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Influence of Cu on the mechanical and tribological properties of Ti3SiC2
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Influence of Cu on the mechanical and tribological properties of Ti3SiC2 Wentao Dang a,b, Shufang Ren a,n, Jiansong Zhou a,n, Youjun Yu a, Zhen Li a, Lingqian Wang a a b
State Key Laboratory of Solid Lubrication, Lanzhou Institute of Chemical Physics, Chinese Academy of Sciences, Lanzhou 730000, PR China Graduate School of Chinese Academy of Sciences, Beijing 100039, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 6 February 2016 Received in revised form 1 March 2016 Accepted 12 March 2016
Ti3SiC2/Cu composites with different contents of Cu were fabricated by mechanical alloying and spark plasma sintering method. The phase composition and structure of the composites were analyzed by X-ray diffractometry and scanning electron microscopy equipped with energy dispersive spectroscopy. The mechanical and tribological properties of Ti3SiC2/Cu composites were tested and analyzed compared with monolithic Ti3SiC2 in details. The results show that the Cu leads to the decomposition of Ti3SiC2 to produce TiCx, Ti5Si3Cy, Cu3Si, and TiSi2Cz. The friction coefficient and wear rate of the composites are lower than that of monolithic Ti3SiC2, which is ascribed to the fixing effect of hard TiCx, Ti5Si3Cy, and Cu3Si to inhibit the abrasive friction and wear. However, at elevated temperatures (ranging from room temperature to 600 °C) the friction and wear of the composites are higher than those at room temperature. Plastic flowing and tribo-oxidation wear accompanied by material transference contribute to the increased friction and wear at elevated temperatures. & 2016 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Keywords: C. Mechanical properties Ti3SiC2/Cu composites Tribological property Abrasive wear Tribo-oxidation wear SPS
1. Introduction Ti3SiC2 has recently attracted extensive research attention due to its combination of the merits of both metals and ceramics. The unique properties of Ti3SiC2 are derived from the unusual valence bond structure that is there are covalent, ionic, and metal three bonds in its crystalline structure [1–3]. Though Ti3SiC2 possesses layered structure it is not intrinsically self-lubricating for the polycrystalline Ti3SiC2. In fact, many studies [4–7] have found that the three-body abrasive wear that derives from the fractured and pulled-out Ti3SiC2 grains is the main wear mechanism at room temperature, which leads to the high friction and wear. The fundamental reason why Ti3SiC2 grains are inclined to fracture and pull out is that the grain boundary force of Ti3SiC2 grains is weak [8]. This intrinsic structure determines its some apparent defective performances, such as the low hardness (4–6 GPa), moderate flexural strength (200–600 MPa), and the brittleness like ceramics at 25–1100 °C [2,3]. These characteristics limit the wide applications of Ti3SiC2 as structural and functional material. Many researchers focus on the strengthening of Ti3SiC2 by the hard phases such as Al2O3 [9], TiC [10], Ti5Si3 [11], and WC-Co etc [12]. The mechanical and tribological properties of Ti3SiC2/ceramic composites are better than that of monolithic Ti3SiC2. Nevertheless, as for Ti3SiC2/metal composite, the corresponding research is very limited. To the knowledge of authors, Ti3SiC2 is chemically n
Corresponding authors. E-mail addresses: [email protected] (S. Ren), [email protected] (J. Zhou).
reactive when sintering with metals at high temperature. Li et al. [13] synthesized Ti3SiC2/Ni and Ti3SiC2/Co through vacuum sintering and they found the metals aggregated towards the sample surfaces resulting in layered microstructural features because the wettability between Ti3SiC2 and Ni or Co are both poor. They suggested some adequate sintering should be added to improve the wetting performance and interface properties between Ti3SiC2 and metals to achieve homogeneous microstructures and excellent mechanical properties. Many researchers reported the reactivity between Ti3SiC2 and metals. Lu et al. [14] found chemical reaction between Cu and Ti3SiC2 contributes to the wettability. Zhou and his coworkers [15] investigated chemical reactions and stability of Ti3SiC2 in Cu during high-temperature processing of Cu/Ti3SiC2 composites. The results indicates that Ti3SiC2 can react with Cu above 900 °C. When the low content of Ti3SiC2 or below 1000 °C, Cu(Si) solid solution and TiCx were formed, whereas at high temperature or high content of Ti3SiC2, Cu–Si intermetallic compounds like Cu5Si, Cu15Si4, and TiCx were formed. So it is critical to control the content of Cu to prevent total Ti3SiC2 from reacting with Cu. Gu et al. [16] investigated the chemical reactions between Ti and Ti3SiC2 during the preparation of Ti–Ti3SiC2 composites in the temperature range of 1000–1300 °C. The results demonstrate that Ti3SiC2 reacts with Ti to form TiCx, Ti5Si3, and TiSi2. They proposed that it is difficult to synthesize composite with titanium and Ti3SiC2 because the reactions between Ti and Ti3SiC2 occur at 1000 °C, which is lower than the melting point of Ti (1670 °C). Copper is a soft metal with good ductility, thermal and electrical conductivity, and excellent mechanical properties, which is commonly used as the cable, electrical and electronic components,
http://dx.doi.org/10.1016/j.ceramint.2016.03.099 0272-8842/& 2016 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
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and building materials. Ti3SiC2 is a potential structure/functional material with many unique properties such as high modulus and strength, high electrical conductivity, good thermal conductivity, good thermal stability, and excellent high-temperature oxidation resistance and so on. However, the poor tribological properties that are traced to the weak intergranular bonding strength limit its application as sliding components that serves in sliding friction conditions. Given the good ductility, electrical conductivity, and the excellent mechanical properties, in the present paper, copper was selected to strengthen the intergranular bonding strength of Ti3SiC2 grains and to then improve the tribological properties, low hardness, and the room-temperature brittleness for Ti3SiC2. The aim of this study is to synthesize Ti3SiC2/Cu composite and investigate the phase composition, mechanical and tribological properties. We hope the mechanical properties and the wear resistance of monolithic Ti3SiC2 can be improved by adding Cu element into Ti3SiC2.
be referred to in two ways: μm to refer to the mean friction coefficient over the entire sliding distance, and μr to refer to its real-time value as a function of sliding time. The μrs were recorded automatically by the tribometer. The wear rates (WR) of the Ti3SiC2/Cu samples were calculated by the following formula [17]:
W = V /PS
(1)
where W is the wear rate, V the wear volume (mm3), P the applied load (N), and S the total sliding distance (m). The profile of cross section of worn surface was outlined using NaNomap 500LS scanning three-dimension surface profiler to determine the wear volume as
(2)
V = AL 3
where V is the wear volume (mm ), A the cross section area of the worn surface (mm), and L the perimeter of wear track (mm). 2.4. Characterization
2. Material and methods 2.1. Samples preparation The commercially available Ti3SiC2 (YueHuan New Material Technology Co., Ltd. 99 wt%, 200 mesh) powder and Cu (Sinopharm Chemical Reagent Co., Ltd, Shanghai, China, 99%, 200 mesh) powder were used in this study. In order to determine the optimum Cu content in the Ti3SiC2/Cu composites, samples with different mass fractions of Cu (5, 10, 15, 20 wt%) were prepared by mechanical alloying (MA) and the following accompanied spark plasma sintering (SPS). The mechanical alloying was carried out by high energy ball-milling machine (Pulverisettes 4, Germany) with a speed of 300 rpm and a ratio of ball to powder of 12 for 6 h. The mixed powder was sintered at 1050 °C for 5 min with the heating rate of 60 °C/min at a cylindrical graphite mold with a 25 mm diameter in vacuum ( o10 Pa) under a pressure of 30 MPa by SPS (SPS-20T-10, Shanghai Chen Hua Technology Co., Ltd.). 2.2. Mechanical properties The Vickers hardness was measured by the MH-5-VM microhardness tester under applied 10 N load with the dwell time of 5 s. The hardness was calculated by averaging at least 8 measurements. Flexural strength and compression strength were measured using a universal material tester (SANS-CMT5205, MTS, China). The three-point bending test was conducted to obtain the flexural strength. The sample was cut to be 3 4 20 mm3. The cross head speed was 0.5 mm/min and the support span was 16 mm. The sizes of samples for compressive strength test were Φ5 mm 12.5 mm and the cross head speed was 0.2 mm/min. 2.3. Friction and wear test The friction and wear tests were performed using a high temperature ball-on-disc tribometer (HT-1000, Zhong Ke Kai Te Ltd. China) capable of going up to 900 °C. All the tests were done at a linear velocity of 0.2 m/s and a load of 5 N corresponding to a stress of ∼0.1 MPa under a non-lubrication condition with a rotation radius of 5 mm. The sliding distance for all the friction and wear tests was 360 m. The Ti3SiC2/Cu composites were shaped into cylindrical having a diameter of 25 mm working face and ∼5 mm thick. The counter surfaces were SiC ceramic ball with a diameter of 6.43 mm (density of 3.2 g/cm3 and a hardness of 28 GPa, K&H, Japan). All surfaces were polished to a 1 μm diamond finish, washed with ethanol and dried prior to testing. The measured μ's will
The phase compositions of the composites were determined by D/MAX-2000 X-ray diffractometer (XRD) with Cukɑ radiation in wavelength λ ¼ 0.1540598 nm. The microstructures of the specimens were examined by scanning electron microscopy (SEM, JEOL-5600LV) equipped with energy dispersive spectroscopy (EDS). The chemical states of typical elements on the worn surfaces were determined using a PHI-5702 X-ray photoelectron spectroscope (XPS) with a resolution of 0.2 eV using Al Ka radiation as the excitation source and the binding energy of carbon (C1s: 284.8 eV) as the reference.
3. Results and discussion 3.1. Phase composition and microstructure Fig. 1a and b shows the SEM micrographs of Ti3SiC2 powder and Ti3SiC2/10 wt% Cu (noted as TSC10) mixed powder after MA, respectively. It can be observed that the grain size of Ti3SiC2 particles ranges from 2 to 12 μm and the grain size of mixed powder of Ti3SiC2 and 10 wt% Cu becomes smaller than that of Ti3SiC2, ranging from 1 to 6 um after MA. The XRD patterns for Ti3SiC2 powder and Ti3SiC2 with 10 wt% Cu after MA are shown in Fig. 2. As for the MA mixed powder the diffraction peaks of Ti3SiC2 obviously broaden, which means Ti3SiC2 grains were refined during highenergy ball milling. Besides the peaks of Ti3SiC2 and Cu, the diffraction peaks of TiCx and Cu(Si) solid solution appear, which indicates that Cu reacts with Ti3SiC2 producing TiCx and Cu(Si) solid solution during high-energy ball milling. Mechanical alloying (MA) is a solid-state powder processing technique involving repeated welding, fracturing, and rewelding of powder particles in a highenergy ball mill [18]. MA has now been shown to be capable of synthesizing a variety of non-equilibrium alloy phases, such as supersaturated solid solutions, intermetallics, nanostructures, and amorphous alloys and so on starting from blended elementals powders. Recent mechanochemical synthesis of materials have been critically reviewed after discussing the process and process variables involved in MA [19]. So in the present study solid solutions of Cu(Si), mechanochemical product of TiCx derived from the decomposition of Ti3SiC2 were found besides the powders refining phenomenon. Additionally, according to our experience of preparation Ti3SiC2/Cu the MA processing can help to improve the density and the purity of the composites. Fig. 3 shows the XRD patterns of Ti3SiC2/Cu composites with different Cu contents sintered by SPS. It can be seen that the diffraction peaks of elementary Cu does not emerge in all the
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Fig. 1. SEM micrographs of (a) Ti3SiC2 powder and (b) TSC10 mixed powder after MA.
Fig. 2. XRD patterns for (a) Ti3SiC2 powder and (b) Ti3SiC2 with 10 wt% Cu powder after MA.
composites owing to the reaction of Cu with Ti3SiC2 to produce TiCx, Ti5Si3Cy, and Cu3Si during sintering. The intensity of the diffraction peaks of Ti3SiC2 weakens with the content of Cu increasing in the composites. The major ingredient of Ti3SiC2/5 wt% Cu (TSC5) and TSC10 composites is Ti3SiC2 with a small amount of Ti5Si3Cy, TiCx, and Cu3Si. The relative content of Cu3Si in TSC10 composite is higher than that of TSC5 by semiquantitative analysis. With the increase of Cu, the diffraction peaks of Ti3SiC2 matrix almost disappear and the major phase composition becomes TiCx, Ti5Si3Cy, TiSi2Cz, and Cu3Si (as shown in Fig. 3d and e). The detail products are listed in Table 1. In conclusion, with the content increase of Cu, the chemical reaction between Ti3SiC2 matrix and Cu takes place more and more intensively during the sintering by SPS. Fig. 4 shows the backscattered electron topography of sintered TSC10 composite. In the image we can see that the white punctate and linear regions distribute between the grey regions. According to the analysis of corresponding EDS, the white region abounds in Cu element and lacks Ti element while the grey region abounds in Ti element and lacks Cu element. So it can be referred that the white regions are mainly composed of Cu3Si and the grey regions are consist of the major Ti3SiC2 and the minor TiCx and Ti5Si3Cy. Cu3Si phase distributes around the Ti3SiC2 particles. From above results of phase and microstructure of the composites it can be concluded that the Cu element reacts with Ti3SiC2 matrix during the process of sintering. The reactive products vary
when different contents of Cu were added. The monolithic Ti3SiC2 decomposed to produce minor Ti5Si3Cy during sintering at 1050 °C with the soaking time of 5 min by SPS. For the 5 and 10 wt% Cu, the main phase for the composites is Ti3SiC2 with minor TiCx, Ti5Si3Cy, and Cu3Si. When the content of Cu is above 10 wt%, the Ti3SiC2 matrix completely decomposes to produce major TiCx and minor Ti5Si3Cy, TiSi2, and Cu3Si. Zhou et al. [15] investigated the chemical reactions of Ti3SiC2 with 20 to 70 vol% in Cu matrix. They found that Cu reacts with Ti3SiC2 above 900 °C, and the reaction products are closely related to the relative ratio of Cu and Ti3SiC2. At low Ti3SiC2 content or temperatures below 1000 °C, Cu(Si) solid solution and TiCx are formed, whereas at high Ti3SiC2 content, Cu– Si intermetallic compounds like Cu5Si, Cu15Si4, and (Cu, Si)η' as well as TiCx were observed. In our study, Ti3SiC2 is the matrix and metallic Cu is the added phase with 5 to 20 wt%, i.e. 2.6 to 11.3 vol% in the matrix. Due to the relative low content of Cu, we only found the intermetallic compound Cu3Si as the reaction product besides the decomposed products TiCx, Ti5Si3Cy, and TiSi2Cz of Ti3SiC2. The solid solution of Cu(Si) was found in the MA powder (Fig. 2) but not in the SPS sintered Ti3SiC2/Cu composites. Barsoum et al. [1] declared that Ti3SiC2 is thermodynamically stable up to at least 1600 °C in vacuum for 24 h and in an argon atmosphere for 4 h. Pang et al. [20] found Ti3SiC2 decomposes through the sublimation of Ti and Si elements forming a surface coating of TiC above 1300 °C. Low and his colleagues [21,22] reported that the surface of Ti3SiC2 starts decomposing to form nonstoichiometric TiC and/or Ti5Si3Cy at 1200 °C and conspicuously decomposes at 1500 °C in the presence of a low oxygen partial pressure. In the present study, Ti3SiC2 decomposes at 1050 °C (Fig. 2) when sintered by SPS not as the reported [23] at 1300 °C. Distinctively, the Cu aggravates the decomposition of Ti3SiC2, and the higher the content of Cu, the more intensively the reaction take place. In fact, the high reaction activity for Ti3SiC2 is closely related with the position of Si atoms in the Ti3SiC2 crystalline structure. As is known, Ti3SiC2 crystalline structure can be described as a layer of Si intercalates into the {111} twin boundary of TiC and the chemical bonding between Si and Ti is relatively weak compared to the strong Ti–C bonding [14]. Consequently, at high temperature the Si atoms with high reactivity favor to de-intercalate from Ti3SiC2 crystalline structure resulting in the structural rearrangement or/and the decomposition of Ti3SiC2. If the added second metal phase can form solid solution or take chemical reaction with Si atoms, the second phase will be a trigger for decomposition of Ti3SiC2. According to the report [13], Si atoms release form Ti3SiC2 crystalline structure leaving Ti3C2 skeleton and diffuse to the adjacent Ti3SiC2 grains and react with them to form Ti5Si3Cy and/or TiSi2Cz. The detailed reactions in the present study are referred as
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Fig. 3. XRD patterns for (a) Ti3SiC2, (b) TSC5, (c) TSC10, (d) TSC15, and (e) TSC20 composites. Table 1 The phases in the prepared composites with the different Cu phase. The content of Cu phase/wt%
The major phase in the prepared composite
The minor phase in the prepared composite
0 5 10 15 20
Ti3SiC2 Ti3SiC2 Ti3SiC2 TiCx TiCx
Ti5Si3Cy TiCx, Ti5Si3Cy, Cu3Si TiCx, Ti5Si3Cy, Cu3Si Ti5Si3Cy, Cu3Si, TiSi2Cz Ti5Si3Cy, Cu3Si, TiSi2Cz
3.2. Mechanical properties
follows: Ti3SiC2-Ti3C2(TiCx)þ Si
(3)
Ti3SiC2 þSi-Ti5Si3Cy
(4)
Ti3SiC2 þSi-TiSi2Cz þTiCx
(5)
Si þCu-Cu3Si
(6)
Because when the content of Cu is higher than 10 wt%, the Ti3SiC2 matrix has completely decomposed, the mechanical and tribological properties of Ti3SiC2/15 and 20 wt% Cu (TSC15, TSC20) are not considered in the present study. The mechanical properties consist of Vickers hardness, flexural and compressive strength of Ti3SiC2/Cu composites and monolithic Ti3SiC2 are listed in Table 2. The measured Vickers hardness of the monolithic Ti3SiC2 is approximately 5.02 GPa. The Vickers hardness of both the composites is higher than that of monolithic Ti3SiC2, i.e. 7.24 GPa for TSC5 and 6.43 GPa for TSC10. The Vickers hardness of Ti3SiC2/Cu composites is the comprehensive result of Ti3SiC2, TiCx, Ti5Si3Cy and Cu3Si, whose Vickers hardness is 5.0, 27.0, 9.5 and 6.70 GPa [24,25]. The higher hardness of the composites than that of monolithic Ti3SiC2 is due to the potentiation of the hard TiCx, Ti5Si3Cy, and Cu3Si. The content of the relatively soft Cu3Si in TSC5 is lower than that in TSC10 so the former is harder than the latter. In addition, the grain refinement, which is caused by MA and fast
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Fig. 4. (a) Backscattered electron image and the (b) EDS analytic result of different regions of backscattered electron image of TSC10 composite.
Table 2 Mechanical properties of monolithic Ti3SiC2 and Ti3SiC2/Cu composites.
Monolithic Ti3SiC2 TSC5 TSC10
Vickers hardness/GPa
Flexural strength/ Ultimate compressive MPa strength/MPa
5.02 7 0.23
434.077 6.8
7.247 0.44 6.43 7 0.57
391.317 8.64 370.677 10.8
949.317 23.28 1651.38 7 102.83 1603.687 39.14
heating and cooling of SPS, may play a positive important role in improving the hardness of Ti3SiC2 composites according to the Hall–Petch relation [26]. The flexural strength of the composites is lower than the monolithic Ti3SiC2 (Table 2). The fracture toughness of monolithic Ti3SiC2 is 7 MPa m1/2 at room temperature [1], which is relatively higher than that of common carbide ceramic materials, such as TiC, SiC, and WC and so on. The multiple energy-absorbing mechanisms, such as diffuse micro-cracking, delamination, crack deflection, grain push-out, grain pull-out, and the buckling of individual grains were found to contribute to the high flexural strength [27,28]. These multiple energy-absorbing mechanisms are fundamentally and closely related to the unique valence bond and crystalline structure. Fig. 5 shows the SEM images of fracture surfaces of monolithic Ti3SiC2 and TSC10. The intergranular fracture is the main failure mechanism accompanied by a few of transgranular fracture for the monolithic Ti3SiC2 (Fig. 5a). As for the fracture surface of TSC10 (Fig. 5b) many fine equiaxed grain particles distribute on the large grains. The fine equiaxed grains in microstructure belong to the reactive products of Cu3Si, TiCx, and Ti5Si3Cy, while the large size particles assign to Ti3SiC2 grains. The intergranular and transgranular fracture can be found for the Ti3SiC2 grains, while intergranular fracture is the main fracture mechanism for the fine-grain reactive products. Some pores and micro-cracks can be found in the fracture surface (as denoted in Fig. 5b). The formed impurities of Cu3Si, TiCx, and Ti5Si3Cy distributed among the Ti3SiC2 grains partly obstacle multiple energy-
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absorbing mechanism, which makes fracture occur between the fine impurities grains. Another important reason for the decrease of fracture strength of the composites is the compatibility (e.g. wettability) between the matrix and the impurities. The mismatch such as the coefficients of thermal expansion (CTEs) easily leads to the weak bonding between grains, intrinsic stresses and micropores and micro-crack in the composite. Grain boundaries play an important role in crack propagation. In the present study the CTE of impurity TiC (7.4 10 6/K) is lower than that of Ti3SiC2 (9.1 10 6/K) matrix, which may lead to the formation of micropores and crack as denoted in Fig. 5b. Crack initiation usually takes place by intrinsic stresses, based on the impurities and the structural anisotropy. Additionally, the silicides (the reaction products) of Cu3Si and Ti5Si3Cy are brittle phases and exhibited low fracture toughness at room temperature, especially the fracture toughness of Ti5Si3Cy is only 2.1 MPa m1/2 [29]. The ultimate compressive strengths for TSC5 and TSC10 composites are 1651.38 and 1603.68 MPa, respectively, which is much higher than that of monolithic Ti3SiC2 (Table 2). From the typical stress-strain curve of TSC10 (Fig. 6), the composite exhibit linearity elastic deformation stage with no distinctly plastic deformation stage, an upper yield point followed by a stress drop, which indicates the composites experience brittle fracture behavior. As was proposed in [30], the yield point arise from the high particle volume fraction, which yields hard phase particle percolation. In fact, yield stress is related to the slip systems for dislocation glide and jogged screw dislocations. Many fine Cu3Si, TiCx, and Ti5Si3Cy grains distributed in the relatively coarse Ti3SiC2 grains can effectively prohibit the activated dislocation glide and jogged screw dislocations, which leads to the higher compressive strength. 3.3. Tribological properties The real-time friction coefficient μr, as a function of sliding time of monolithic Ti3SiC2 and Ti3SiC2/Cu composites at room temperature (RT) are shown in Fig. 7a. The μr of the monolithic Ti3SiC2 is stable to be about 0.79 after a 3 min running-in period, while the μrs of TSC5 and TSC10 composites are stable to be 0.48 and 0.54, respectively, after about a 5 min running-in period. The mean μms are shown in Table 3. The WR of TSC5 is 2.47 10 4 mm3/ NUm and TSC10 4.11 10 4 mm3/N Um, both of which are lower than that of monolithic Ti3SiC2, 6.39 10 4 mm3/N Um (Table 3). It can be referred that the lubrication and anti-wear properties of the composites both are superior to the monolithic Ti3SiC2. Because Ti3SiC2 possesses excellent mechanical and thermal properties at high temperature [1,15] and may be used as high-temperature structure and function material, we also investigated the tribology properties of at elevated temperatures. The time dependencies of μrs for TSC10 at different temperatures are shown in Fig. 7b and the mean μm and the WRs of TSC10 are listed in Table 3. At elevated temperatures volatility of μr is relatively larger than that at RT. The mean μms at 200, 400 and 600 °C are 0.85 70.122, 0.65 70.19 and 0.57 70.10, all of which are higher than that at RT. The WRs of TSC10 at elevated temperatures are also higher than that at RT (Table 3). It can be concluded that the tribological properties of Ti3SiC2/Cu composite at elevated temperatures are worse than that at RT. Fig. 8a shows the SEM image of the worn surface of monolithic Ti3SiC2. Mass abrasive particles scatter on the worn surface accompanied by local plastic deformation. The fractured and pulledout grains in the worn surface can be distinguished from the plastic-deformation region and the edge. We can also see a mass of lamellar peeling structures which are going to spalling off (as denoted in Fig. 8a). This may be because Ti3SiC2 has a lamellar structure which has a characteristic of being prone to cleave. These lamellar peeling structures come to spalling off gradually with the
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Fig. 5. SEM images of the fractured surface of (a) Ti3SiC2 and (b) TSC10 composite.
Fig. 6. The ultimate force as a function of strain curve of TSC10 composites.
action of circulatory shear stress. The fractured and pulled-out Ti3SiC2 grains could evolve into three-body abrasive wear in the process of rubbing. In addition, there also exists some plastic deformations in some locations of wear scar of Ti3SiC2. It may be due to the lower hardness of Ti3SiC2 and Ti3SiC2 will suffer heavy shear stress when it rubs with SiC ball. The threeUbody abrasive wear and lamellar peeling will lead to heavy loss of material in line with the high WR. So the wear mechanism for monolithic Ti3SiC2 bulk is abrasive wear, which derives from the fracture and pull out of grains accompanied by local plastic deformation. These results we acquired are consistent with others [5–7]. The fracture and pull out of grains are attributed to the relatively weak interfacial bonding force between Ti3SiC2 grains. Fig. 8b shows the worn surface of TSC10 composite at RT. The worn surface of TSC10 is smoother than that of monolithic Ti3SiC2. We can see plentiful of wear furrows on the worn surface, which is the symptom of abrasive wear. It indicates that abrasive wear takes place on the worn surface of TSC10 composite, while no abundant lamellar peeling structures as that on the worn surface of Ti3SiC2. It can be speculated that the fine grains of TiCx, Ti5Si3Cy, and Cu3Si with high hardness can pin the relatively soft Ti3SiC2 matrix, which prevents grains lamellar peeling and pulling out. The wear resistance of TSC10 has a significant improvement compared with the monolithic Ti3SiC2 bulk due to lighter abrasive wear and significant increase in hardness. In addition, the hardness increase can inhibit the plastic deformation effectively, which can also decrease friction and wear. The worn surface of TSC5
Fig. 7. Real-time friction coefficients curves versus sliding time for (a) monolithic Ti3SiC2 and its composites at RT and (b) TSC10 composite at elevated temperatures.
composite is similar to that of TSC10 (the SEM image of worn surface of TSC5 is omitted for short). The SEM images of worn surfaces of TSC10 at elevated temperatures are shown in Fig. 8c–e. With the temperature increasing, the character of plastic flow on the worn surfaces becomes increasingly obvious. The area of the islands of plastic flow is
Please cite this article as: W. Dang, et al., Influence of Cu on the mechanical and tribological properties of Ti3SiC2, Ceramics International (2016), http://dx.doi.org/10.1016/j.ceramint.2016.03.099i
W. Dang et al. / Ceramics International ∎ (∎∎∎∎) ∎∎∎–∎∎∎ Table 3 The mean μms and WRs of monolithic Ti3SiC2 and Ti3SiC2/Cu composites at RT and elevated temperatures.
Monolithic Ti3SiC2 TSC5 TSC10 TSC10 TSC10 TSC10
Mean WR (mm3/N ·m)
Test temperature (°C)
μm
RT
0.79 7 0.023 6.39 10 4 71.25 10 5
RT RT 200 400 600
0.487 0.040 0.54 7 0.085 0.85 7 0.12 0.65 7 0.19 0.577 0.10
2.79 10 4 72.76 10 5 4.11 10 4 7 7.40 10 6 4.44 10 3 7 2.38 10 4 1.03 10 3 7 2.10 10 5 6.96 10 4 71.17 10 4
growing with the decrease of scattered abrasive particles. Plastic flow results to the adhesive friction and wear for the composite at elevated temperatures. Additionally, on the worn scar of the tribo-
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pair SiC ball at 600 °C covers a mechanical mixed layer as shown in Fig. 9a. The EDS analysis (Fig. 9b) shows that the mechanical layer consists Ti and Cu elements, which indicates that material transference from Ti3SiC2/Cu composite to the counter SiC ball occurred during the rubbing. Plastic adhesive wear and material transference replacing the abrasive wear becomes the dominant wear mechanism at elevated temperatures, which leads to the higher WRs than that at RT. The chemical states of Ti2p (Fig. 10a), Si2p (Fig. 10b), and Cu2p (Fig. 10c) on the worn surface of TSC10 composite at RT and 600 °C were analyzed by XPS and the results are shown in Fig. 10. It can be seen that the Ti, Si and Cu species are partly oxidized to form TiO2, SiO2, and CuO at RT. At 600 °C, tribo-oxidation is more severe. The elements of Ti, Si, and Cu are all completely oxidized. So tribooxidation wear is another wear mechanism for the composites. As is well known, the unique valence bond structure makes
Fig. 8. SEM micrographs of worn surfaces of (a) Ti3SiC2, (b) TSC10 at RT and TSC10 at (c) 200, (d) 400, and (e) 600 °C.
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Fig. 9. (a) The SEM image and (b) the EDS analysis of worn scar of SiC ball rubbing with TSC10 at 600 °C.
Ti3SiC2 be characterized by the unusual combination of metallic and ceramic properties [1]. When the temperature reaches 1200 °C, the brittle-ductile transition will take place for Ti3SiC2. While in the process of friction, the worn surface undergoes ductile under the flash temperature derived from the friction shearing action. On the other hand, it was reported that continuous and homogeneous tribo-oxidation scales formed on the worn surface at high temperature or high sliding speed can endow Ti3SiC2 excellent lubricating and anti-wear properties [31–33]. However, in the present study though tribo-oxidation occur on the worn surface, the oxidation sale is not continuous and homogeneous and can not act as lubrication layer to reduce friction and wear. In conclusion, there exists different level of abrasive wear for Ti3SiC2 and the Ti3SiC2/Cu composites at RT. For Ti3SiC2, there is local plastic deformation besides abrasive wear. For the TSC5 and TSC10 composites, the main wear mechanism is abrasive wear. In a word, the addition of Cu to the polycrystalline Ti3SiC2 effectively prevents the pull-out and fracture of grains and so leads to the decrease of friction and improvement of wear resistance. At elevated temperatures, plastic adhesive wear and tribo-oxidation wear along with material transference become the dominant wear mechanism, which leads to the higher WRs than that at RT.
4. Conclusions Ti3SiC2/Cu composites with different contents of Cu were fabricated and the mechanical and tribological properties of the composites were investigated compared with monolithic Ti3SiC2 in details. The conclusions are as follows: a. The chemical reactions take place to produce Cu3Si and TiCx at the process of fabrication of the composites by SPS and are
Fig. 10. X-ray photoelectron spectroscopy for (a) Ti2p (a), (b) Si2p, and (c) Cu2p on the worn surface of TSC10 composite at RT and 600 °C.
more intense for the composites with high content of Cu (15 and 20 wt%). The reaction products consist of TiCx, Ti5Si3Cy, Cu3Si, and TiSi2Cz. Cu3Si phase distributes around the Ti3SiC2 particles. b. The hardness of Ti3SiC2 composites is much higher than that of monolithic Ti3SiC2, which is derived from the enhancement of
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TiCx, Ti5Si3Cy, and Cu3Si. The flexural strength of the composites is lower than the monolithic Ti3SiC2. The fine impurities of Cu3Si, TiCx, and Ti5Si3Cy partly obstacle multiple energyabsorbing mechanism. Many fine Cu3Si, TiCx, and Ti5Si3Cy grains distributed in the relatively coarse Ti3SiC2 grains effectively prohibit the activated dislocation glide and jogged screw dislocations, which leads to the higher compressive strength for the composites. c. The friction and wear of Ti3SiC2/Cu composites is lower than that of monolithic Ti3SiC2 at RT, which is ascribed to the fixing effect of hard particle of Cu3Si, TiCx, and Ti5Si3Cy to reduce the abrasive friction and wear. While at elevated temperatures plastic adhesive wear and tribo-oxidation wear accompanied by material transference contribute to the increased friction and wear for the composites.
Acknowledgements The authors acknowledge the financial support from the National Natural Science Foundation of China (Grant nos. 51405475, 51475444, and 51505464).
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Ceramics International journal homepage: www.elsevier.com/locate/ceramint
Influence of Cu on the mechanical and tribological properties of Ti3SiC2 Wentao Dang a,b, Shufang Ren a,n, Jiansong Zhou a,n, Youjun Yu a, Zhen Li a, Lingqian Wang a a b
State Key Laboratory of Solid Lubrication, Lanzhou Institute of Chemical Physics, Chinese Academy of Sciences, Lanzhou 730000, PR China Graduate School of Chinese Academy of Sciences, Beijing 100039, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 6 February 2016 Received in revised form 1 March 2016 Accepted 12 March 2016
Ti3SiC2/Cu composites with different contents of Cu were fabricated by mechanical alloying and spark plasma sintering method. The phase composition and structure of the composites were analyzed by X-ray diffractometry and scanning electron microscopy equipped with energy dispersive spectroscopy. The mechanical and tribological properties of Ti3SiC2/Cu composites were tested and analyzed compared with monolithic Ti3SiC2 in details. The results show that the Cu leads to the decomposition of Ti3SiC2 to produce TiCx, Ti5Si3Cy, Cu3Si, and TiSi2Cz. The friction coefficient and wear rate of the composites are lower than that of monolithic Ti3SiC2, which is ascribed to the fixing effect of hard TiCx, Ti5Si3Cy, and Cu3Si to inhibit the abrasive friction and wear. However, at elevated temperatures (ranging from room temperature to 600 °C) the friction and wear of the composites are higher than those at room temperature. Plastic flowing and tribo-oxidation wear accompanied by material transference contribute to the increased friction and wear at elevated temperatures. & 2016 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Keywords: C. Mechanical properties Ti3SiC2/Cu composites Tribological property Abrasive wear Tribo-oxidation wear SPS
1. Introduction Ti3SiC2 has recently attracted extensive research attention due to its combination of the merits of both metals and ceramics. The unique properties of Ti3SiC2 are derived from the unusual valence bond structure that is there are covalent, ionic, and metal three bonds in its crystalline structure [1–3]. Though Ti3SiC2 possesses layered structure it is not intrinsically self-lubricating for the polycrystalline Ti3SiC2. In fact, many studies [4–7] have found that the three-body abrasive wear that derives from the fractured and pulled-out Ti3SiC2 grains is the main wear mechanism at room temperature, which leads to the high friction and wear. The fundamental reason why Ti3SiC2 grains are inclined to fracture and pull out is that the grain boundary force of Ti3SiC2 grains is weak [8]. This intrinsic structure determines its some apparent defective performances, such as the low hardness (4–6 GPa), moderate flexural strength (200–600 MPa), and the brittleness like ceramics at 25–1100 °C [2,3]. These characteristics limit the wide applications of Ti3SiC2 as structural and functional material. Many researchers focus on the strengthening of Ti3SiC2 by the hard phases such as Al2O3 [9], TiC [10], Ti5Si3 [11], and WC-Co etc [12]. The mechanical and tribological properties of Ti3SiC2/ceramic composites are better than that of monolithic Ti3SiC2. Nevertheless, as for Ti3SiC2/metal composite, the corresponding research is very limited. To the knowledge of authors, Ti3SiC2 is chemically n
Corresponding authors. E-mail addresses: [email protected] (S. Ren), [email protected] (J. Zhou).
reactive when sintering with metals at high temperature. Li et al. [13] synthesized Ti3SiC2/Ni and Ti3SiC2/Co through vacuum sintering and they found the metals aggregated towards the sample surfaces resulting in layered microstructural features because the wettability between Ti3SiC2 and Ni or Co are both poor. They suggested some adequate sintering should be added to improve the wetting performance and interface properties between Ti3SiC2 and metals to achieve homogeneous microstructures and excellent mechanical properties. Many researchers reported the reactivity between Ti3SiC2 and metals. Lu et al. [14] found chemical reaction between Cu and Ti3SiC2 contributes to the wettability. Zhou and his coworkers [15] investigated chemical reactions and stability of Ti3SiC2 in Cu during high-temperature processing of Cu/Ti3SiC2 composites. The results indicates that Ti3SiC2 can react with Cu above 900 °C. When the low content of Ti3SiC2 or below 1000 °C, Cu(Si) solid solution and TiCx were formed, whereas at high temperature or high content of Ti3SiC2, Cu–Si intermetallic compounds like Cu5Si, Cu15Si4, and TiCx were formed. So it is critical to control the content of Cu to prevent total Ti3SiC2 from reacting with Cu. Gu et al. [16] investigated the chemical reactions between Ti and Ti3SiC2 during the preparation of Ti–Ti3SiC2 composites in the temperature range of 1000–1300 °C. The results demonstrate that Ti3SiC2 reacts with Ti to form TiCx, Ti5Si3, and TiSi2. They proposed that it is difficult to synthesize composite with titanium and Ti3SiC2 because the reactions between Ti and Ti3SiC2 occur at 1000 °C, which is lower than the melting point of Ti (1670 °C). Copper is a soft metal with good ductility, thermal and electrical conductivity, and excellent mechanical properties, which is commonly used as the cable, electrical and electronic components,
http://dx.doi.org/10.1016/j.ceramint.2016.03.099 0272-8842/& 2016 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
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and building materials. Ti3SiC2 is a potential structure/functional material with many unique properties such as high modulus and strength, high electrical conductivity, good thermal conductivity, good thermal stability, and excellent high-temperature oxidation resistance and so on. However, the poor tribological properties that are traced to the weak intergranular bonding strength limit its application as sliding components that serves in sliding friction conditions. Given the good ductility, electrical conductivity, and the excellent mechanical properties, in the present paper, copper was selected to strengthen the intergranular bonding strength of Ti3SiC2 grains and to then improve the tribological properties, low hardness, and the room-temperature brittleness for Ti3SiC2. The aim of this study is to synthesize Ti3SiC2/Cu composite and investigate the phase composition, mechanical and tribological properties. We hope the mechanical properties and the wear resistance of monolithic Ti3SiC2 can be improved by adding Cu element into Ti3SiC2.
be referred to in two ways: μm to refer to the mean friction coefficient over the entire sliding distance, and μr to refer to its real-time value as a function of sliding time. The μrs were recorded automatically by the tribometer. The wear rates (WR) of the Ti3SiC2/Cu samples were calculated by the following formula [17]:
W = V /PS
(1)
where W is the wear rate, V the wear volume (mm3), P the applied load (N), and S the total sliding distance (m). The profile of cross section of worn surface was outlined using NaNomap 500LS scanning three-dimension surface profiler to determine the wear volume as
(2)
V = AL 3
where V is the wear volume (mm ), A the cross section area of the worn surface (mm), and L the perimeter of wear track (mm). 2.4. Characterization
2. Material and methods 2.1. Samples preparation The commercially available Ti3SiC2 (YueHuan New Material Technology Co., Ltd. 99 wt%, 200 mesh) powder and Cu (Sinopharm Chemical Reagent Co., Ltd, Shanghai, China, 99%, 200 mesh) powder were used in this study. In order to determine the optimum Cu content in the Ti3SiC2/Cu composites, samples with different mass fractions of Cu (5, 10, 15, 20 wt%) were prepared by mechanical alloying (MA) and the following accompanied spark plasma sintering (SPS). The mechanical alloying was carried out by high energy ball-milling machine (Pulverisettes 4, Germany) with a speed of 300 rpm and a ratio of ball to powder of 12 for 6 h. The mixed powder was sintered at 1050 °C for 5 min with the heating rate of 60 °C/min at a cylindrical graphite mold with a 25 mm diameter in vacuum ( o10 Pa) under a pressure of 30 MPa by SPS (SPS-20T-10, Shanghai Chen Hua Technology Co., Ltd.). 2.2. Mechanical properties The Vickers hardness was measured by the MH-5-VM microhardness tester under applied 10 N load with the dwell time of 5 s. The hardness was calculated by averaging at least 8 measurements. Flexural strength and compression strength were measured using a universal material tester (SANS-CMT5205, MTS, China). The three-point bending test was conducted to obtain the flexural strength. The sample was cut to be 3 4 20 mm3. The cross head speed was 0.5 mm/min and the support span was 16 mm. The sizes of samples for compressive strength test were Φ5 mm 12.5 mm and the cross head speed was 0.2 mm/min. 2.3. Friction and wear test The friction and wear tests were performed using a high temperature ball-on-disc tribometer (HT-1000, Zhong Ke Kai Te Ltd. China) capable of going up to 900 °C. All the tests were done at a linear velocity of 0.2 m/s and a load of 5 N corresponding to a stress of ∼0.1 MPa under a non-lubrication condition with a rotation radius of 5 mm. The sliding distance for all the friction and wear tests was 360 m. The Ti3SiC2/Cu composites were shaped into cylindrical having a diameter of 25 mm working face and ∼5 mm thick. The counter surfaces were SiC ceramic ball with a diameter of 6.43 mm (density of 3.2 g/cm3 and a hardness of 28 GPa, K&H, Japan). All surfaces were polished to a 1 μm diamond finish, washed with ethanol and dried prior to testing. The measured μ's will
The phase compositions of the composites were determined by D/MAX-2000 X-ray diffractometer (XRD) with Cukɑ radiation in wavelength λ ¼ 0.1540598 nm. The microstructures of the specimens were examined by scanning electron microscopy (SEM, JEOL-5600LV) equipped with energy dispersive spectroscopy (EDS). The chemical states of typical elements on the worn surfaces were determined using a PHI-5702 X-ray photoelectron spectroscope (XPS) with a resolution of 0.2 eV using Al Ka radiation as the excitation source and the binding energy of carbon (C1s: 284.8 eV) as the reference.
3. Results and discussion 3.1. Phase composition and microstructure Fig. 1a and b shows the SEM micrographs of Ti3SiC2 powder and Ti3SiC2/10 wt% Cu (noted as TSC10) mixed powder after MA, respectively. It can be observed that the grain size of Ti3SiC2 particles ranges from 2 to 12 μm and the grain size of mixed powder of Ti3SiC2 and 10 wt% Cu becomes smaller than that of Ti3SiC2, ranging from 1 to 6 um after MA. The XRD patterns for Ti3SiC2 powder and Ti3SiC2 with 10 wt% Cu after MA are shown in Fig. 2. As for the MA mixed powder the diffraction peaks of Ti3SiC2 obviously broaden, which means Ti3SiC2 grains were refined during highenergy ball milling. Besides the peaks of Ti3SiC2 and Cu, the diffraction peaks of TiCx and Cu(Si) solid solution appear, which indicates that Cu reacts with Ti3SiC2 producing TiCx and Cu(Si) solid solution during high-energy ball milling. Mechanical alloying (MA) is a solid-state powder processing technique involving repeated welding, fracturing, and rewelding of powder particles in a highenergy ball mill [18]. MA has now been shown to be capable of synthesizing a variety of non-equilibrium alloy phases, such as supersaturated solid solutions, intermetallics, nanostructures, and amorphous alloys and so on starting from blended elementals powders. Recent mechanochemical synthesis of materials have been critically reviewed after discussing the process and process variables involved in MA [19]. So in the present study solid solutions of Cu(Si), mechanochemical product of TiCx derived from the decomposition of Ti3SiC2 were found besides the powders refining phenomenon. Additionally, according to our experience of preparation Ti3SiC2/Cu the MA processing can help to improve the density and the purity of the composites. Fig. 3 shows the XRD patterns of Ti3SiC2/Cu composites with different Cu contents sintered by SPS. It can be seen that the diffraction peaks of elementary Cu does not emerge in all the
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Fig. 1. SEM micrographs of (a) Ti3SiC2 powder and (b) TSC10 mixed powder after MA.
Fig. 2. XRD patterns for (a) Ti3SiC2 powder and (b) Ti3SiC2 with 10 wt% Cu powder after MA.
composites owing to the reaction of Cu with Ti3SiC2 to produce TiCx, Ti5Si3Cy, and Cu3Si during sintering. The intensity of the diffraction peaks of Ti3SiC2 weakens with the content of Cu increasing in the composites. The major ingredient of Ti3SiC2/5 wt% Cu (TSC5) and TSC10 composites is Ti3SiC2 with a small amount of Ti5Si3Cy, TiCx, and Cu3Si. The relative content of Cu3Si in TSC10 composite is higher than that of TSC5 by semiquantitative analysis. With the increase of Cu, the diffraction peaks of Ti3SiC2 matrix almost disappear and the major phase composition becomes TiCx, Ti5Si3Cy, TiSi2Cz, and Cu3Si (as shown in Fig. 3d and e). The detail products are listed in Table 1. In conclusion, with the content increase of Cu, the chemical reaction between Ti3SiC2 matrix and Cu takes place more and more intensively during the sintering by SPS. Fig. 4 shows the backscattered electron topography of sintered TSC10 composite. In the image we can see that the white punctate and linear regions distribute between the grey regions. According to the analysis of corresponding EDS, the white region abounds in Cu element and lacks Ti element while the grey region abounds in Ti element and lacks Cu element. So it can be referred that the white regions are mainly composed of Cu3Si and the grey regions are consist of the major Ti3SiC2 and the minor TiCx and Ti5Si3Cy. Cu3Si phase distributes around the Ti3SiC2 particles. From above results of phase and microstructure of the composites it can be concluded that the Cu element reacts with Ti3SiC2 matrix during the process of sintering. The reactive products vary
when different contents of Cu were added. The monolithic Ti3SiC2 decomposed to produce minor Ti5Si3Cy during sintering at 1050 °C with the soaking time of 5 min by SPS. For the 5 and 10 wt% Cu, the main phase for the composites is Ti3SiC2 with minor TiCx, Ti5Si3Cy, and Cu3Si. When the content of Cu is above 10 wt%, the Ti3SiC2 matrix completely decomposes to produce major TiCx and minor Ti5Si3Cy, TiSi2, and Cu3Si. Zhou et al. [15] investigated the chemical reactions of Ti3SiC2 with 20 to 70 vol% in Cu matrix. They found that Cu reacts with Ti3SiC2 above 900 °C, and the reaction products are closely related to the relative ratio of Cu and Ti3SiC2. At low Ti3SiC2 content or temperatures below 1000 °C, Cu(Si) solid solution and TiCx are formed, whereas at high Ti3SiC2 content, Cu– Si intermetallic compounds like Cu5Si, Cu15Si4, and (Cu, Si)η' as well as TiCx were observed. In our study, Ti3SiC2 is the matrix and metallic Cu is the added phase with 5 to 20 wt%, i.e. 2.6 to 11.3 vol% in the matrix. Due to the relative low content of Cu, we only found the intermetallic compound Cu3Si as the reaction product besides the decomposed products TiCx, Ti5Si3Cy, and TiSi2Cz of Ti3SiC2. The solid solution of Cu(Si) was found in the MA powder (Fig. 2) but not in the SPS sintered Ti3SiC2/Cu composites. Barsoum et al. [1] declared that Ti3SiC2 is thermodynamically stable up to at least 1600 °C in vacuum for 24 h and in an argon atmosphere for 4 h. Pang et al. [20] found Ti3SiC2 decomposes through the sublimation of Ti and Si elements forming a surface coating of TiC above 1300 °C. Low and his colleagues [21,22] reported that the surface of Ti3SiC2 starts decomposing to form nonstoichiometric TiC and/or Ti5Si3Cy at 1200 °C and conspicuously decomposes at 1500 °C in the presence of a low oxygen partial pressure. In the present study, Ti3SiC2 decomposes at 1050 °C (Fig. 2) when sintered by SPS not as the reported [23] at 1300 °C. Distinctively, the Cu aggravates the decomposition of Ti3SiC2, and the higher the content of Cu, the more intensively the reaction take place. In fact, the high reaction activity for Ti3SiC2 is closely related with the position of Si atoms in the Ti3SiC2 crystalline structure. As is known, Ti3SiC2 crystalline structure can be described as a layer of Si intercalates into the {111} twin boundary of TiC and the chemical bonding between Si and Ti is relatively weak compared to the strong Ti–C bonding [14]. Consequently, at high temperature the Si atoms with high reactivity favor to de-intercalate from Ti3SiC2 crystalline structure resulting in the structural rearrangement or/and the decomposition of Ti3SiC2. If the added second metal phase can form solid solution or take chemical reaction with Si atoms, the second phase will be a trigger for decomposition of Ti3SiC2. According to the report [13], Si atoms release form Ti3SiC2 crystalline structure leaving Ti3C2 skeleton and diffuse to the adjacent Ti3SiC2 grains and react with them to form Ti5Si3Cy and/or TiSi2Cz. The detailed reactions in the present study are referred as
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Fig. 3. XRD patterns for (a) Ti3SiC2, (b) TSC5, (c) TSC10, (d) TSC15, and (e) TSC20 composites. Table 1 The phases in the prepared composites with the different Cu phase. The content of Cu phase/wt%
The major phase in the prepared composite
The minor phase in the prepared composite
0 5 10 15 20
Ti3SiC2 Ti3SiC2 Ti3SiC2 TiCx TiCx
Ti5Si3Cy TiCx, Ti5Si3Cy, Cu3Si TiCx, Ti5Si3Cy, Cu3Si Ti5Si3Cy, Cu3Si, TiSi2Cz Ti5Si3Cy, Cu3Si, TiSi2Cz
3.2. Mechanical properties
follows: Ti3SiC2-Ti3C2(TiCx)þ Si
(3)
Ti3SiC2 þSi-Ti5Si3Cy
(4)
Ti3SiC2 þSi-TiSi2Cz þTiCx
(5)
Si þCu-Cu3Si
(6)
Because when the content of Cu is higher than 10 wt%, the Ti3SiC2 matrix has completely decomposed, the mechanical and tribological properties of Ti3SiC2/15 and 20 wt% Cu (TSC15, TSC20) are not considered in the present study. The mechanical properties consist of Vickers hardness, flexural and compressive strength of Ti3SiC2/Cu composites and monolithic Ti3SiC2 are listed in Table 2. The measured Vickers hardness of the monolithic Ti3SiC2 is approximately 5.02 GPa. The Vickers hardness of both the composites is higher than that of monolithic Ti3SiC2, i.e. 7.24 GPa for TSC5 and 6.43 GPa for TSC10. The Vickers hardness of Ti3SiC2/Cu composites is the comprehensive result of Ti3SiC2, TiCx, Ti5Si3Cy and Cu3Si, whose Vickers hardness is 5.0, 27.0, 9.5 and 6.70 GPa [24,25]. The higher hardness of the composites than that of monolithic Ti3SiC2 is due to the potentiation of the hard TiCx, Ti5Si3Cy, and Cu3Si. The content of the relatively soft Cu3Si in TSC5 is lower than that in TSC10 so the former is harder than the latter. In addition, the grain refinement, which is caused by MA and fast
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Fig. 4. (a) Backscattered electron image and the (b) EDS analytic result of different regions of backscattered electron image of TSC10 composite.
Table 2 Mechanical properties of monolithic Ti3SiC2 and Ti3SiC2/Cu composites.
Monolithic Ti3SiC2 TSC5 TSC10
Vickers hardness/GPa
Flexural strength/ Ultimate compressive MPa strength/MPa
5.02 7 0.23
434.077 6.8
7.247 0.44 6.43 7 0.57
391.317 8.64 370.677 10.8
949.317 23.28 1651.38 7 102.83 1603.687 39.14
heating and cooling of SPS, may play a positive important role in improving the hardness of Ti3SiC2 composites according to the Hall–Petch relation [26]. The flexural strength of the composites is lower than the monolithic Ti3SiC2 (Table 2). The fracture toughness of monolithic Ti3SiC2 is 7 MPa m1/2 at room temperature [1], which is relatively higher than that of common carbide ceramic materials, such as TiC, SiC, and WC and so on. The multiple energy-absorbing mechanisms, such as diffuse micro-cracking, delamination, crack deflection, grain push-out, grain pull-out, and the buckling of individual grains were found to contribute to the high flexural strength [27,28]. These multiple energy-absorbing mechanisms are fundamentally and closely related to the unique valence bond and crystalline structure. Fig. 5 shows the SEM images of fracture surfaces of monolithic Ti3SiC2 and TSC10. The intergranular fracture is the main failure mechanism accompanied by a few of transgranular fracture for the monolithic Ti3SiC2 (Fig. 5a). As for the fracture surface of TSC10 (Fig. 5b) many fine equiaxed grain particles distribute on the large grains. The fine equiaxed grains in microstructure belong to the reactive products of Cu3Si, TiCx, and Ti5Si3Cy, while the large size particles assign to Ti3SiC2 grains. The intergranular and transgranular fracture can be found for the Ti3SiC2 grains, while intergranular fracture is the main fracture mechanism for the fine-grain reactive products. Some pores and micro-cracks can be found in the fracture surface (as denoted in Fig. 5b). The formed impurities of Cu3Si, TiCx, and Ti5Si3Cy distributed among the Ti3SiC2 grains partly obstacle multiple energy-
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absorbing mechanism, which makes fracture occur between the fine impurities grains. Another important reason for the decrease of fracture strength of the composites is the compatibility (e.g. wettability) between the matrix and the impurities. The mismatch such as the coefficients of thermal expansion (CTEs) easily leads to the weak bonding between grains, intrinsic stresses and micropores and micro-crack in the composite. Grain boundaries play an important role in crack propagation. In the present study the CTE of impurity TiC (7.4 10 6/K) is lower than that of Ti3SiC2 (9.1 10 6/K) matrix, which may lead to the formation of micropores and crack as denoted in Fig. 5b. Crack initiation usually takes place by intrinsic stresses, based on the impurities and the structural anisotropy. Additionally, the silicides (the reaction products) of Cu3Si and Ti5Si3Cy are brittle phases and exhibited low fracture toughness at room temperature, especially the fracture toughness of Ti5Si3Cy is only 2.1 MPa m1/2 [29]. The ultimate compressive strengths for TSC5 and TSC10 composites are 1651.38 and 1603.68 MPa, respectively, which is much higher than that of monolithic Ti3SiC2 (Table 2). From the typical stress-strain curve of TSC10 (Fig. 6), the composite exhibit linearity elastic deformation stage with no distinctly plastic deformation stage, an upper yield point followed by a stress drop, which indicates the composites experience brittle fracture behavior. As was proposed in [30], the yield point arise from the high particle volume fraction, which yields hard phase particle percolation. In fact, yield stress is related to the slip systems for dislocation glide and jogged screw dislocations. Many fine Cu3Si, TiCx, and Ti5Si3Cy grains distributed in the relatively coarse Ti3SiC2 grains can effectively prohibit the activated dislocation glide and jogged screw dislocations, which leads to the higher compressive strength. 3.3. Tribological properties The real-time friction coefficient μr, as a function of sliding time of monolithic Ti3SiC2 and Ti3SiC2/Cu composites at room temperature (RT) are shown in Fig. 7a. The μr of the monolithic Ti3SiC2 is stable to be about 0.79 after a 3 min running-in period, while the μrs of TSC5 and TSC10 composites are stable to be 0.48 and 0.54, respectively, after about a 5 min running-in period. The mean μms are shown in Table 3. The WR of TSC5 is 2.47 10 4 mm3/ NUm and TSC10 4.11 10 4 mm3/N Um, both of which are lower than that of monolithic Ti3SiC2, 6.39 10 4 mm3/N Um (Table 3). It can be referred that the lubrication and anti-wear properties of the composites both are superior to the monolithic Ti3SiC2. Because Ti3SiC2 possesses excellent mechanical and thermal properties at high temperature [1,15] and may be used as high-temperature structure and function material, we also investigated the tribology properties of at elevated temperatures. The time dependencies of μrs for TSC10 at different temperatures are shown in Fig. 7b and the mean μm and the WRs of TSC10 are listed in Table 3. At elevated temperatures volatility of μr is relatively larger than that at RT. The mean μms at 200, 400 and 600 °C are 0.85 70.122, 0.65 70.19 and 0.57 70.10, all of which are higher than that at RT. The WRs of TSC10 at elevated temperatures are also higher than that at RT (Table 3). It can be concluded that the tribological properties of Ti3SiC2/Cu composite at elevated temperatures are worse than that at RT. Fig. 8a shows the SEM image of the worn surface of monolithic Ti3SiC2. Mass abrasive particles scatter on the worn surface accompanied by local plastic deformation. The fractured and pulledout grains in the worn surface can be distinguished from the plastic-deformation region and the edge. We can also see a mass of lamellar peeling structures which are going to spalling off (as denoted in Fig. 8a). This may be because Ti3SiC2 has a lamellar structure which has a characteristic of being prone to cleave. These lamellar peeling structures come to spalling off gradually with the
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Fig. 5. SEM images of the fractured surface of (a) Ti3SiC2 and (b) TSC10 composite.
Fig. 6. The ultimate force as a function of strain curve of TSC10 composites.
action of circulatory shear stress. The fractured and pulled-out Ti3SiC2 grains could evolve into three-body abrasive wear in the process of rubbing. In addition, there also exists some plastic deformations in some locations of wear scar of Ti3SiC2. It may be due to the lower hardness of Ti3SiC2 and Ti3SiC2 will suffer heavy shear stress when it rubs with SiC ball. The threeUbody abrasive wear and lamellar peeling will lead to heavy loss of material in line with the high WR. So the wear mechanism for monolithic Ti3SiC2 bulk is abrasive wear, which derives from the fracture and pull out of grains accompanied by local plastic deformation. These results we acquired are consistent with others [5–7]. The fracture and pull out of grains are attributed to the relatively weak interfacial bonding force between Ti3SiC2 grains. Fig. 8b shows the worn surface of TSC10 composite at RT. The worn surface of TSC10 is smoother than that of monolithic Ti3SiC2. We can see plentiful of wear furrows on the worn surface, which is the symptom of abrasive wear. It indicates that abrasive wear takes place on the worn surface of TSC10 composite, while no abundant lamellar peeling structures as that on the worn surface of Ti3SiC2. It can be speculated that the fine grains of TiCx, Ti5Si3Cy, and Cu3Si with high hardness can pin the relatively soft Ti3SiC2 matrix, which prevents grains lamellar peeling and pulling out. The wear resistance of TSC10 has a significant improvement compared with the monolithic Ti3SiC2 bulk due to lighter abrasive wear and significant increase in hardness. In addition, the hardness increase can inhibit the plastic deformation effectively, which can also decrease friction and wear. The worn surface of TSC5
Fig. 7. Real-time friction coefficients curves versus sliding time for (a) monolithic Ti3SiC2 and its composites at RT and (b) TSC10 composite at elevated temperatures.
composite is similar to that of TSC10 (the SEM image of worn surface of TSC5 is omitted for short). The SEM images of worn surfaces of TSC10 at elevated temperatures are shown in Fig. 8c–e. With the temperature increasing, the character of plastic flow on the worn surfaces becomes increasingly obvious. The area of the islands of plastic flow is
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W. Dang et al. / Ceramics International ∎ (∎∎∎∎) ∎∎∎–∎∎∎ Table 3 The mean μms and WRs of monolithic Ti3SiC2 and Ti3SiC2/Cu composites at RT and elevated temperatures.
Monolithic Ti3SiC2 TSC5 TSC10 TSC10 TSC10 TSC10
Mean WR (mm3/N ·m)
Test temperature (°C)
μm
RT
0.79 7 0.023 6.39 10 4 71.25 10 5
RT RT 200 400 600
0.487 0.040 0.54 7 0.085 0.85 7 0.12 0.65 7 0.19 0.577 0.10
2.79 10 4 72.76 10 5 4.11 10 4 7 7.40 10 6 4.44 10 3 7 2.38 10 4 1.03 10 3 7 2.10 10 5 6.96 10 4 71.17 10 4
growing with the decrease of scattered abrasive particles. Plastic flow results to the adhesive friction and wear for the composite at elevated temperatures. Additionally, on the worn scar of the tribo-
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pair SiC ball at 600 °C covers a mechanical mixed layer as shown in Fig. 9a. The EDS analysis (Fig. 9b) shows that the mechanical layer consists Ti and Cu elements, which indicates that material transference from Ti3SiC2/Cu composite to the counter SiC ball occurred during the rubbing. Plastic adhesive wear and material transference replacing the abrasive wear becomes the dominant wear mechanism at elevated temperatures, which leads to the higher WRs than that at RT. The chemical states of Ti2p (Fig. 10a), Si2p (Fig. 10b), and Cu2p (Fig. 10c) on the worn surface of TSC10 composite at RT and 600 °C were analyzed by XPS and the results are shown in Fig. 10. It can be seen that the Ti, Si and Cu species are partly oxidized to form TiO2, SiO2, and CuO at RT. At 600 °C, tribo-oxidation is more severe. The elements of Ti, Si, and Cu are all completely oxidized. So tribooxidation wear is another wear mechanism for the composites. As is well known, the unique valence bond structure makes
Fig. 8. SEM micrographs of worn surfaces of (a) Ti3SiC2, (b) TSC10 at RT and TSC10 at (c) 200, (d) 400, and (e) 600 °C.
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Fig. 9. (a) The SEM image and (b) the EDS analysis of worn scar of SiC ball rubbing with TSC10 at 600 °C.
Ti3SiC2 be characterized by the unusual combination of metallic and ceramic properties [1]. When the temperature reaches 1200 °C, the brittle-ductile transition will take place for Ti3SiC2. While in the process of friction, the worn surface undergoes ductile under the flash temperature derived from the friction shearing action. On the other hand, it was reported that continuous and homogeneous tribo-oxidation scales formed on the worn surface at high temperature or high sliding speed can endow Ti3SiC2 excellent lubricating and anti-wear properties [31–33]. However, in the present study though tribo-oxidation occur on the worn surface, the oxidation sale is not continuous and homogeneous and can not act as lubrication layer to reduce friction and wear. In conclusion, there exists different level of abrasive wear for Ti3SiC2 and the Ti3SiC2/Cu composites at RT. For Ti3SiC2, there is local plastic deformation besides abrasive wear. For the TSC5 and TSC10 composites, the main wear mechanism is abrasive wear. In a word, the addition of Cu to the polycrystalline Ti3SiC2 effectively prevents the pull-out and fracture of grains and so leads to the decrease of friction and improvement of wear resistance. At elevated temperatures, plastic adhesive wear and tribo-oxidation wear along with material transference become the dominant wear mechanism, which leads to the higher WRs than that at RT.
4. Conclusions Ti3SiC2/Cu composites with different contents of Cu were fabricated and the mechanical and tribological properties of the composites were investigated compared with monolithic Ti3SiC2 in details. The conclusions are as follows: a. The chemical reactions take place to produce Cu3Si and TiCx at the process of fabrication of the composites by SPS and are
Fig. 10. X-ray photoelectron spectroscopy for (a) Ti2p (a), (b) Si2p, and (c) Cu2p on the worn surface of TSC10 composite at RT and 600 °C.
more intense for the composites with high content of Cu (15 and 20 wt%). The reaction products consist of TiCx, Ti5Si3Cy, Cu3Si, and TiSi2Cz. Cu3Si phase distributes around the Ti3SiC2 particles. b. The hardness of Ti3SiC2 composites is much higher than that of monolithic Ti3SiC2, which is derived from the enhancement of
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TiCx, Ti5Si3Cy, and Cu3Si. The flexural strength of the composites is lower than the monolithic Ti3SiC2. The fine impurities of Cu3Si, TiCx, and Ti5Si3Cy partly obstacle multiple energyabsorbing mechanism. Many fine Cu3Si, TiCx, and Ti5Si3Cy grains distributed in the relatively coarse Ti3SiC2 grains effectively prohibit the activated dislocation glide and jogged screw dislocations, which leads to the higher compressive strength for the composites. c. The friction and wear of Ti3SiC2/Cu composites is lower than that of monolithic Ti3SiC2 at RT, which is ascribed to the fixing effect of hard particle of Cu3Si, TiCx, and Ti5Si3Cy to reduce the abrasive friction and wear. While at elevated temperatures plastic adhesive wear and tribo-oxidation wear accompanied by material transference contribute to the increased friction and wear for the composites.
Acknowledgements The authors acknowledge the financial support from the National Natural Science Foundation of China (Grant nos. 51405475, 51475444, and 51505464).
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