braking speed. At 28 m/s under 0.8 MPa pressure, the weight wear rate of sample B was 5.55 mg/cycle, but the weight wear rate of sample A was 16.43 mg/cycle, which was almost two times higher than that of sample B. On the other side, the variation of the weight wear rate for sample B was less sensitive to the increase of braking speed compared to that of sample A. The weight wear rates of samples A and B at 28 m/s under different initial braking speeds are shown in Table 3, which were compared to those in the previous literatures. The weight wear rates of sample B at 28 m/s under 0.3 MPa and 0.8 MPa pressure
Fig. 5. Macroscopic friction surface of (a) sample A and (b) sample B after the braking test at 28 m/s under 0.8 MPa pressure.
Fig. 6. Comparison on the temperature of rubbing surfaces of samples A and B under different braking pressures.
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Table 3 The weight wear rate of sample A and sample B compared to previous literature at an initial braking speed of 28 m/s. Samples
Density (g/cm3 )
Braking pressure (MPa)
Average S
cp
Wear rate (mg/cycle)
f (MPa)
A B
1.90 2.18
1.93 2.10
0.43 ± 0.02 0.48 ± 0.01 0.49 ± 0.02 0.48 ± 0.02 0.56 0.42
0.29 ± 0.02 0.35 ± 0.01 0.33 ± 0.01 0.29 ± 0.01 0.36 0.32
16.43 5.59 5.36 5.55 22.5 6.66
178 200 [11]
Ref. [5] Ref. [12]
0.8 0.3 0.5 0.8 0.8 0.3
Fig. 7. cp as a function of number of braking tests at an initial braking speed of 28 m/s under different braking pressures.
were 5.59 mg/cycle and 5.55 mg/cycle, respectively, which were both lower than the weight wear rate in Ref. [12] (6.66 mg/cycle) and Ref. [5] (22.5 mg/cycle). In addition, the mechanical property of sample B was higher than others, which have been explained in detail in the previous work [11]. Known from the above results, the introduction of Ti3 SiC2 had improved the stability of the friction behavior and reduced efficiently the wear rate of C/C-SiC composites. In the present work, both materials have the same reinforcement geometry and the content of carbon fibers, and the similar content of carbon and SiC matrix. The considerable difference was that sample A contained 8 vol% residual silicon and sample B had
Fig. 8. Effect of braking pressure and braking speed on the weight wear rate of samples A and B.
– 82 [13]
5 vol% Ti3 SiC2 . As a result, both friction and wear properties of the latter were improved by the introduction of Ti3 SiC2 matrix, which replaced residual silicon. Ti3 SiC2 as a self-lubricious material [9] may be oxidized at the temperature beyond 400 ◦ C [14]. During braking, the temperature of the friction surface increased with increasing initial braking speed (Fig. 6). The hole in the brake disk for measuring temperature by the thermocouples had a certain distance away from the friction surface, so the measured data was lower than the actual temperature of the friction surface. Additionally, there were some hot spots existed on the friction surface. Therefore, the oxidation of Ti3 SiC2 may start even though the measurement temperature was lower than 400 ◦ C, which led to the formation of self-lubricious oxide and improved the friction and wear properties. After braking under 0.8 MPa pressure, the debris ejected from the disks was characterized by SEM microscopy (Fig. 9). For sample B, the size of the debris became gradually smaller until the initial braking speed increased to 15 m/s, and then the size increased with increasing initial braking speed. This is different from the trend of sample A, which became small gradually due to the shear stress increasing as increasing initial braking speed [15]. The considerable difference between the debris of sample B and sample A can be found in Fig. 9d and e. The debris of sample A was mainly composed of fine particles, and the debris of sample B was composed of film-like ones (Fig. 9d). The existence of filmlike debris on sample B was consistent with its improved friction and wear properties. In the previous report, the introduction of flake graphite made the friction surface smooth and be easy to form the film-type debris [6]. Ti3 SiC2 has the same nanolaminate structure like graphite, and also has similar plasticity and lubrication with graphite. Similarly, the introduction of Ti3 SiC2 also led to the formation of the film-like debris, which had been demonstrated by Fig. 9d. Because the lubrication effect of Ti3 SiC2 was decided by its oxide, the film-like debris was only seen at high braking speeds due to their enough high temperature on the friction surface. For the C/C-SiC composites, the wear rate increased with increasing initial braking speed. Especially at high braking speeds, the oxidation caused by increasing temperature will accelerate the wear, making the wear rate increase quickly. Both samples A and B had shown the similar trend. However, for sample B, the film-like debris was formed by the oxidation of Ti3 SiC2 , and the increasing trend of the weight wear rate was slowed down. Therefore, at lower initial braking speeds, the surface temperature was not high enough to lead to the oxidation of Ti3 SiC2 , the difference of the weight wear rate between sample A and sample B was very little. Once the temperature was increased to an enough high one, the self-lubricious film-like debris was formed due to the oxidation of Ti3 SiC2 , which made the difference of the weight wear rate between sample A and sample B become obvious. Fig. 10 presents the EDS patterns of wear debris of samples A and B at 28 m/s under 0.8 MPa pressure. There existed oxygen in the debris of samples A and B, which implied the oxidation of the debris. Some metals, i.e. Ti, Zn, W, Ta, Al, Ni, etc. and their alloys and several carbides and nitrides such as TiC, TiN, TiCN and
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Fig. 9. The typical micrographs of debris at (a) 5 m/s, (b) 15 m/s, (c) 20 m/s, (d) 28 m/s for sample B and (e) 28 m/s for sample A under 0.8 MPa pressure.
Ti3 SiC2 exhibit good tribological properties at elevated temperatures owing to the formation of lubricious oxides by tribo-oxidation [16–18]. In Ref. [18], tribofilms, which were mainly comprised of amorphous oxides of the M and A elements were formed on the contact surfaces of MAX phase and Al2 O3 . SiC has higher hardness than Al2 O3 , so it is possible to form amorphous oxides between Ti3 SiC2 and SiC. As a comparison, no self-lubricious oxide could be generated during braking for C/C-SiC composites, and the existence of silicon would increase the wear and lead to the instability [4,5]. Therefore, the weight wear rate of C/C-SiC composites was much higher than that of composites containing Ti3 SiC2 at high braking speeds. Owing to the replacement of residual silicon by Ti3 SiC2 , which led to the formation of self-lubricious film-like debris at high braking speeds, the friction and wear properties were improved.
3.3. Effect of braking speed and braking pressure on tribological properties of C/C-SiC composites containing Ti3 SiC2 The effect of braking speed and braking pressure on the cp is shown in Fig. 11. Under the same braking pressure, the cp increased to a maximum value at 15 m/s, and then decreased with increasing braking speed. The braking energy increased with the increase of braking speed, which made the friction force increased. When the speed was lower than 15 m/s, the friction force was not enough to shear the asperities completed. The area energy in the friction surface increased with the increase of the braking speed, and created the new asperities, which caused the cp increased. When the braking speed was higher than 15 m/s, the friction force between the friction surfaces increased enough to shear the asperities completely, which could cut the asperities and form a lot of debris. The formed debris could rapidly fill the grooves on
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Fig. 12. Effect of braking pressure and braking speed on the braking time of sample B.
ing speed, the friction force increased, which would accelerate the plough effect of the friction surface. The braking time also increased with increasing braking speed under the same braking pressure (Fig. 12), so that the wear rate increased constantly with the increase of braking speed under the same pressure. With increasing braking pressure, the friction force increased, so the wear rate in unit time increased. At the same time, the braking time decreased with increasing braking pressure at the same braking speed (Fig. 12). Therefore, the wear rate was affected less by braking pressure. Fig. 10. EDS of the debris of (a) sample A and (b) sample B at an initial braking speed of 28 m/s under 0.8 MPa pressure.
the friction surfaces, and reduced the plough effect, which made the cp decreased. At the same braking speed, the cp decreased with the increase of the braking pressure. The friction force rapidly increased with increasing braking pressure, and made the debris formed quickly, and then filled the grooves and reduced the plough effect, which made the cp decreased with the increase of the braking pressure. The effect of braking speed and braking pressure on the wear rate is investigated. As shown in Fig. 8, the wear rate increased constantly with increasing braking speed. With increasing brak-
Fig. 11. Effect of braking pressure and braking speed on the friction coefficient of sample B.
4. Conclusions (1) C/C-SiC composites and C/C-SiC composites containing Ti3 SiC2 were fabricated using C/C preforms by the process of liquid silicon infiltration and a joint process of slurry infiltration and liquid silicon infiltration, respectively. Through the chemical analysis combined with the calculation, it can be known that C/C-SiC composites were composed of 70 vol% carbon, 22 vol% SiC and 8 vol% silicon, and the composites containing Ti3 SiC2 were composed of 73 vol% carbon, 22 vol% SiC and 5 vol% Ti3 SiC2 . Ti3 SiC2 distributed in the inter-bundle SiC matrix and showed the typical nanolaminated structure. (2) The improved friction and wear properties of the composites containing Ti3 SiC2 were investigated. The weight wear rate of the composites was reduced obviously by the introduction of Ti3 SiC2 and therefore the absence of silicon, which would otherwise increase the wear rate. At an initial braking speed of 28 m/s under 0.8 MPa pressure, the weight wear rate of the composites containing Ti3 SiC2 was 5.55 mg/cycle, which was one-third of the weight wear rate of C/C-SiC composites. The COF of the composites containing Ti3 SiC2 had shown more stable than the C/C-SiC composites. Self-lubricious film-like debris was formed on the composites containing Ti3 SiC2 , which led to the improvement of the wear and friction properties. (3) The effect of braking speed and braking pressure on the tribological properties on the C/C-SiC composites containing Ti3 SiC2 were studied. Under the same pressure, the cp increased to maximum value at 15 m/s, and then decreased with the increase of braking speed. The cp decreased with increasing braking pressure at the same braking speed. The wear rate increased constantly with the increase of braking speed under the same braking pressure, but was less affected by braking pressure at the same braking speed.
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Acknowledgements The authors acknowledge the Natural Science Foundation of China (Grant: 50802074), the 863 National High-tech Research Development Plan (863 plan) (Project No. 2007AA03Z542) and the Program for New Century Excellent Talents in University (NCET08-0465) for the financial supports. References [1] W. Krenkel, B. Heidenreich, R. Renz, C/C-SiC composites for advanced friction systems, Adv. Eng. Mater. 4 (2002) 427–436. [2] S. Fan, L. Zhang, Y. Xu, L. Cheng, J. Lou, J. Zhang, L. Yu, Microstructure and properties of 3D needle-punched carbon/silicon carbide brake materials, Compos. Sci. Technol. 67 (2007) 2390–2398. [3] S. Fouquet, M. Rollin, R. Pailler, X. Bourrat, Tribological behaviour of composites made of carbon fibres and ceramic matrix in the Si–C system, Wear 264 (2008) 850–856. [4] Y. Zhang, Z. Xiao, J. Wang, R. Xing, Z. Peng, J. Su, Z. Jin, Effect of pyrocarbon content on thermal and frictional properties in C/C preforms of C/C-SiC composites, Wear 269 (2010) 132–138. [5] G. Jiang, J. Yang, Y. Xu, J. Gao, J. Zhang, L. Zhang, L. Cheng, J. Lou, Effect of graphitization on microstructure and tribological properties of C/SiC composites prepared by reactive melt infiltration, Compos. Sci. Technol. 68 (2008) 2468–2473. [6] Y. Cai, S. Fan, H. Liu, L. Zhang, L. Cheng, B. Dong, J. Jiang, Microstructures and improved wear resistance of 3D needled C/SiC composites with graphite filler, Compos. Sci. Technol. 69 (2009) 2447–2453.
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Wear journal homepage: www.elsevier.com/locate/wear
Friction and wear behaviors of C/C-SiC composites containing Ti3 SiC2 Xiaomeng Fan, Xiaowei Yin ∗ , Shanshan He, Litong Zhang, Laifei Cheng Northwestern Polytechnical University, National Key Laboratory of Thermostructure Composite Materials, Xi’an, Shaanxi 710072, PR China
a r t i c l e
i n f o
Article history: Received 1 August 2010 Received in revised form 22 August 2011 Accepted 23 August 2011 Available online 30 August 2011 Keywords: Ceramic–matrix composite Ti3 SiC2 MAX-phase Brake materials Wear testing
a b s t r a c t In the present paper, friction and wear behaviors of a carbon fiber reinforced carbon–silicon carbide–titanium silicon carbide (C-SiC–Ti3 SiC2 ) hybrid matrix composites fabricated by slurry infiltration and liquid silicon infiltration were studied for potential application as brake materials. The properties were compared with those of C/C-SiC composites. The composites containing Ti3 SiC2 had not only higher friction stability coefficient but also much higher wear resistance than C/C-SiC composites. At an initial braking speed of 28 m/s under 0.8 MPa pressure, the weight wear rate of the composites containing 5 vol% Ti3 SiC2 was 5.55 mg/cycle, which was only one-third of C/C-SiC composites. Self-lubricious filmlike debris was formed on the composites containing Ti3 SiC2 , leading to the improvement of friction and wear properties. The effect of braking speed and braking pressure on the tribological properties of modified composites were investigated. The average friction coefficient was significantly affected by braking speed and braking pressure, but the wear rate was less affected by braking pressure. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Carbon fiber reinforced carbon/silicon carbide binary matrix (C/C-SiC) composites fabricated by liquid silicon infiltration process (LSI) possess excellent friction properties and low sensitivity to service environments, making them promising brake materials, which is also named as C/SiC composites [1–3]. The LSI C/C-SiC composites comprise carbon fibers, C-SiC binary matrix and more than 8 vol% residual silicon mainly locating in the interbundle pores [2]. As a residual phase in brake materials, silicon in the matrix increases the wear and leads to the instability of friction behaviors [3,4]. This is supposed to be avoided by replacing silicon with alloys to give stable silicides or carbides [3]. The effect of graphitization on tribological properties of C/CSiC composites was studied, and it was found the stability of COF increased significantly [5]. Graphite filler was added into the C/CSiC brake materials, it was found that C/C-SiC composites with graphite filler was relatively insensitive to changes in braking speeds and displayed higher braking performance at high braking speeds due to the formation of the film-type wear debris [6]. Owing to its special nanolaminate structure and excellent properties, Ti3 SiC2 is extensively researched in recent years [7]. Ti3 SiC2 ceramics have been proved to be promising tribological materials [8–10]. The dependence of friction and wear properties of Ti3 SiC2 material on its grain size was investigated [8]. The wear resistance increases with increasing grain size for sliding and abrasive
∗ Corresponding author. Tel.: +86 29 88494947; fax: +86 29 88494620. E-mail address: [email protected] (X. Yin). 0043-1648/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.wear.2011.08.029
wear tests. Under dry conditions at room temperature, the friction coefficient of Ti3 SiC2 /Ti3 SiC2 pair was 1.16–1.43, but that of Ti3 SiC2 plate/diamond pin pair was 0.05–0.1, which was even lower than that of graphite plate/diamond pin pair [9]. The difference may depend on whether lubricious oxide tribofilms, TiCx Oy , can be formed or not [10]. For the Ti3 SiC2 ceramics, the oxide is important for its lubricating effect in the tribological tests. In our previous work [11], in order to eliminate residual free silicon in the LSI C/C-SiC composites, forming Ti3 SiC2 , a porous C/C preform fabricated by CVI was infiltrated with TiC particles, and then the preform containing TiC was infiltrated with liquid silicon, resulting in the formation of a carbon fiber reinforced CSiC–Ti3 SiC2 matrix composite. The effects of fabrication process on phase composition and microstructure of the composites were studied, and mechanical properties of the composites containing Ti3 SiC2 were compared with those of C/C and C/C-SiC composites, which indicated the introduction of Ti3 SiC2 was beneficial to improve the mechanical properties of C/C-SiC composites. The aim of the present study was to understand friction and wear behaviors of C/C-SiC composites containing Ti3 SiC2 , which was compared with those of C/C-SiC composites. 2. Experimental procedures 2.1. Sample preparation The fabrication procedure of C/C-SiC–Ti3 SiC2 composites included three steps. In the first step, porous C/C composite was prepared by chemical vapor infiltration (CVI) using 3D needled fiber felts as preform, and propylene was used as the precursor of
X. Fan et al. / Wear 274–275 (2012) 188–195 Table 1 Braking conditions for friction and wear performance tests. Test conditions
I
Inertia (kg m2 ) Braking pressure (MPa) Braking speeda (m/s) Braking timesb
0.235 0.3, 0.5, 0.8 5 10 30 30
a b
II
III
IV
V
VI
20 15
25 15
28 15
Braking speed: the linear velocity of outer edge of rotate disk. Braking times: the number of repeated test under the applied conditions.
carbon and Ar was used as carrying and diluting gas. In the second step, the CVI C/C composite was infiltrated by TiC slurry using the process parameter depicted in the previous work [11], which led to the formation of C/C-TiC composite. At last, the as-received C/C-TiC composites and C/C composites were infiltrated by liquid silicon at 1500 ◦ C in a vacuum furnace, and the C/C-SiC–Ti3 SiC2 composites and C/C-SiC composites were obtained. For ease of the comparison, the C/C-SiC composites and the C/C-SiC–Ti3 SiC2 composites were designated as sample A and sample B. 2.2. Characterizations of materials 2.2.1. Friction and wear properties Both samples A and B were machined into 76.5 mm in outer diameter, 52.6 mm in inner diameter and 12 mm in thickness as rotating disk, and into 90 mm in outer diameter, 55 mm in inner diameter and 12 mm in thickness as stationary disk, respectively. Friction and wear properties of the samples were performed on a disk-on-disk MM-1000 testing machine (Xi’an ShunTong Technical Research Institute, China), which was described previously in detail [2]. The braking conditions for friction and wear performance tests are shown in Table 1. The friction stability coefficient, S, during each braking test can be calculated from Eq. (1). S=
cp max
Table 2 Density and open porosity of samples A and B before and after infiltration with liquid silicon. Samples
15 30
(1)
where max is the maximum COF during each braking test, and cp is the average COF, which is automatically recorded by the machine. The temperature of the samples during braking was on-line measured using NiCr–NiSi thermal–electrical couples, which were inserted into a hole of 2 mm in diameter and 7 mm in depth of the stationary disk, and the distance between the hole and the friction surface was 2 mm. In every braking cycle, the braking speed of the brake disk fell from a certain speed to static, and the energy was transformed from kinetic energy into heat energy. The kinetic energy was determined by the braking speed, so that the kinetic energy increased with increasing braking speed under the same braking pressure. The friction resistance was determined by braking pressure, so that the friction force increased with the increase of braking pressure at the same braking speed. 2.2.2. Open porosity, density, phase composition, and microstructure The open porosity and the bulk density of the samples were measured by Archimedes method according to ASTM C-20 standard. The samples were crushed into power and analyzed by X-ray diffraction (XRD), via a computer-controlled diffractometer (X’Pert Pro, Philips, Netherlands) with Cu K␣ radiation at 40 kV and 100 mA. Data was digitally recorded in a continuous scanning mode in the angle (2) ranging from 8 to 80◦ with a scanning rate of 0.12◦ /s. Microstructure of the polished-surface and fracture-surface of the samples was characterized by Scanning Electron Microscope (SEM, S-2700, Hitachi, Japan) at 15 kV and 10 mA, and the element composition was conducted by Energy-Dispersive X-Ray Spectrometer (EDS).
189
A B
Before infiltration with silicon
After infiltration with silicon
(g/cm3 )
Po (%)
(g/cm3 )
Po (%)
1.48 1.67
24 20
1.90 2.18
11 8
3. Results and discussion 3.1. Microstructure and phase composition The effect of infiltration of TiC on open porosity and density of samples A and B is shown in Table 2. Density and open porosity of sample A before infiltration with liquid silicon were 1.48 g/cm3 and 24%, respectively, and those of sample B before infiltration with liquid silicon were 1.67 g/cm3 and 20%, respectively. The increase in density and the decrease in open porosity of sample B were owing to the filling of 4 vol% TiC in the C/C composites. TiC particles mainly distributed in the interbundle pores of the C/C composites. After liquid silicon infiltration, the porosity of sample B was slightly lower than that of sample A owing to the good wetting ability of liquid silicon on TiC particles, which led to a larger capillary force driving the infiltration of liquid silicon through the micron-size pores among TiC particles. Therefore, the existence of TiC particles facilitated the infiltration of liquid silicon into the porous C/C preform. After infiltration of silicon, the density of sample B, 2.18 g/cm3 , was higher than that of sample A, 1.90 g/cm3 . Both carbon matrix and TiC particles of sample B reacted with silicon melt at high temperature, leading to the formation of Ti3 SiC2 and SiC, which can be described as: 9TiC(s) + 8Si(l) + 2C(s) = 3Ti3 SiC2 (s) + 5SiC(s) G1773 = −338.7 kJ/mol
K
(2)
Known from the thermodynamic calculations, the formation of Ti3 SiC2 in the existence of carbon according to Eq. (2) is much more favorable than that without carbon. The XRD analysis confirms the formation of Ti3 SiC2 in sample B. As shown in Fig. 1, sample A was composed of carbon, SiC and residual silicon. As a comparison, Ti3 SiC2 was detected in sample B, instead of residual silicon. The contents of carbon, silicon, and SiC
Fig. 1. XRD patterns of samples A and B.
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Fig. 2. Backscattered electron (BSE) images of morphologies of (a) sample A and (b) sample B.
in the composites were determined by gravimetric analysis. Carbon phase was removed by burning it off at 700 ◦ C for 20 h in air, and silicon was removed by dissolving the samples in a mixture of hydrofluoric and nitric acid (HNO3 :HF = 4:1) at 25 ◦ C for 48 h. By analyzing the weight variation during oxidation and etching, the content of carbon, silicon and SiC can be determined. For sample B, the residual phases after removing carbon phase were Ti3 SiC2 and SiC. The content of Ti3 SiC2 was calculated from Eq. (2) by the content of TiC, which was measured after slurry infiltration. According to the above way, it was revealed that sample A was composed of 70 vol% carbon, 22 vol% SiC and 8 vol% silicon, and sample B was composed of 73 vol% carbon, 22 vol% SiC and 5 vol% Ti3 SiC2 . Fig. 2 shows the typical microstructure of samples A and B observed in backscattered electron (BSE) images. Both samples A and B were composed of the layers of 0◦ non-woven fiber cloth, short fiber web, 90◦ non-woven fiber cloth, which were repeatedly overlapped, and restricted by needle fibers. Confirmed by EDS analysis, SiC and silicon in sample A were mainly distributed in the short fiber web layers and the inter-bundle matrix (Fig. 2a), as well as SiC and Ti3 SiC2 in sample B, as shown in Fig. 2b. SEM images of fracture surface of sample B are presented in Fig. 3, which further revealed that Ti3 SiC2 and SiC distributed in the matrix among carbon fiber bundles (Fig. 3a). The high magnification SEM image shows the typical nanolaminated structure of Ti3 SiC2 (Fig. 3b). Though the content of Ti3 SiC2 was only 5 vol%, it distributed in the SiC matrix, and the volume ratio between Ti3 SiC2 and SiC was approximately 1:4. Therefore, the properties of the C/C-SiC composites can be altered by the formation of Ti3 SiC2 owing to the variation in the composition and microstructure of the ceramic matrix. The existence of Ti3 SiC2 had been proved to be beneficial to improving the mechanical properties of C/C-SiC composites, which
Fig. 3. SEM images of fracture surface of sample B, showing (a) the distribution of Ti3 SiC2 in the matrix, and (b) the nanolaminated structure of Ti3 SiC2 .
exhibited characteristics of delamination, deformation, kinking, layered fracture and intergranular cracking during bending [11]. 3.2. The improved friction and wear properties by the introduction of Ti3 SiC2 The variation curves of COF of samples A and B at different initial braking speeds under 0.8 MPa pressure are shown in Fig. 4, which had shown similar tendency. At lower braking speeds, both of samples A and B showed a persistent increase of the COF. When the braking speed exceeded 20 m/s, the COF curve exhibited a “saddle” shape, a sharp peak appeared at the start of the curve, a relatively smooth middle stage, and the COF reached a highest value at the end of braking, which decreased with increasing braking speed. The friction properties depend on the friction surface of the samples. As shown in Fig. 5, there existed more grooves on the surface of sample A after the braking test at 28 m/s under 0.8 MPa pressure. As a comparison, the friction surface of sample B was relatively smooth. The grooves had an important effect on the friction behaviors. During braking, the micro-peaks composed of brittle SiC and carbon were pressed to contact each other, first under pressing stress and then together with shear stress, which spalled off into large debris. SiC particles in the debris had high hardness, which resulted in the formation of plough grooves, leading to a high COF. When the asperities were destroyed and the debris was grinded and filled into the grooves, the COF decreased. With increasing initial braking speed, the temperature on the surface of the braking disks increased, as shown in Fig. 6. Since high temperature was produced, the degree of adhesion of the contact surfaces was enhanced, which led to a quick increase of COF at the end of braking. Ti3 SiC2 has a much lower hardness than SiC, so it can reduce the formation of plough grooves and be beneficial to filling the grooves.
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191
Fig. 4. Friction coefficient curves of (a) sample A and (b) sample B with different braking speeds under 0.8 MPa pressure.
The variation of cp with number of braking tests at 28 m/s is shown in Fig. 7. From the calculated data according to 15 times of braking, at 28 m/s under 0.8 MPa, the cp of sample A was changed from 0.33 to 0.26, but the cp of sample B was changed from 0.30 to 0.27, which had shown more stable than that of sample A. Known from the calculated results, the average S of samples A and B under 0.8 MPa pressure were 0.43 and 0.48, respectively, which confirmed that the COF of sample B was more stable than that of sample A. The above results indicated that the existence of residual silicon could cause the instability in the friction properties of C/C-SiC composite, which is consistent with the previous report [3]. For sample B, residual silicon was replaced by the in situ formed Ti3 SiC2 , and it can be found that the stability of the friction behavior was improved efficiently. The weight wear rates of samples A and B are shown in Fig. 8. Both samples A and B showed increasing tendency in weight wear rate with increasing initial braking speed. Under the same braking pressure 0.8 MPa, the weight wear rate of sample B were much lower than those of sample A at higher braking speeds, and the difference increased significantly with increasing initial braking speed. At 28 m/s under 0.8 MPa pressure, the weight wear rate of sample B was 5.55 mg/cycle, but the weight wear rate of sample A was 16.43 mg/cycle, which was almost two times higher than that of sample B. On the other side, the variation of the weight wear rate for sample B was less sensitive to the increase of braking speed compared to that of sample A. The weight wear rates of samples A and B at 28 m/s under different initial braking speeds are shown in Table 3, which were compared to those in the previous literatures. The weight wear rates of sample B at 28 m/s under 0.3 MPa and 0.8 MPa pressure
Fig. 5. Macroscopic friction surface of (a) sample A and (b) sample B after the braking test at 28 m/s under 0.8 MPa pressure.
Fig. 6. Comparison on the temperature of rubbing surfaces of samples A and B under different braking pressures.
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Table 3 The weight wear rate of sample A and sample B compared to previous literature at an initial braking speed of 28 m/s. Samples
Density (g/cm3 )
Braking pressure (MPa)
Average S
cp
Wear rate (mg/cycle)
f (MPa)
A B
1.90 2.18
1.93 2.10
0.43 ± 0.02 0.48 ± 0.01 0.49 ± 0.02 0.48 ± 0.02 0.56 0.42
0.29 ± 0.02 0.35 ± 0.01 0.33 ± 0.01 0.29 ± 0.01 0.36 0.32
16.43 5.59 5.36 5.55 22.5 6.66
178 200 [11]
Ref. [5] Ref. [12]
0.8 0.3 0.5 0.8 0.8 0.3
Fig. 7. cp as a function of number of braking tests at an initial braking speed of 28 m/s under different braking pressures.
were 5.59 mg/cycle and 5.55 mg/cycle, respectively, which were both lower than the weight wear rate in Ref. [12] (6.66 mg/cycle) and Ref. [5] (22.5 mg/cycle). In addition, the mechanical property of sample B was higher than others, which have been explained in detail in the previous work [11]. Known from the above results, the introduction of Ti3 SiC2 had improved the stability of the friction behavior and reduced efficiently the wear rate of C/C-SiC composites. In the present work, both materials have the same reinforcement geometry and the content of carbon fibers, and the similar content of carbon and SiC matrix. The considerable difference was that sample A contained 8 vol% residual silicon and sample B had
Fig. 8. Effect of braking pressure and braking speed on the weight wear rate of samples A and B.
– 82 [13]
5 vol% Ti3 SiC2 . As a result, both friction and wear properties of the latter were improved by the introduction of Ti3 SiC2 matrix, which replaced residual silicon. Ti3 SiC2 as a self-lubricious material [9] may be oxidized at the temperature beyond 400 ◦ C [14]. During braking, the temperature of the friction surface increased with increasing initial braking speed (Fig. 6). The hole in the brake disk for measuring temperature by the thermocouples had a certain distance away from the friction surface, so the measured data was lower than the actual temperature of the friction surface. Additionally, there were some hot spots existed on the friction surface. Therefore, the oxidation of Ti3 SiC2 may start even though the measurement temperature was lower than 400 ◦ C, which led to the formation of self-lubricious oxide and improved the friction and wear properties. After braking under 0.8 MPa pressure, the debris ejected from the disks was characterized by SEM microscopy (Fig. 9). For sample B, the size of the debris became gradually smaller until the initial braking speed increased to 15 m/s, and then the size increased with increasing initial braking speed. This is different from the trend of sample A, which became small gradually due to the shear stress increasing as increasing initial braking speed [15]. The considerable difference between the debris of sample B and sample A can be found in Fig. 9d and e. The debris of sample A was mainly composed of fine particles, and the debris of sample B was composed of film-like ones (Fig. 9d). The existence of filmlike debris on sample B was consistent with its improved friction and wear properties. In the previous report, the introduction of flake graphite made the friction surface smooth and be easy to form the film-type debris [6]. Ti3 SiC2 has the same nanolaminate structure like graphite, and also has similar plasticity and lubrication with graphite. Similarly, the introduction of Ti3 SiC2 also led to the formation of the film-like debris, which had been demonstrated by Fig. 9d. Because the lubrication effect of Ti3 SiC2 was decided by its oxide, the film-like debris was only seen at high braking speeds due to their enough high temperature on the friction surface. For the C/C-SiC composites, the wear rate increased with increasing initial braking speed. Especially at high braking speeds, the oxidation caused by increasing temperature will accelerate the wear, making the wear rate increase quickly. Both samples A and B had shown the similar trend. However, for sample B, the film-like debris was formed by the oxidation of Ti3 SiC2 , and the increasing trend of the weight wear rate was slowed down. Therefore, at lower initial braking speeds, the surface temperature was not high enough to lead to the oxidation of Ti3 SiC2 , the difference of the weight wear rate between sample A and sample B was very little. Once the temperature was increased to an enough high one, the self-lubricious film-like debris was formed due to the oxidation of Ti3 SiC2 , which made the difference of the weight wear rate between sample A and sample B become obvious. Fig. 10 presents the EDS patterns of wear debris of samples A and B at 28 m/s under 0.8 MPa pressure. There existed oxygen in the debris of samples A and B, which implied the oxidation of the debris. Some metals, i.e. Ti, Zn, W, Ta, Al, Ni, etc. and their alloys and several carbides and nitrides such as TiC, TiN, TiCN and
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Fig. 9. The typical micrographs of debris at (a) 5 m/s, (b) 15 m/s, (c) 20 m/s, (d) 28 m/s for sample B and (e) 28 m/s for sample A under 0.8 MPa pressure.
Ti3 SiC2 exhibit good tribological properties at elevated temperatures owing to the formation of lubricious oxides by tribo-oxidation [16–18]. In Ref. [18], tribofilms, which were mainly comprised of amorphous oxides of the M and A elements were formed on the contact surfaces of MAX phase and Al2 O3 . SiC has higher hardness than Al2 O3 , so it is possible to form amorphous oxides between Ti3 SiC2 and SiC. As a comparison, no self-lubricious oxide could be generated during braking for C/C-SiC composites, and the existence of silicon would increase the wear and lead to the instability [4,5]. Therefore, the weight wear rate of C/C-SiC composites was much higher than that of composites containing Ti3 SiC2 at high braking speeds. Owing to the replacement of residual silicon by Ti3 SiC2 , which led to the formation of self-lubricious film-like debris at high braking speeds, the friction and wear properties were improved.
3.3. Effect of braking speed and braking pressure on tribological properties of C/C-SiC composites containing Ti3 SiC2 The effect of braking speed and braking pressure on the cp is shown in Fig. 11. Under the same braking pressure, the cp increased to a maximum value at 15 m/s, and then decreased with increasing braking speed. The braking energy increased with the increase of braking speed, which made the friction force increased. When the speed was lower than 15 m/s, the friction force was not enough to shear the asperities completed. The area energy in the friction surface increased with the increase of the braking speed, and created the new asperities, which caused the cp increased. When the braking speed was higher than 15 m/s, the friction force between the friction surfaces increased enough to shear the asperities completely, which could cut the asperities and form a lot of debris. The formed debris could rapidly fill the grooves on
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Fig. 12. Effect of braking pressure and braking speed on the braking time of sample B.
ing speed, the friction force increased, which would accelerate the plough effect of the friction surface. The braking time also increased with increasing braking speed under the same braking pressure (Fig. 12), so that the wear rate increased constantly with the increase of braking speed under the same pressure. With increasing braking pressure, the friction force increased, so the wear rate in unit time increased. At the same time, the braking time decreased with increasing braking pressure at the same braking speed (Fig. 12). Therefore, the wear rate was affected less by braking pressure. Fig. 10. EDS of the debris of (a) sample A and (b) sample B at an initial braking speed of 28 m/s under 0.8 MPa pressure.
the friction surfaces, and reduced the plough effect, which made the cp decreased. At the same braking speed, the cp decreased with the increase of the braking pressure. The friction force rapidly increased with increasing braking pressure, and made the debris formed quickly, and then filled the grooves and reduced the plough effect, which made the cp decreased with the increase of the braking pressure. The effect of braking speed and braking pressure on the wear rate is investigated. As shown in Fig. 8, the wear rate increased constantly with increasing braking speed. With increasing brak-
Fig. 11. Effect of braking pressure and braking speed on the friction coefficient of sample B.
4. Conclusions (1) C/C-SiC composites and C/C-SiC composites containing Ti3 SiC2 were fabricated using C/C preforms by the process of liquid silicon infiltration and a joint process of slurry infiltration and liquid silicon infiltration, respectively. Through the chemical analysis combined with the calculation, it can be known that C/C-SiC composites were composed of 70 vol% carbon, 22 vol% SiC and 8 vol% silicon, and the composites containing Ti3 SiC2 were composed of 73 vol% carbon, 22 vol% SiC and 5 vol% Ti3 SiC2 . Ti3 SiC2 distributed in the inter-bundle SiC matrix and showed the typical nanolaminated structure. (2) The improved friction and wear properties of the composites containing Ti3 SiC2 were investigated. The weight wear rate of the composites was reduced obviously by the introduction of Ti3 SiC2 and therefore the absence of silicon, which would otherwise increase the wear rate. At an initial braking speed of 28 m/s under 0.8 MPa pressure, the weight wear rate of the composites containing Ti3 SiC2 was 5.55 mg/cycle, which was one-third of the weight wear rate of C/C-SiC composites. The COF of the composites containing Ti3 SiC2 had shown more stable than the C/C-SiC composites. Self-lubricious film-like debris was formed on the composites containing Ti3 SiC2 , which led to the improvement of the wear and friction properties. (3) The effect of braking speed and braking pressure on the tribological properties on the C/C-SiC composites containing Ti3 SiC2 were studied. Under the same pressure, the cp increased to maximum value at 15 m/s, and then decreased with the increase of braking speed. The cp decreased with increasing braking pressure at the same braking speed. The wear rate increased constantly with the increase of braking speed under the same braking pressure, but was less affected by braking pressure at the same braking speed.
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Acknowledgements The authors acknowledge the Natural Science Foundation of China (Grant: 50802074), the 863 National High-tech Research Development Plan (863 plan) (Project No. 2007AA03Z542) and the Program for New Century Excellent Talents in University (NCET08-0465) for the financial supports. References [1] W. Krenkel, B. Heidenreich, R. Renz, C/C-SiC composites for advanced friction systems, Adv. Eng. Mater. 4 (2002) 427–436. [2] S. Fan, L. Zhang, Y. Xu, L. Cheng, J. Lou, J. Zhang, L. Yu, Microstructure and properties of 3D needle-punched carbon/silicon carbide brake materials, Compos. Sci. Technol. 67 (2007) 2390–2398. [3] S. Fouquet, M. Rollin, R. Pailler, X. Bourrat, Tribological behaviour of composites made of carbon fibres and ceramic matrix in the Si–C system, Wear 264 (2008) 850–856. [4] Y. Zhang, Z. Xiao, J. Wang, R. Xing, Z. Peng, J. Su, Z. Jin, Effect of pyrocarbon content on thermal and frictional properties in C/C preforms of C/C-SiC composites, Wear 269 (2010) 132–138. [5] G. Jiang, J. Yang, Y. Xu, J. Gao, J. Zhang, L. Zhang, L. Cheng, J. Lou, Effect of graphitization on microstructure and tribological properties of C/SiC composites prepared by reactive melt infiltration, Compos. Sci. Technol. 68 (2008) 2468–2473. [6] Y. Cai, S. Fan, H. Liu, L. Zhang, L. Cheng, B. Dong, J. Jiang, Microstructures and improved wear resistance of 3D needled C/SiC composites with graphite filler, Compos. Sci. Technol. 69 (2009) 2447–2453.
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