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C composites on their tribological properties
NEW CARBON MATERIALS
Volume 34, Issue 5, Oct 2019 Online English edition of the Chinese language journal Cite this article as: New Carbon Materials, 2019, 34(5): 472-481
RESEARCH PAPER
Effects of the high-temperature treatment of C/C composites on their tribological properties Xi Ou-yang1, Zhuan Li1,2*, Peng Xiao1,2, Guan-yi Chen1, jin-wei Li1, Peng-fei Liu1 China State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China;
1
Science and Technology on High Strength Structural Materials Laboratory, Central South University, Changsha 410083, China
2
Abstract:
C/C composites prepared by chemical vapor infiltration (CVI) were subjected to high temperature treatment (HTT) at 2300 °C under
an Ar atmosphere. C/C-SiC composites were prepared by liquid silicon infiltration (LSI) of the C/C composites before HTT. The C/C composites before (A) and after (B) HTT were mated with the C/C-SiC to form friction pairs to investigate their tribological properties. Results indicated that the average coefficient of friction (COF) and stable COF of the C/C-B are 0.280 and 0.65 while that of the C/C-B are 0.451 and 0.55, respectively. The average linear wear rate of the C/C-A and the C/C-B are 3.7 and 8.9 μm/(slide cycle), respectively. SEM reveals that the wear debris of the C/C-A is particulate while that of the C/C-B are the sheets, and the friction surface is more flat for C/C-B than C/C-A. HTT leads to a softening of the carbon, increases the wear rate, but improves the stable COF of the C/C composites. The wear mechanisms of the composites are mainly grain-abrasion, oxidation loss and fatigue wear. Key Words: C/C; C/C-SiC; High-temperature treatment; Tribological properties
1 Introduction With the rapid development of aerospace, military, automobile and high-speed train industry, the speed becomes more and more quick and the load is also heavier than past, which place high demands on the brake materials. Furthermore, brake materials must be better adapted to harsh environments such as sand and rainwater. C/C composites have become one of the excellent choices as brake materials owing to their light weight, excellent thermal and mechanical properties, and high energy absorption [1]. C/C composites were first investigated in 1958 [2, 3] with the purpose to compensate the weaknesses of powder metallurgical materials [4-7], and then aroused the interest of the majority of scientific research workers. Y. Cai et al. [8] have studied the microstructures and mechanical properties of three-dimensional ceramic filler-modified C/C composites. H. Kasem et al. [9] have researched the characterization of surface grooves and scratches of C/C composites induced by friction at low and high temperatures. S. Sarkar et al. [6] have reported the impact abrasive wear response of C/C composites at elevated temperatures. Because of their low-density and excellent high-temperature properties, C/C composites have been introduced in the aircraft brake system [10, 11]. However, C/C composites were commonly considered to have some fatal weakness such as poor oxidation resistance under ambient environment, very low coefficient of friction (COF) in a wet environment and relatively high wear rates. In response to these problems, researchers have done a lot of
improvement work [12-14]. C/C-SiC composites display the theoretically excellent high-temperature mechanical and thermal properties, and now they are widely used in sports cars, luxury sedans and airplane [15-17] . In addition, C/C-SiC composites containing at least 20 wt% of SiC also have many other advantages such as the lower wear rate, longer service life, especially lower sensibility to surroundings and temperature as compared with C/C composites [18, 19]. Since 1990s, W. Krenkel et al. [15, 20, 21] have made a systematic study on C/C-SiC composites in the field of tribology, such as C/C-SiC composites for advanced friction systems, integration of ceramic matrix composite （ CMC ） -brake disks in automotive brake systems and full-ceramic brake systems for high performance friction applications. In order to optimize the properties of C/C-SiC composites, Z. Li et al. [22, 23] have done some modified work. Though C/C-SiC composites have overcome many disadvantages of C/C composites, they also have shortcomings. On the one hand, the COF of self-mated C/C-SiC is too high, which is easy to cause the lock and lead to traffic accidents. On the other hand, C/C-SiC composites show a higher density compared with C/C composites. In this paper, in order to overcome the disadvantages of C/C and C/C-SiC composite brakes, the idea of using C/C mated with C/C-SiC for brake assembly was proposed. Meanwhile, to further enhance the COF stability, one of the
Received date: 10 Aug 2019; Revised date: 30 Sep 2019 *Corresponding author. E-mail: [email protected] Copyright©2019, Institute of Coal Chemistry, Chinese Academy of Sciences. Published by Elsevier Limited. All rights reserved. DOI: 10.1016/S1872-5805(19)60025-9
Xi Ou-yang et al. / New Carbon Materials, 2019, 34(5): 472-481
C/C composites was subjected to high temperature treatment (HTT) at 2 300 °C. R. Luo et al. [24] have reported that the stability of the braking moment-time curves of C/C composites increased with the heat treatment temperature. The correlation of the phase composition, microstructure, friction surfaces to the friction and wear performance was investigated. The characteristics of friction surfaces and wear debris were evaluated and the friction mechanism was proposed.
vacuum. The liquid Si reacted with a small amount of the C/C and formed SiC ceramic matrix. The C/C preform specimen without HTT was named as C/C A, and another one with 2300 °C HTT was named C/C B. The densities and porosities of the C/C A were 1.68 g/cm3 and 8.1%, respectively, those of the C/C B were 1.72 g/cm3 and 14.4% while those of the C/C-SiC were 2.08 g/cm3 and 2.0%. The preparation steps of C/C and C/C-SiC composites are shown in Fig. 1.
2 Experimental
2.3 Brake tests
2.1 Raw materials
The tribological properties of the C/C mated with the C/C-SiC were tested on a MM-3 000 disk-on-disk type laboratory scale dynamometer (Shuntong, Xi’an, China). The tested two C/C specimens acted as a stator, and the rotor (i.e. counterpart) was a C/C-SiC sample. The dimensions of the stator discs were 17 mm in thickness, 37.5 mm and 26.5 mm in outer and inner radii, respectively. The outer and inner radii of the rotor discs were the same as the stator discs, but the thickness was 20 mm. There were three independent braking parameters-velocity, brake pressure and inertia, and the values were 6 500 r/min (21.7 m/s), 0.6 MPa, and 0.35 kg · m2, respectively. Braking speed, braking time and braking moment were recorded by a computer. The dimensional changes of the discs before and after the test were measured with a spiral micrometer, and the accuracy of the spiral micrometer is 0.01 mm. The average value at five random points was taken as the linear wear rate of the discs. Each brake test was repeated 10 times and the average COF was expressed as the results. The COF stability could be calculated from the following equation (1) [22]:
Polyacrylonitrile (PAN) based carbon fibers (T700，12 K) were supplied by the Toray, Japan. The Si powder (the diameters and purity are 30-50 μm and 99.0% respectively) was supplied by the Da Zelin-silicon Co., LTD, Beijing, China. 3D carbon fiber needled integrated felts were used as preforms for both C/C and C/C-SiC composites. 2.2 Sample preparation The carbon fiber preforms were prepared by a needling technique, starting from repeatedly overlapping the layers of short-cut fiber web and then needle-punching step by step. The density of the carbon fiber preforms was 0.45 ± 0.05 g/cm3. The carbon fiber preforms were transformed into C/C composites or C/C preforms by chemical vapor infiltration (CVI). The CVI was conducted at about 1000 °C for 300-500 h in an argon atmosphere of 0.1 MPa. C3H6 was employed as a carbon precursor and H2 as a carrier and diluting gas (C3H6/H2 =10mL•min − 1:20 mL•min −1). After CVI, two C/C composite samples were obtained. Following the CVI, one of the C/C composites was subjected to HTT at 2 300°C for 1-2 h in Ar. C/C-SiC composites were prepared by a liquid silicon infiltration method (LSI) to infiltrate Si into the C/C preforms. The LSI was carried out at 1 600±100 °C for 0.5-2.0 h under
S
μcp μ max
(1)
where μ cp is the average COF and μ max is the maximum COF.
Fig. 1 The preparation process of the C/C composites and the C/C-SiC composites.
Xi Ou-yang et al. / New Carbon Materials, 2019, 34(5): 472-481
2.4 Analysis methods The open porosities and bulk densities were measured by the Archimedes ’ method at room temperature, and the accuracy of the electronic balance was 0.1mg. The phase compositions, microstructures and morphologies and element distribution of C/C and C/C-SiC were investigated by a D/max2550VB + 18 kW rotating target X-ray diffraction analyzer (XRD, Rigaku Ltd., Japan, Cu Ka radiation), scanning electron microscopy (SEM, FEI Nova Nano SEM-230) and electron probe microanalysis (EPMA, JEOL CO., Jxa8230), respectively. The microstructures and morphologies of friction surfaces of the C/C and C/C-SiC were observed by an optical metallographic microscope (OM, LEICA DM 2700 M) and three-dimensional rotational digital video microscope (HIROX KH-7700). The wear debris was analyzed by XRD and SEM equipped with a Link energy-dispersive X-ray analyzer (EDX).
element distribution in the C/C-SiC (Fig. 3). Three different kinds of phases (light white phase, grey phase and dark phase) are observed in Fig. 3 (a). The light white phase is surrounded by the grey phase in the C/C-SiC. Based on the element distribution in Fig. 3 (b, c) and the XRD results in Fig. 2, it is found that the grey phase and the light white phase is SiC and residual Si, respectively while the remaining dark phase is C. During the LSI, the region within carbon fiber bundles is hardly accessible to the molten Si. However, the pores within both the inter-bundles and inter-plies regions could act as a communicating channel. Therefore, the SiC and the residual Si are mainly distributed between the inter-bundles and inter-plies regions, while the region intra-carbon fiber bundle is very little infiltrated by Si. The molten silicon reacts with matrix carbon and carbon fiber bundles to form SiC but there must be residual Si surrounding SiC in the specimen as the reaction of Si with C is not complete.
3 Results and discussion 3.1 Phase analysis The XRD pattern of the C/C-SiC is shown in Fig. 2. The phase analysis shows the presence of C, SiC and residual Si in the C/C-SiC. The carbon phase contains carbon fibers and pyrocarbon, and the structure of most SiC phase is a face-centered cubic (fcc; β) type. The contents of C, SiC and Si are 51.6, 38.7 and 9.7 wt.%, respectively by the K-value method. In order to fully understand the distribution of different phases in the C/C-SiC , EMPA is used to investigate the
Fig. 2 X-ray diffraction pattern of the C/C-SiC composite.
Fig. 3 Morphology and element distribution between carbon fiber bundles of C/C-SiC composite: (a) Morphology; (b) C; (c) Si.
Xi Ou-yang et al. / New Carbon Materials, 2019, 34(5): 472-481
Fig. 4 SEM images of C/C and C/C SiC before braking: (a) C/C A; (b) C/C B; (c) C/C SiC. Table 1 The g values and densities of C/C composites. Samples
d002 (nm)
g (%)
Density (g/cm3)
Open porosity (%)
C/C A
0.3435
5.8
1.68
8.1
C/C B
0.3397
50.0
1.72
14.4
3.2 Microstructure characterization Morphologies of the C/C and C/C-SiC are shown in Fig. 4. It can be seen from Fig. 4 (a, b) that only few micropores are distributed on the surface of the C/C A, but the number of micropores of the C/C B are more. Furthermore, it is obvious that the average size of micropores of the C/C A is smaller than that of the C/C B. After the HTT of the C/C, carbon with a turbostratic graphite structure is converted into the ideal three-dimensional ordered graphite crystal structure. In addition, the thermal expansion mismatch between the pyrolytic carbon and the carbon fibers during the HTT may cause interface cracking. The densities and opening porosities of the two specimens measured by the Archimedes method are shown in Table 1. The data also show that the densities and open porosities of the C/C increase with the degree of graphitization. The microstructures of the C/C have a great impact on friction and wear properties, such as wear rate. Fig. 4 (c) shows the microstructures of the C/C-SiC. Due to the infiltration of elemental silicon, most of the micropores in the C/C preform are filled with silicon to produce SiC and the residual silicon that does not react, so the densities of the C/C-SiC are relatively higher than that of the C/C. In other words, the degree of densification of the C/C-SiC is higher than that of the C/C. The results are in good agreement with EMPA, which indicates that the gray region is SiC and the white region is residual Si.
than that of the C/C A, indicating the crystallinity is higher in the C/C B. Table 1 shows the change of interlayer spacing (d002) and graphitization degrees of the two C/C. The interlayer spacing (d002) of the C/C A and C/C B is 0.3435 and 0.3397 nm, , respectively, corresponding to the graphitization degrees of 5.8 and 50.0%. The degree of graphitization degrees (g) is calculated by the following equation (2) [25]:
g = (0.3440 - d 002 )/(0.3440 - 0.3354)
(2)
where d002 is the interlayer distance determined by X-ray diffraction. The data reveal that the HTT results in a decrease in d002 and an improvement in the graphitization degree.
3.3 Graphitization degrees of the C/C composites Fig. 5 shows the XRD results of the C/C. The (002) peak of the C/C A is broad, indicating a low crystallinity. On the contrary, the (002) peak of the C/C B is higher and sharper
Fig. 5 X-ray diffraction pattern of the C/C composites.
Xi Ou-yang et al. / New Carbon Materials, 2019, 34(5): 472-481
The degree of graphitization is an important structural parameter, which influences the friction and wear properties of the C/C greatly. When the C/C is in a high temperature environment, the graphite lamellar structure is prone to relative slip under the action of thermal stress, which changes from the turbostratic graphite structure to the ideal three-dimensional ordered graphite crystal structure [26]. The orientation of the graphite lamellar structure tends to be parallel in the process of transformation, which means the degree of graphitization is improved. In the graphite crystal structure, the bonding force between the carbon atoms in a graphite layer is the covalent bond of a high bond energy, and the bonding force between the adjacent graphite layers is van der Waals force, which is a weak intermolecular force. During the brake test, due to the great pressure between the friction pairs, the graphite material between the layers prone to relative slip and easy to split along the layers, resulting in self-lubricating graphite microcrystalline fragments. These microcrystalline fragments in the friction surface will form a friction film and fill the friction surface defects, which flattens the uneven friction surface and improves the stability of the friction. 3.4 Tribological characteristics
spacing of the matrix carbon decreases, which means an increase of the graphitization degree. Some papers have reported [24, 27] that graphite is very soft, it is prone to deformation and crushed into friction film during the brake process, which will decrease the COF but increase the COF stability. The high COF stability helps to stabilize the brakes, but the low COF extends the braking time. Fig. 7 displays the average COF of a single braking versus the braking time and the average linear wear rate of the C/C mated with the C/C-SiC. As shown in Fig. 7 (a), the average COF of the C/C B mated with the C/C-SiC has a smaller fluctuation than that of the C/C A mated with the C/C-SiC. So the single average COF of the C/C B is more stable, which demonstrates that the braking stability of the C/C B for multiple braking is improved. In addition, as shown in Fig. 7 (b), the average linear wear rates of the C/C A and the C/C B are 3.7 μm/(slide cycle) and 8.9 μm/(slide cycle), respectively. The average linear wear rate of the C/C A is lower than that of the C/C B, which indicates that the wear resistance of the C/C B is declined because of the HTT. The average linear wear rate of the C/C-SiC mated with the C/C A is not much different from the C/C-SiC mated with the C/C B, which is 0.5 μm/(slide cycle) and 0.6 μm/(slide cycle), respectively.
3.4.1 Friction and wear properties The typical brake curves of the C/C mated with the C/C-SiC are shown in Fig. 6. It is noted that the brake curves exhibit a shape of horse-saddle. At the beginning of the brake test, the COF dramatically increases with time with a peak in the curve. After reaching the front peak, it decreases gradually and keeps a balance with varying degrees of fluctuations. At the end of the brake test, the COF increases again with a tail-up phenomenon in the curves. The brake curve of the C/C B mated with the C/C-SiC is flat, but the COF is lower than the C/C A. Shear forces on the crystallites and the nature of the lubricating films determine the friction and wear properties of the C/C composites [24]. As mentioned earlier, with increasing heat treatment temperature, the interlayer
Fig. 6 Typical brake curves of C/C mated with C/C-SiC.
Fig. 7 (a) The average COF of a single braking and (b) the linear wear rates of C/C and C/C-SiC.
Xi Ou-yang et al. / New Carbon Materials, 2019, 34(5): 472-481
3.4.2 Friction surface morphologies Friction and wear performances are not inherent properties of brake materials but a function of parameters of a brake system, and the tribological mechanism is very complicated[28]. The type of braking parameters, environmental and interface conditions and microstructures are important factors that determine the tribological properties. The friction surface morphologies are investigated by an OM as shown in Fig. 8, cracks and defects are observed on the friction surface of the C/C and C/C-SiC obviously. Frictional traction forces repeated for a large number of times on the friction surface in the course of braking, which leads to the generation of cracks. Under the cyclic stress, cracks are generated and propagated, resulting in the peeling of the matrix and fibers, which is the fatigue wear [28]. In other words, the presence of the fatigue wear causes cracks on the friction surface. As the cracks expand, they will lead to the defects and the high wear rate. At the starting of the friction experiment, asperities of the friction surface are mechanically interlocked with each other due to mutual embedding under the normal stress [29]. The asperities break off by the shearing and compressive force with the progress of the friction
experiment . The C/C-SiC consists of C, SiC and residual Si, with SiC and Si having a very high hardness. During the brake process, asperities with a high hardness plough on the friction surface and result in scratches, as displayed in Fig. 8. After the HTT of the C/C B, plenty of the structures of carbon change and the carbon with the ordered graphite crystal structure is easy to form friction films on the friction surfaces, which reduces the ploughing effect of the hard phases (Fig. 8 (b, d)). The weakening of the ploughing effect reduces the COF, which is well documented in Fig. 6. The flatness degree of friction surfaces of the C/C mated with the C/C-SiC is investigated by a 3D digital video microscope, which is shown in Fig. 9. It is apparent that the friction surface of the C/C B (Fig. 9 (b)) is more flat than the C/C A (Fig. 9 (a)), which indicates that the friction films of the C/C B are relatively uniform. Meanwhile, there are many asperities on the friction surface in Fig. 9 (a), which demonstrates the reason of fluctuation of the brake curve of the C/C A. In addition, it also can be found that the actual contact area of the friction surface of the C/C B (Fig. 9 (b)) is larger than that of the C/C A (Fig. 9 (a)), resulting in a higher wear rate of the C/C B.
Fig. 8 Optical photographs of friction surfaces of C/C mated with C/C-SiC: (a) C/C A; (b) C/C B; (c) C/C-SiC (mated with C/C A); (d) C/C-SiC (mated with C/C B).
Fig. 9 3D morphologies of friction surface of C/C composites: (a) C/C A; (b) C/C B.
Xi Ou-yang et al. / New Carbon Materials, 2019, 34(5): 472-481
Fig. 10 SEM morphologies and EDS spectra of wear debris of C/C mated with C/C-SiC: (a, b) C/C A mated with C/C-SiC; (c, d) C/C B mated with C/C-SiC.
3.4.3 Wear debris The SEM morphologies of the wear debris from the C/C mated with the C/C-SiC is investigated as shown in Fig. 10. It is observed that there are the particulate-type debris for the C/C A mated with the C/C-SiC (Fig. 10 (a)) and the sheet-type debris for the C/C B mated with the C/C-SiC (Fig. 10 (b)). Because of the C/C A without HTT, its carbon exists mainly in the turbostratic graphite structure. During the braking process, asperities break into small masses due to the shear stress and brake pressure, but the carbon with the turbostratic graphite structure is hard and it is difficult to form into a flake. What is more, the Si and SiC in the C/C-SiC are the hard phases, wear debris formed during the friction process mainly exists in small masses. Therefore, the observed wear debris of the C/C A is mainly irregular small particulates. After the HTT of the C/C B, a large amount of carbon with the turbostratic graphite structure is converted into the ideal three-dimensional ordered graphite crystal structure. The carbon with the ideal three-dimensional ordered graphite crystal structure is soft and it is easily pressed into a sheet shape under the braking pressure after falling off from the matrix. In addition, it can be found from the energy spectrum of the wear debris (Fig. 10 (b, d)) that the content of Si is small, indicating that the wear rate of the C/C-SiC is small. Fig. 11 shows the XRD patterns of the wear debris. It is seen that the C/C-SiC is composed of C, Si, and SiC (Fig. 2), while the wear debris contains not only C, Si, and SiC but also SiO2. It can be concluded that the Si or SiC in the C/C-SiC is
oxidized to form SiO2 during the braking process. Further, the oxidation of C can be inferred from the oxidation of Si or SiC. Because the temperature of the friction surface can reach 600 °C in each brake process, even the temperature of some local “hot spot” of the friction surface can reach over 800 °C [28, 30] . Whereas the anti-oxidation of carbon over 400 °C is poor, as well as silicon over 700 °C. Therefore, the oxidation of C can be inferred. The reactions of C, Si or SiC with oxygen are with the following equations:
C (s) ＋ O 2(g) → CO 2(g) Si (l) ＋ O 2(g) → SiO 2(l) SiC(l) ＋ O 2(g) → SiO 2 (l) ＋ CO 2(g)
(3) (4) (5)
However, CO2 exists in the gaseous state at high temperatures. Therefore, only SiO2 formed on the debris surface can be observed.
Fig. 11 XRD patterns of wear debris of C/C mated with C/C-SiC.
Xi Ou-yang et al. / New Carbon Materials, 2019, 34(5): 472-481
4 Conclusions
International, 2010, 43(11): 1951-1959.
The C/C specimens are prepared by CVI, and the C/C-SiC specimen is produced by the LSI method after CVI. The C/C samples are subjected to HTT at different temperatures. The tribological and wear properties of the C/C mated with the C/C-SiC are investigated. Main conclusions are summarized as follows. The densities and porosities of the C/C A without HTT are 1.68 g/cm3 and 8.1%, that of the C/C B with 2 300 °C HTT are 1.72 g/cm3 and 14.4%, while that of the C/C-SiC are 2.08 g/cm3 and 2.0%, respectively. The graphitization degrees of the C/C without and with 2300 °C HTT are 5.8 and 50.0%, respectively. Compared with the C/C A without HTT, the COF stability (0.65) of the C/C B with 2 300 °C HTT is improved, but the COF (0.280) is greatly attenuated. In addition, due to the softness of graphite carbon and high porosities, the wear rate of C/C B with 2 300 °C (8.9 μm/(slide cycle)) mated with the C/C-SiC is increased. The wear mechanisms of the C/C mated with the C/C-SiC are mainly grain-abrasion, oxidation-abrasion and fatigue wear. The grain-abrasive of the C/C B with 2 300 °C HTT is relatively mild, which is one of the reasons for the low stable COF.
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Volume 34, Issue 5, Oct 2019 Online English edition of the Chinese language journal Cite this article as: New Carbon Materials, 2019, 34(5): 472-481
RESEARCH PAPER
Effects of the high-temperature treatment of C/C composites on their tribological properties Xi Ou-yang1, Zhuan Li1,2*, Peng Xiao1,2, Guan-yi Chen1, jin-wei Li1, Peng-fei Liu1 China State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China;
1
Science and Technology on High Strength Structural Materials Laboratory, Central South University, Changsha 410083, China
2
Abstract:
C/C composites prepared by chemical vapor infiltration (CVI) were subjected to high temperature treatment (HTT) at 2300 °C under
an Ar atmosphere. C/C-SiC composites were prepared by liquid silicon infiltration (LSI) of the C/C composites before HTT. The C/C composites before (A) and after (B) HTT were mated with the C/C-SiC to form friction pairs to investigate their tribological properties. Results indicated that the average coefficient of friction (COF) and stable COF of the C/C-B are 0.280 and 0.65 while that of the C/C-B are 0.451 and 0.55, respectively. The average linear wear rate of the C/C-A and the C/C-B are 3.7 and 8.9 μm/(slide cycle), respectively. SEM reveals that the wear debris of the C/C-A is particulate while that of the C/C-B are the sheets, and the friction surface is more flat for C/C-B than C/C-A. HTT leads to a softening of the carbon, increases the wear rate, but improves the stable COF of the C/C composites. The wear mechanisms of the composites are mainly grain-abrasion, oxidation loss and fatigue wear. Key Words: C/C; C/C-SiC; High-temperature treatment; Tribological properties
1 Introduction With the rapid development of aerospace, military, automobile and high-speed train industry, the speed becomes more and more quick and the load is also heavier than past, which place high demands on the brake materials. Furthermore, brake materials must be better adapted to harsh environments such as sand and rainwater. C/C composites have become one of the excellent choices as brake materials owing to their light weight, excellent thermal and mechanical properties, and high energy absorption [1]. C/C composites were first investigated in 1958 [2, 3] with the purpose to compensate the weaknesses of powder metallurgical materials [4-7], and then aroused the interest of the majority of scientific research workers. Y. Cai et al. [8] have studied the microstructures and mechanical properties of three-dimensional ceramic filler-modified C/C composites. H. Kasem et al. [9] have researched the characterization of surface grooves and scratches of C/C composites induced by friction at low and high temperatures. S. Sarkar et al. [6] have reported the impact abrasive wear response of C/C composites at elevated temperatures. Because of their low-density and excellent high-temperature properties, C/C composites have been introduced in the aircraft brake system [10, 11]. However, C/C composites were commonly considered to have some fatal weakness such as poor oxidation resistance under ambient environment, very low coefficient of friction (COF) in a wet environment and relatively high wear rates. In response to these problems, researchers have done a lot of
improvement work [12-14]. C/C-SiC composites display the theoretically excellent high-temperature mechanical and thermal properties, and now they are widely used in sports cars, luxury sedans and airplane [15-17] . In addition, C/C-SiC composites containing at least 20 wt% of SiC also have many other advantages such as the lower wear rate, longer service life, especially lower sensibility to surroundings and temperature as compared with C/C composites [18, 19]. Since 1990s, W. Krenkel et al. [15, 20, 21] have made a systematic study on C/C-SiC composites in the field of tribology, such as C/C-SiC composites for advanced friction systems, integration of ceramic matrix composite （ CMC ） -brake disks in automotive brake systems and full-ceramic brake systems for high performance friction applications. In order to optimize the properties of C/C-SiC composites, Z. Li et al. [22, 23] have done some modified work. Though C/C-SiC composites have overcome many disadvantages of C/C composites, they also have shortcomings. On the one hand, the COF of self-mated C/C-SiC is too high, which is easy to cause the lock and lead to traffic accidents. On the other hand, C/C-SiC composites show a higher density compared with C/C composites. In this paper, in order to overcome the disadvantages of C/C and C/C-SiC composite brakes, the idea of using C/C mated with C/C-SiC for brake assembly was proposed. Meanwhile, to further enhance the COF stability, one of the
Received date: 10 Aug 2019; Revised date: 30 Sep 2019 *Corresponding author. E-mail: [email protected] Copyright©2019, Institute of Coal Chemistry, Chinese Academy of Sciences. Published by Elsevier Limited. All rights reserved. DOI: 10.1016/S1872-5805(19)60025-9
Xi Ou-yang et al. / New Carbon Materials, 2019, 34(5): 472-481
C/C composites was subjected to high temperature treatment (HTT) at 2 300 °C. R. Luo et al. [24] have reported that the stability of the braking moment-time curves of C/C composites increased with the heat treatment temperature. The correlation of the phase composition, microstructure, friction surfaces to the friction and wear performance was investigated. The characteristics of friction surfaces and wear debris were evaluated and the friction mechanism was proposed.
vacuum. The liquid Si reacted with a small amount of the C/C and formed SiC ceramic matrix. The C/C preform specimen without HTT was named as C/C A, and another one with 2300 °C HTT was named C/C B. The densities and porosities of the C/C A were 1.68 g/cm3 and 8.1%, respectively, those of the C/C B were 1.72 g/cm3 and 14.4% while those of the C/C-SiC were 2.08 g/cm3 and 2.0%. The preparation steps of C/C and C/C-SiC composites are shown in Fig. 1.
2 Experimental
2.3 Brake tests
2.1 Raw materials
The tribological properties of the C/C mated with the C/C-SiC were tested on a MM-3 000 disk-on-disk type laboratory scale dynamometer (Shuntong, Xi’an, China). The tested two C/C specimens acted as a stator, and the rotor (i.e. counterpart) was a C/C-SiC sample. The dimensions of the stator discs were 17 mm in thickness, 37.5 mm and 26.5 mm in outer and inner radii, respectively. The outer and inner radii of the rotor discs were the same as the stator discs, but the thickness was 20 mm. There were three independent braking parameters-velocity, brake pressure and inertia, and the values were 6 500 r/min (21.7 m/s), 0.6 MPa, and 0.35 kg · m2, respectively. Braking speed, braking time and braking moment were recorded by a computer. The dimensional changes of the discs before and after the test were measured with a spiral micrometer, and the accuracy of the spiral micrometer is 0.01 mm. The average value at five random points was taken as the linear wear rate of the discs. Each brake test was repeated 10 times and the average COF was expressed as the results. The COF stability could be calculated from the following equation (1) [22]:
Polyacrylonitrile (PAN) based carbon fibers (T700，12 K) were supplied by the Toray, Japan. The Si powder (the diameters and purity are 30-50 μm and 99.0% respectively) was supplied by the Da Zelin-silicon Co., LTD, Beijing, China. 3D carbon fiber needled integrated felts were used as preforms for both C/C and C/C-SiC composites. 2.2 Sample preparation The carbon fiber preforms were prepared by a needling technique, starting from repeatedly overlapping the layers of short-cut fiber web and then needle-punching step by step. The density of the carbon fiber preforms was 0.45 ± 0.05 g/cm3. The carbon fiber preforms were transformed into C/C composites or C/C preforms by chemical vapor infiltration (CVI). The CVI was conducted at about 1000 °C for 300-500 h in an argon atmosphere of 0.1 MPa. C3H6 was employed as a carbon precursor and H2 as a carrier and diluting gas (C3H6/H2 =10mL•min − 1:20 mL•min −1). After CVI, two C/C composite samples were obtained. Following the CVI, one of the C/C composites was subjected to HTT at 2 300°C for 1-2 h in Ar. C/C-SiC composites were prepared by a liquid silicon infiltration method (LSI) to infiltrate Si into the C/C preforms. The LSI was carried out at 1 600±100 °C for 0.5-2.0 h under
S
μcp μ max
(1)
where μ cp is the average COF and μ max is the maximum COF.
Fig. 1 The preparation process of the C/C composites and the C/C-SiC composites.
Xi Ou-yang et al. / New Carbon Materials, 2019, 34(5): 472-481
2.4 Analysis methods The open porosities and bulk densities were measured by the Archimedes ’ method at room temperature, and the accuracy of the electronic balance was 0.1mg. The phase compositions, microstructures and morphologies and element distribution of C/C and C/C-SiC were investigated by a D/max2550VB + 18 kW rotating target X-ray diffraction analyzer (XRD, Rigaku Ltd., Japan, Cu Ka radiation), scanning electron microscopy (SEM, FEI Nova Nano SEM-230) and electron probe microanalysis (EPMA, JEOL CO., Jxa8230), respectively. The microstructures and morphologies of friction surfaces of the C/C and C/C-SiC were observed by an optical metallographic microscope (OM, LEICA DM 2700 M) and three-dimensional rotational digital video microscope (HIROX KH-7700). The wear debris was analyzed by XRD and SEM equipped with a Link energy-dispersive X-ray analyzer (EDX).
element distribution in the C/C-SiC (Fig. 3). Three different kinds of phases (light white phase, grey phase and dark phase) are observed in Fig. 3 (a). The light white phase is surrounded by the grey phase in the C/C-SiC. Based on the element distribution in Fig. 3 (b, c) and the XRD results in Fig. 2, it is found that the grey phase and the light white phase is SiC and residual Si, respectively while the remaining dark phase is C. During the LSI, the region within carbon fiber bundles is hardly accessible to the molten Si. However, the pores within both the inter-bundles and inter-plies regions could act as a communicating channel. Therefore, the SiC and the residual Si are mainly distributed between the inter-bundles and inter-plies regions, while the region intra-carbon fiber bundle is very little infiltrated by Si. The molten silicon reacts with matrix carbon and carbon fiber bundles to form SiC but there must be residual Si surrounding SiC in the specimen as the reaction of Si with C is not complete.
3 Results and discussion 3.1 Phase analysis The XRD pattern of the C/C-SiC is shown in Fig. 2. The phase analysis shows the presence of C, SiC and residual Si in the C/C-SiC. The carbon phase contains carbon fibers and pyrocarbon, and the structure of most SiC phase is a face-centered cubic (fcc; β) type. The contents of C, SiC and Si are 51.6, 38.7 and 9.7 wt.%, respectively by the K-value method. In order to fully understand the distribution of different phases in the C/C-SiC , EMPA is used to investigate the
Fig. 2 X-ray diffraction pattern of the C/C-SiC composite.
Fig. 3 Morphology and element distribution between carbon fiber bundles of C/C-SiC composite: (a) Morphology; (b) C; (c) Si.
Xi Ou-yang et al. / New Carbon Materials, 2019, 34(5): 472-481
Fig. 4 SEM images of C/C and C/C SiC before braking: (a) C/C A; (b) C/C B; (c) C/C SiC. Table 1 The g values and densities of C/C composites. Samples
d002 (nm)
g (%)
Density (g/cm3)
Open porosity (%)
C/C A
0.3435
5.8
1.68
8.1
C/C B
0.3397
50.0
1.72
14.4
3.2 Microstructure characterization Morphologies of the C/C and C/C-SiC are shown in Fig. 4. It can be seen from Fig. 4 (a, b) that only few micropores are distributed on the surface of the C/C A, but the number of micropores of the C/C B are more. Furthermore, it is obvious that the average size of micropores of the C/C A is smaller than that of the C/C B. After the HTT of the C/C, carbon with a turbostratic graphite structure is converted into the ideal three-dimensional ordered graphite crystal structure. In addition, the thermal expansion mismatch between the pyrolytic carbon and the carbon fibers during the HTT may cause interface cracking. The densities and opening porosities of the two specimens measured by the Archimedes method are shown in Table 1. The data also show that the densities and open porosities of the C/C increase with the degree of graphitization. The microstructures of the C/C have a great impact on friction and wear properties, such as wear rate. Fig. 4 (c) shows the microstructures of the C/C-SiC. Due to the infiltration of elemental silicon, most of the micropores in the C/C preform are filled with silicon to produce SiC and the residual silicon that does not react, so the densities of the C/C-SiC are relatively higher than that of the C/C. In other words, the degree of densification of the C/C-SiC is higher than that of the C/C. The results are in good agreement with EMPA, which indicates that the gray region is SiC and the white region is residual Si.
than that of the C/C A, indicating the crystallinity is higher in the C/C B. Table 1 shows the change of interlayer spacing (d002) and graphitization degrees of the two C/C. The interlayer spacing (d002) of the C/C A and C/C B is 0.3435 and 0.3397 nm, , respectively, corresponding to the graphitization degrees of 5.8 and 50.0%. The degree of graphitization degrees (g) is calculated by the following equation (2) [25]:
g = (0.3440 - d 002 )/(0.3440 - 0.3354)
(2)
where d002 is the interlayer distance determined by X-ray diffraction. The data reveal that the HTT results in a decrease in d002 and an improvement in the graphitization degree.
3.3 Graphitization degrees of the C/C composites Fig. 5 shows the XRD results of the C/C. The (002) peak of the C/C A is broad, indicating a low crystallinity. On the contrary, the (002) peak of the C/C B is higher and sharper
Fig. 5 X-ray diffraction pattern of the C/C composites.
Xi Ou-yang et al. / New Carbon Materials, 2019, 34(5): 472-481
The degree of graphitization is an important structural parameter, which influences the friction and wear properties of the C/C greatly. When the C/C is in a high temperature environment, the graphite lamellar structure is prone to relative slip under the action of thermal stress, which changes from the turbostratic graphite structure to the ideal three-dimensional ordered graphite crystal structure [26]. The orientation of the graphite lamellar structure tends to be parallel in the process of transformation, which means the degree of graphitization is improved. In the graphite crystal structure, the bonding force between the carbon atoms in a graphite layer is the covalent bond of a high bond energy, and the bonding force between the adjacent graphite layers is van der Waals force, which is a weak intermolecular force. During the brake test, due to the great pressure between the friction pairs, the graphite material between the layers prone to relative slip and easy to split along the layers, resulting in self-lubricating graphite microcrystalline fragments. These microcrystalline fragments in the friction surface will form a friction film and fill the friction surface defects, which flattens the uneven friction surface and improves the stability of the friction. 3.4 Tribological characteristics
spacing of the matrix carbon decreases, which means an increase of the graphitization degree. Some papers have reported [24, 27] that graphite is very soft, it is prone to deformation and crushed into friction film during the brake process, which will decrease the COF but increase the COF stability. The high COF stability helps to stabilize the brakes, but the low COF extends the braking time. Fig. 7 displays the average COF of a single braking versus the braking time and the average linear wear rate of the C/C mated with the C/C-SiC. As shown in Fig. 7 (a), the average COF of the C/C B mated with the C/C-SiC has a smaller fluctuation than that of the C/C A mated with the C/C-SiC. So the single average COF of the C/C B is more stable, which demonstrates that the braking stability of the C/C B for multiple braking is improved. In addition, as shown in Fig. 7 (b), the average linear wear rates of the C/C A and the C/C B are 3.7 μm/(slide cycle) and 8.9 μm/(slide cycle), respectively. The average linear wear rate of the C/C A is lower than that of the C/C B, which indicates that the wear resistance of the C/C B is declined because of the HTT. The average linear wear rate of the C/C-SiC mated with the C/C A is not much different from the C/C-SiC mated with the C/C B, which is 0.5 μm/(slide cycle) and 0.6 μm/(slide cycle), respectively.
3.4.1 Friction and wear properties The typical brake curves of the C/C mated with the C/C-SiC are shown in Fig. 6. It is noted that the brake curves exhibit a shape of horse-saddle. At the beginning of the brake test, the COF dramatically increases with time with a peak in the curve. After reaching the front peak, it decreases gradually and keeps a balance with varying degrees of fluctuations. At the end of the brake test, the COF increases again with a tail-up phenomenon in the curves. The brake curve of the C/C B mated with the C/C-SiC is flat, but the COF is lower than the C/C A. Shear forces on the crystallites and the nature of the lubricating films determine the friction and wear properties of the C/C composites [24]. As mentioned earlier, with increasing heat treatment temperature, the interlayer
Fig. 6 Typical brake curves of C/C mated with C/C-SiC.
Fig. 7 (a) The average COF of a single braking and (b) the linear wear rates of C/C and C/C-SiC.
Xi Ou-yang et al. / New Carbon Materials, 2019, 34(5): 472-481
3.4.2 Friction surface morphologies Friction and wear performances are not inherent properties of brake materials but a function of parameters of a brake system, and the tribological mechanism is very complicated[28]. The type of braking parameters, environmental and interface conditions and microstructures are important factors that determine the tribological properties. The friction surface morphologies are investigated by an OM as shown in Fig. 8, cracks and defects are observed on the friction surface of the C/C and C/C-SiC obviously. Frictional traction forces repeated for a large number of times on the friction surface in the course of braking, which leads to the generation of cracks. Under the cyclic stress, cracks are generated and propagated, resulting in the peeling of the matrix and fibers, which is the fatigue wear [28]. In other words, the presence of the fatigue wear causes cracks on the friction surface. As the cracks expand, they will lead to the defects and the high wear rate. At the starting of the friction experiment, asperities of the friction surface are mechanically interlocked with each other due to mutual embedding under the normal stress [29]. The asperities break off by the shearing and compressive force with the progress of the friction
experiment . The C/C-SiC consists of C, SiC and residual Si, with SiC and Si having a very high hardness. During the brake process, asperities with a high hardness plough on the friction surface and result in scratches, as displayed in Fig. 8. After the HTT of the C/C B, plenty of the structures of carbon change and the carbon with the ordered graphite crystal structure is easy to form friction films on the friction surfaces, which reduces the ploughing effect of the hard phases (Fig. 8 (b, d)). The weakening of the ploughing effect reduces the COF, which is well documented in Fig. 6. The flatness degree of friction surfaces of the C/C mated with the C/C-SiC is investigated by a 3D digital video microscope, which is shown in Fig. 9. It is apparent that the friction surface of the C/C B (Fig. 9 (b)) is more flat than the C/C A (Fig. 9 (a)), which indicates that the friction films of the C/C B are relatively uniform. Meanwhile, there are many asperities on the friction surface in Fig. 9 (a), which demonstrates the reason of fluctuation of the brake curve of the C/C A. In addition, it also can be found that the actual contact area of the friction surface of the C/C B (Fig. 9 (b)) is larger than that of the C/C A (Fig. 9 (a)), resulting in a higher wear rate of the C/C B.
Fig. 8 Optical photographs of friction surfaces of C/C mated with C/C-SiC: (a) C/C A; (b) C/C B; (c) C/C-SiC (mated with C/C A); (d) C/C-SiC (mated with C/C B).
Fig. 9 3D morphologies of friction surface of C/C composites: (a) C/C A; (b) C/C B.
Xi Ou-yang et al. / New Carbon Materials, 2019, 34(5): 472-481
Fig. 10 SEM morphologies and EDS spectra of wear debris of C/C mated with C/C-SiC: (a, b) C/C A mated with C/C-SiC; (c, d) C/C B mated with C/C-SiC.
3.4.3 Wear debris The SEM morphologies of the wear debris from the C/C mated with the C/C-SiC is investigated as shown in Fig. 10. It is observed that there are the particulate-type debris for the C/C A mated with the C/C-SiC (Fig. 10 (a)) and the sheet-type debris for the C/C B mated with the C/C-SiC (Fig. 10 (b)). Because of the C/C A without HTT, its carbon exists mainly in the turbostratic graphite structure. During the braking process, asperities break into small masses due to the shear stress and brake pressure, but the carbon with the turbostratic graphite structure is hard and it is difficult to form into a flake. What is more, the Si and SiC in the C/C-SiC are the hard phases, wear debris formed during the friction process mainly exists in small masses. Therefore, the observed wear debris of the C/C A is mainly irregular small particulates. After the HTT of the C/C B, a large amount of carbon with the turbostratic graphite structure is converted into the ideal three-dimensional ordered graphite crystal structure. The carbon with the ideal three-dimensional ordered graphite crystal structure is soft and it is easily pressed into a sheet shape under the braking pressure after falling off from the matrix. In addition, it can be found from the energy spectrum of the wear debris (Fig. 10 (b, d)) that the content of Si is small, indicating that the wear rate of the C/C-SiC is small. Fig. 11 shows the XRD patterns of the wear debris. It is seen that the C/C-SiC is composed of C, Si, and SiC (Fig. 2), while the wear debris contains not only C, Si, and SiC but also SiO2. It can be concluded that the Si or SiC in the C/C-SiC is
oxidized to form SiO2 during the braking process. Further, the oxidation of C can be inferred from the oxidation of Si or SiC. Because the temperature of the friction surface can reach 600 °C in each brake process, even the temperature of some local “hot spot” of the friction surface can reach over 800 °C [28, 30] . Whereas the anti-oxidation of carbon over 400 °C is poor, as well as silicon over 700 °C. Therefore, the oxidation of C can be inferred. The reactions of C, Si or SiC with oxygen are with the following equations:
C (s) ＋ O 2(g) → CO 2(g) Si (l) ＋ O 2(g) → SiO 2(l) SiC(l) ＋ O 2(g) → SiO 2 (l) ＋ CO 2(g)
(3) (4) (5)
However, CO2 exists in the gaseous state at high temperatures. Therefore, only SiO2 formed on the debris surface can be observed.
Fig. 11 XRD patterns of wear debris of C/C mated with C/C-SiC.
Xi Ou-yang et al. / New Carbon Materials, 2019, 34(5): 472-481
4 Conclusions
International, 2010, 43(11): 1951-1959.
The C/C specimens are prepared by CVI, and the C/C-SiC specimen is produced by the LSI method after CVI. The C/C samples are subjected to HTT at different temperatures. The tribological and wear properties of the C/C mated with the C/C-SiC are investigated. Main conclusions are summarized as follows. The densities and porosities of the C/C A without HTT are 1.68 g/cm3 and 8.1%, that of the C/C B with 2 300 °C HTT are 1.72 g/cm3 and 14.4%, while that of the C/C-SiC are 2.08 g/cm3 and 2.0%, respectively. The graphitization degrees of the C/C without and with 2300 °C HTT are 5.8 and 50.0%, respectively. Compared with the C/C A without HTT, the COF stability (0.65) of the C/C B with 2 300 °C HTT is improved, but the COF (0.280) is greatly attenuated. In addition, due to the softness of graphite carbon and high porosities, the wear rate of C/C B with 2 300 °C (8.9 μm/(slide cycle)) mated with the C/C-SiC is increased. The wear mechanisms of the C/C mated with the C/C-SiC are mainly grain-abrasion, oxidation-abrasion and fatigue wear. The grain-abrasive of the C/C B with 2 300 °C HTT is relatively mild, which is one of the reasons for the low stable COF.
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