Grain Refinement and Mechanical Properties of Cu–Cr–Zr Alloys with Different Nano-Sized TiCp Addition

The TiCp/Cu master alloy was prepared via thermal explosion reaction. Afterwards, the nano-sized TiCp/Cu master alloy was dispersed by electromagnetic stirring casting into the melting Cu–Cr–Zr alloys to fabricate the nano-sized TiCp-reinforced Cu–Cr–Zr composites. Results show that nano-sized TiCp can effectively refine the grain size of Cu–Cr–Zr alloys. The morphologies of grain in Cu–Cr–Zr composites changed from dendritic grain to equiaxed crystal because of the addition and dispersion of nano-sized TiCp. The grain size decreased from 82 to 28 μm with the nano-sized TiCp content. Compared with Cu–Cr–Zr alloys, the ultimate compressive strength (σUCS) and yield strength (σ0.2) of 4 wt% TiCp-reinforced Cu–Cr–Zr composites increased by 6.7% and 9.4%, respectively. The wear resistance of the nano-sized TiCp-reinforced Cu–Cr–Zr composites increased with the increasing nano-sized TiCp content. The wear loss of the nano-sized TiCp-reinforced Cu–Cr–Zr composites decreased with the increasing TiCp content under abrasive particles. The eletrical conductivity of Cu–Cr–Zr alloys, 2% and 4% nano-sized TiCp-reinforced Cu–Cr–Zr composites are 64.71% IACS, 56.77% IACS and 52.93% IACS, respectively.


Introduction
Cu-matrix composites and Cu alloys were widely used as functional and structural materials [1][2][3][4][5][6], such as electrodes for electrical-resistance welding, electric switches, lead frames, friction pieces, as well as a cooling medium of the magnetic channel, encapsulating material on account of their good wear resistance [7], excellent electrical and thermal conductivities [8,9], and good corrosion resistance [10,11] etc. Recent research indicates that the dispersion of ceramic particles in the copper matrix could play a role in pinning dislocation movement and improving the strength of the composites, in addition, alloying elements can also enhance Cu alloy [12][13][14][15]. For example, Lu et al. [15] fabricated 40-60 vol % TiC x -TiB 2 /Cu composites by thermal explosion reaction synthesis. They found that the ultimate compressive strength increased with the ceramic content and the microhardness of the 40-60 vol % TiC x -TiB 2 /Cu composites reached 339, 404 and 448 HV, respectively. However, the electrical and thermal conductivity of Cu was reduced because of the high content of ceramic particles. Zhang et al. [2] revealed that the microhardness of solid soluted Cu-Cr-Zr alloys was 181.8 HV, which was much higher than that of pure copper (68 HV), with less decreasing electrical conductivity (70.8% IACS). Nevertheless, the microhardness is much less than those of particle-reinforced Cu matrix composites. Qiu et al. [12] fabricated the 50 vol % TiC-TiB 2 /Cu composites with different Cr content. Abrasive wear results showed that the volume loss of the 50 vol % TiC-TiB 2 /Cu composites decreased with Cr content and the composites had an extremely low wear rate compared with pure Cu. Afterwards, It is clear from the above that the enhancing principles of PRMMCs are the exposed ceramic playing as barriers, which enhanced the wear resistance of the composites.
In this research, nano-sized TiCp-reinforced Cu-Cr-Zr composites were prepared by the dilution of TiC p /Cu master alloy into Cu-Cr-Zr alloys; the TiC p /Cu master alloy was fabricated by thermal explosion reaction in Cu-Ti-CNTs system. The electrical conductivity and compressive properties of the TiC p reinforced Cu-Cr-Zr composites were studied. Furthermore, the wear behavior and law of the TiC p -reinforced Cu-Cr-Zr composite was studied. This will provide some guidance for the fabrication of TiC p -reinforced Cu-Cr-Zr composites and its application.

Materials and Methods
Cu powder with a particle size of~45 µm, CNTs~15-80 µm in length and~10-20 nm in diameter and Ti powder~25 µm in size were used to fabricate TiC p /Cu master alloys via combustion synthesis. Cu, Ti and CNTs powders were weighed accurately, which was calculated according to the C/Ti mole ratio of 0.8:1. The reactants' compositions was 70% Cu in volume. The prepared powders were milled on a planetary mixer with six cylindrical ceramic inwall containers at the speed of 50 rpm for 24 h. The blended powder was compressed into cylindrical compacts under the pressure of 100 MPa. The size of the compacts was 29 mm in diameter and 45 mm in height. Afterwards, the compacts were placed into a high strength graphite mold, which was later put into a self-made vacuum furnace and heated. W5-Re26 thermocouples were used to measure the temperature. When the temperature rose rapidly, stopped heating and pressed with the pressure of 40 MPa. The nano-sized TiC p /Cu master alloy was obtained when it was cooled down to room temperature.
The Cu-Cr-Zr composites reinforced with different TiC p content were prepared by the dilution of TiC p /Cu master alloy into Cu-Cr-Zr alloys via stirring casting. At first, the prepared Cu-Cr-Zr alloys were heated in a high purity graphite crucible, which was later placed in an induction furnace with a water cooling system. Once the Cu-Cr-Zr alloys melted, the TiC p /Cu master alloy was added in the molten Cu-Cr-Zr. Manual stirring was helpful to melt and disperse the master alloy. Once the master alloy melted completely, kept the temperature for 3 min and the melt were poured into a steel mold. After cooling down to indoor temperature, the TiCp-reinforced Cu-Cr-Zr composites with different TiC p content were successfully prepared.
The characterization of phase constitutions of TiC p /Cu master alloy were performed on a X-ray diffraction (XRD, Rigaku D/Max 2500PC, Tokyo, Japan) with Cu Kα radiation at the scanning speed of 4 • /min. The TiC p in TiC p /Cu master alloy were extracted by FeCl 3 -HCl distilled water solution, the morphologies of TiC p were detected by Field Emission Scanning Electron Microscope (FESEM, JSM 6700, Tokyo, Japan). Olympus optical microscope (XJZ-6, Tokyo, Japan) and high resolution transmission electron microscopy (HRTEM, JEM-2100F, Tokyo, Japan) were used to observe the microstructure of the Cu-Cr-Zr alloys and TiCp-reinforced Cu-Cr-Zr composites. The compressive properties were proceeded on a servo hydraulic materials testing system (MTS, MTS 810, Minneapolis, MN, USA) at a strain rate of 3 × 10 −4 s −1 . Microhardness of the Cu-Cr-Zr alloys and TiCp-reinforced Cu-Cr-Zr composites were performed by a Vickers hardness tester (Model 1600-5122VD, Newage, Feasterville, PA, USA) using a static load of 50 gf and a dwell time of 15 s. Abrasive wear tests were carried out on a pin-on-disk machine with the SiC abrasive papers, the load was 5 N and the wear distance was 24.78 m. The abrasive wear samples are 12 mm in height and 6 mm in diameter. The electrical conductivities of the Cu-Cr-Zr alloys and TiCp-reinforced Cu-Cr-Zr composites were measured at room temperature by a digital eddy current electroconductive machine (Sigma 2008b, Xiamen, China), International Annealed Copper Standard (IACS) are quoted as the units of electrical conductivity results. Figure 1 shows the XRD results of Cu-Ti-CNTs master alloy fabricated by combustion synthesis. The diffraction peaks belonging to TiC p and Cu can be seen clearly, indicating that the in-situ nano-sized TiC p /Cu master alloy was successfully fabricated. Figure 2a shows the morphology of the deep etched surfaces of the nano-sized TiC p /Cu master alloy. As can be seen, Cu was removed and the TiC p in the TiC p /Cu master alloy exhibited a homogeneous distribution. The morphologies of extracted nano-sized TiC p were shown in Figure 2b. As indicated, the morphologies of the extracted TiC p in TiC p /Cu master alloy are spherical, with the size of 100 nm on average (Figure 2c). The interspace between TiC p of the magnified image of the selected area can be observed (Figure 2d).   The microstructures of nano-sized TiCp/Cu-Cr-Zr composites with different TiCp content are shown in Figure 3. As illustrated, dendritic grain is the main morphology of Cu-Cr-Zr alloys with    The microstructures of nano-sized TiCp/Cu-Cr-Zr composites with different TiCp content are shown in Figure 3. As illustrated, dendritic grain is the main morphology of Cu-Cr-Zr alloys with The microstructures of nano-sized TiC p /Cu-Cr-Zr composites with different TiC p content are shown in Figure 3. As illustrated, dendritic grain is the main morphology of Cu-Cr-Zr alloys with the grain size of 82 µm on average; the morphologies of the grain in 2 wt % and 4 wt % nano-sized TiCp-reinforced Cu-Cr-Zr composites are equiaxed crystal and the grain size are 36 µm and 28 µm on average, respectively. The morphology of the grain changes from coarse dendrites to equiaxed crystal with increasing TiC p content. An investigation [45] suggested that TiC p were not the effective substrates for the heterogeneous nucleation of the Cu-melt. However, the variation in grain from dendrites to equiaxed crystal was a piece of adverse evidence for the ineffectiveness of TiC p in grain refinement in Cu alloy.

Results and Discussion
As we know, the metal matrix determined the ductility of the composites and the ceramic particles determined the strength of the composites [46]. The fracture surface of (a) Cu-Cr-Zr alloys, (b) 2-wt % nano-sized TiC p /Cu-Cr-Zr composites and (c) 4-wt % nano-sized TiC p /Cu-Cr-Zr composites were shown in Figure 5. As can be seen, the fracture surface of Cu-Cr-Zr alloys consists of dimples, which means that the Cu-Cr-Zr alloy has a good plasticity. With the addition of TiC p , the amount of dimples decreased and some TiCp clusters appeared in the dimples. Uniformly dispersed TiC p can improve the mechanical properties of the Cu-Cr-Zr alloys, but cluster of TiC p will act as the source of crack leading to the reduction of the plasticity. This means that the additon of TiC p can improve the strength of composite but sacrifice the plasticity. According to the above viewpoint, we conject that the σ 0.2 , σ UCS and microhardness could be enhanced and the fracture strain (ε f ) decreased with the increasing TiC p content. Consequently, the 4-wt % TiCp-reinforced Cu-Cr-Zr composites possessed the highest compression strength and microhardness (HV). the grain size of 82 μm on average; the morphologies of the grain in 2 wt % and 4 wt % nano-sized TiCp-reinforced Cu-Cr-Zr composites are equiaxed crystal and the grain size are 36 μm and 28 μm on average, respectively. The morphology of the grain changes from coarse dendrites to equiaxed crystal with increasing TiCp content. An investigation [45] suggested that TiCp were not the effective substrates for the heterogeneous nucleation of the Cu-melt. However, the variation in grain from dendrites to equiaxed crystal was a piece of adverse evidence for the ineffectiveness of TiCp in grain refinement in Cu alloy.  Figure 4 shows the engineering stress-strain curves of compressive test for Cu-Cr-Zr alloys and nano-sized TiCp-reinforced Cu-Cr-Zr composites with different TiCp content. The compressive test data are presented in Table 1. As exhibited in the diagram, the yield strength (σ0.2), ultimate compressive strength (σUCS) and microhardness (HV) of Cu-Cr-Zr alloys were obviously enhanced at the expense of the fracture strain (εf) because of the addition of TiCp. For the nano-sized TiCpreinforced Cu-Cr-Zr composites, the σ0.2, σUCS and microhardness increased with the increasing TiCp content. The 4-wt% TiCp-reinforced Cu-Cr-Zr composites possessed the highest yield strength (σ0.2), ultimate compressive strength (σUCS) and microhardness (HV), which are 190 MPa, 491 MPa and 118.5 HV, respectively. The fracture strain (εf) of Cu-Cr-Zr alloys, 2-wt % and 4-wt % nano-sized TiCpreinforced Cu-Cr-Zr composites are 34.1%, 30.7% and 29.3%, respectively.
As we know, the metal matrix determined the ductility of the composites and the ceramic particles determined the strength of the composites [46]. The fracture surface of (a) Cu-Cr-Zr alloys, (b) 2-wt % nano-sized TiCp/Cu-Cr-Zr composites and (c) 4-wt % nano-sized TiCp/Cu-Cr-Zr composites were shown in Figure 5. As can be seen, the fracture surface of Cu-Cr-Zr alloys consists  Figure 4 shows the engineering stress-strain curves of compressive test for Cu-Cr-Zr alloys and nano-sized TiCp-reinforced Cu-Cr-Zr composites with different TiC p content. The compressive test data are presented in Table 1. As exhibited in the diagram, the yield strength (σ 0.2 ), ultimate compressive strength (σ UCS ) and microhardness (HV) of Cu-Cr-Zr alloys were obviously enhanced at the expense of the fracture strain (ε f ) because of the addition of TiC p . For the nano-sized TiCp-reinforced Cu-Cr-Zr composites, the σ 0.2 , σ UCS and microhardness increased with the increasing TiC p content.   The variation in volume loss with different abrasive particles of 6.5 μm, 10.3 μm and 13 μm at the applied load of 5 N for the Cu-Cr-Zr alloys, 2-wt % and 4-wt % nano-sized TiCp-reinforced Cu-Cr-Zr composites are shown in Figure 6. As indicated, the volume loss of the alloys and composites increased with the increasing abrasive particles. This may be due to the large abrasive particles that penetrated deeply into contact area of the alloys and the composites, leading to the increase in volume loss of composites. Meanwhile, the volume loss of the Cu-Cr-Zr alloys, 2-wt % and 4-wt % TiCpreinforced Cu-Cr-Zr composites decreased with the increasing TiCp content under all abrasive particles. The reason may be the grain refinement, nano-sized particle strengthening and the good interface bonding of composites, which could enhance the hardness and decrease the volume loss of composites. It also can be confirmed by the SEM images of the worn surfaces for Cu-Cr-Zr alloys, 2wt % and 4-wt % nano-sized TiCp-reinforced Cu-Cr-Zr composites at 13 μm abrasive particles, as shown in Figure 7.     The variation in volume loss with different abrasive particles of 6.5 μm, 10.3 μm and 13 μm at the applied load of 5 N for the Cu-Cr-Zr alloys, 2-wt % and 4-wt % nano-sized TiCp-reinforced Cu-Cr-Zr composites are shown in Figure 6. As indicated, the volume loss of the alloys and composites increased with the increasing abrasive particles. This may be due to the large abrasive particles that penetrated deeply into contact area of the alloys and the composites, leading to the increase in volume loss of composites. Meanwhile, the volume loss of the Cu-Cr-Zr alloys, 2-wt % and 4-wt % TiCpreinforced Cu-Cr-Zr composites decreased with the increasing TiCp content under all abrasive particles. The reason may be the grain refinement, nano-sized particle strengthening and the good interface bonding of composites, which could enhance the hardness and decrease the volume loss of composites. It also can be confirmed by the SEM images of the worn surfaces for Cu-Cr-Zr alloys, 2wt % and 4-wt % nano-sized TiCp-reinforced Cu-Cr-Zr composites at 13 μm abrasive particles, as shown in Figure 7. The variation in volume loss with different abrasive particles of 6.5 µm, 10.3 µm and 13 µm at the applied load of 5 N for the Cu-Cr-Zr alloys, 2-wt % and 4-wt % nano-sized TiCp-reinforced Cu-Cr-Zr composites are shown in Figure 6. As indicated, the volume loss of the alloys and composites increased with the increasing abrasive particles. This may be due to the large abrasive particles that penetrated deeply into contact area of the alloys and the composites, leading to the increase in volume loss of composites. Meanwhile, the volume loss of the Cu-Cr-Zr alloys, 2-wt % and 4-wt % TiCp-reinforced Cu-Cr-Zr composites decreased with the increasing TiC p content under all abrasive particles. The reason may be the grain refinement, nano-sized particle strengthening and the good interface bonding of composites, which could enhance the hardness and decrease the volume loss of composites. It also can be confirmed by the SEM images of the worn surfaces for Cu-Cr-Zr alloys, 2-wt % and 4-wt % nano-sized TiCp-reinforced Cu-Cr-Zr composites at 13 µm abrasive particles, as shown in Figure 7.
SiC particles penetrated deeply into the contact area of Cu-Cr-Zr alloys and these abrasive particles (Figure 7a) lead to the deformation of the surfaces of the alloys. The surfaces, after wear and tear of TiCp-reinforced Cu-Cr-Zr composites, became smooth with the increasing nano-sized TiC p contents (Figure 7b,c). As already discussed, microhardness of the TiCp-reinforced Cu-Cr-Zr composites can be improved with an increase in particle content. Therefore, the depth of abrasive particles penetrated into the composites decreased with the increasing hardness. Accordingly, it is evident that TiC p can effectively enhance the microhardness and wear resistance by dispersing TiC p into Cu matrix, in which TiC p played a role as grain refinement in enhancing the composites and a barrier in reducing the cutting efficiency of abrasive particles as well as hammering the deformation of the Cu matrix. Accordingly, the wear resistance of the TiCp-reinforced Cu-Cr-Zr composites was enhanced due to the addition of TiC p . SiC particles penetrated deeply into the contact area of Cu-Cr-Zr alloys and these abrasive particles (Figure 7a) lead to the deformation of the surfaces of the alloys. The surfaces, after wear and tear of TiCp-reinforced Cu-Cr-Zr composites, became smooth with the increasing nano-sized TiCp contents (Figure 7b,c). As already discussed, microhardness of the TiCp-reinforced Cu-Cr-Zr composites can be improved with an increase in particle content. Therefore, the depth of abrasive particles penetrated into the composites decreased with the increasing hardness. Accordingly, it is evident that TiCp can effectively enhance the microhardness and wear resistance by dispersing TiCp into Cu matrix, in which TiCp played a role as grain refinement in enhancing the composites and a barrier in reducing the cutting efficiency of abrasive particles as well as hammering the deformation of the Cu matrix. Accordingly, the wear resistance of the TiCp-reinforced Cu-Cr-Zr composites was enhanced due to the addition of TiCp.    SiC particles penetrated deeply into the contact area of Cu-Cr-Zr alloys and these abrasive particles (Figure 7a) lead to the deformation of the surfaces of the alloys. The surfaces, after wear and tear of TiCp-reinforced Cu-Cr-Zr composites, became smooth with the increasing nano-sized TiCp contents (Figure 7b,c). As already discussed, microhardness of the TiCp-reinforced Cu-Cr-Zr composites can be improved with an increase in particle content. Therefore, the depth of abrasive particles penetrated into the composites decreased with the increasing hardness. Accordingly, it is evident that TiCp can effectively enhance the microhardness and wear resistance by dispersing TiCp into Cu matrix, in which TiCp played a role as grain refinement in enhancing the composites and a barrier in reducing the cutting efficiency of abrasive particles as well as hammering the deformation of the Cu matrix. Accordingly, the wear resistance of the TiCp-reinforced Cu-Cr-Zr composites was enhanced due to the addition of TiCp.     Wetting angle (θ) was a standard to measure whether the reinforced particles can be act as the heterogeneous nuclei to refine the grain during the crystallization of Cu melt. Small wetting angle (0° < θ < 90°) can provide the potential to act as the heterogeneous nuclei and high wetting angle (90° < θ < 180°) can not. TiCp and Cu have a high wetting angle that belongs to non-wetting system. Accordingly, the TiCp could not act as the heterogeneous nuclei to refine the grain of Cu-Cr-Zr alloy. For the purpose of studying the grain refining mechanism of Cu-Cr-Zr alloys, we therefore studied the interfaces between TiCp and Cu nmatrix in the TiCp-reinforced Cu-Cr-Zr composites. Figure 9 shows TEM images of 4 wt % nano-sized TiCp-reinforced Cu-Cr-Zr composites. As indicated in Figure 9a [47] confirmed that the free Ti in the melt was conducive to the formation of Cu3Ti layer. Moreover, Yang et al. [48] indicated that the wear resistance could be improved for the good interface bonding of reinforced particles and the matrix in the composites.
Lattice misfit (δ) used to judge the possibility of the heterogeneous nuclei can be calculated by the flowing mathematical model [49]:  Wetting angle (θ) was a standard to measure whether the reinforced particles can be act as the heterogeneous nuclei to refine the grain during the crystallization of Cu melt. Small wetting angle (0 • < θ < 90 • ) can provide the potential to act as the heterogeneous nuclei and high wetting angle (90 • < θ < 180 • ) can not. TiC p and Cu have a high wetting angle that belongs to non-wetting system. Accordingly, the TiC p could not act as the heterogeneous nuclei to refine the grain of Cu-Cr-Zr alloy. For the purpose of studying the grain refining mechanism of Cu-Cr-Zr alloys, we therefore studied the interfaces between TiC p and Cu nmatrix in the TiCp-reinforced Cu-Cr-Zr composites. Figure 9 shows TEM images of 4 wt % nano-sized TiCp-reinforced Cu-Cr-Zr composites. As indicated in Figure 9a [47] confirmed that the free Ti in the melt was conducive to the formation of Cu 3 Ti layer. Moreover, Yang et al. [48] indicated that the wear resistance could be improved for the good interface bonding of reinforced particles and the matrix in the composites.
Lattice misfit (δ) used to judge the possibility of the heterogeneous nuclei can be calculated by the flowing mathematical model [49]:   Table 2 shows the calculated results of lattice misfit using Equation (1). As shown in Table 2, the lattice misfit (δ) between the (111) of Cu and the (010) of Cu3Ti is 1.9%, indicating that Cu3Ti could act as the heterogeneous nuclei during the crystallization of Cu-melt. On the other hand, the lattice misfit (δ) between the (100) of Cu3Ti and the (100) of TiC is 1.2%. The low lattice misfit (δ) between Cu3Ti and Cu as well as Cu3Ti and TiCp is conducive to good interface bonding. According to the reaction mechanism in Cu-Ti-CNTs system [25], Cu firstly reacts with Ti to form TixCuy around Ti particles. With increasing temperature, TixCuy melts to form a Cu-Ti binary liquid phase, meanwhile, CNTs dissolve in the Cu-Ti binary liquid phase to form an Cu-Ti-C ternary liquid phase. When the concentration of [Ti] and [C] in the Cu-Ti-C ternary liquid phase is high enough for reactions between [Ti] and [C] to occur, TiCx will be synthesized and precipitate out of the melts. However, the reaction is not complete; TixCuy will be remaining in the TiCx/Cu master alloy. When the master alloy remelts and is dispersed into the melting Cu-Cr-Zr alloys, the Cu3Ti layer formed by the adsorption of Ti atoms from Cu-Ti solution, and the Cu3Ti could even exist at lower concentration or at the temperature above the alloy liquidus. Therefore, the formation of Cu3Ti is mainly because of the reduction of the interfacial energy at the interface between TiCp and the Cu-melt. As a result, Cu3Ti formed surrounding the TiCp as the heterogeneous nuclei refined the grain.
As already discussed, the grain refinement of Cu-Cr-Zr alloy is caused by the Cu3Ti, which formed on the surface of TiCp. TiCp with a Cu3Ti layer could be used as the heterogeneous nuclei of the Cu-melt. The refinement of grain, nano-sized particle strengthening and good interface bonding improved the ultimate compressive strength (σUCS), yield strength (σ0.2), microhardness (HV) and wear resistance of different TiCp-content reinforced Cu-Cr-Zr composites with a slight decrease in fracture strain (εf) and electrical conductivity.  Table 2 shows the calculated results of lattice misfit using Equation (1). As shown in Table 2, the lattice misfit (δ) between the (111) of Cu and the (010) of Cu 3 Ti is 1.9%, indicating that Cu 3 Ti could act as the heterogeneous nuclei during the crystallization of Cu-melt. On the other hand, the lattice misfit (δ) between the (100) of Cu 3 Ti and the (100) of TiC is 1.2%. The low lattice misfit (δ) between Cu 3 Ti and Cu as well as Cu 3 Ti and TiCp is conducive to good interface bonding. According to the reaction mechanism in Cu-Ti-CNTs system [25], Cu firstly reacts with Ti to form Ti x Cu y around Ti particles. With increasing temperature, Ti x Cu y melts to form a Cu-Ti binary liquid phase, meanwhile, CNTs dissolve in the Cu-Ti binary liquid phase to form an Cu-Ti-C ternary liquid phase. When the concentration of [Ti] and [C] in the Cu-Ti-C ternary liquid phase is high enough for reactions between [Ti] and [C] to occur, TiC x will be synthesized and precipitate out of the melts. However, the reaction is not complete; Ti x Cu y will be remaining in the TiC x /Cu master alloy. When the master alloy remelts and is dispersed into the melting Cu-Cr-Zr alloys, the Cu 3 Ti layer formed by the adsorption of Ti atoms from Cu-Ti solution, and the Cu 3 Ti could even exist at lower concentration or at the temperature above the alloy liquidus. Therefore, the formation of Cu 3 Ti is mainly because of the reduction of the interfacial energy at the interface between TiCp and the Cu-melt. As a result, Cu 3 Ti formed surrounding the TiCp as the heterogeneous nuclei refined the grain.
As already discussed, the grain refinement of Cu-Cr-Zr alloy is caused by the Cu 3 Ti, which formed on the surface of TiCp. TiCp with a Cu 3 Ti layer could be used as the heterogeneous nuclei of the Cu-melt. The refinement of grain, nano-sized particle strengthening and good interface bonding improved the ultimate compressive strength (σ UCS ), yield strength (σ 0.2 ), microhardness (HV) and wear resistance of different TiCp-content reinforced Cu-Cr-Zr composites with a slight decrease in fracture strain (ε f ) and electrical conductivity.  [010]TiC 0

Conclusions
The compressive properties, wear resistance and electrical conductivities of Cu-Cr-Zr alloys and TiCp-reinforced Cu-Cr-Zr composites fabricated via thermal explosion combined with stirring casting were investigated. The results show that nano-sized TiCp can effectively refine the grain of Cu-Cr-Zr alloys. Nano-sized TiCp was surrounded by a layer of Cu3Ti, which played a role as heterogeneous nuclei of Cu. The microstructure of Cu-Cr-Zr composites changed from dendritic grain to equiaxed crystal with the addition of TiCp/Cu master alloy. The grain size decreased from 82 to 28 μm with the increasing TiCp content. Compared with Cu-Cr-Zr alloys, the ultimate compressive strength (σUCS), yield strength (σ0.2) and hardness (HV) of nano-sized TiCp-reinforced Cu-Cr-Zr composites were improved. The wear resistance of Cu-Cr-Zr alloys and TiCp-reinforced Cu-Cr-Zr composites increased with the increasing content of TiCp. The refinement of grain, nano-sized particle strengthening and the good interface bonding improved the strength, hardness and wear resistance of Cu-Cr-Zr alloys and nano-sized TiCp-reinforced Cu-Cr-Zr composites with different particle content.

Conclusions
The compressive properties, wear resistance and electrical conductivities of Cu-Cr-Zr alloys and TiCp-reinforced Cu-Cr-Zr composites fabricated via thermal explosion combined with stirring casting were investigated. The results show that nano-sized TiC p can effectively refine the grain of Cu-Cr-Zr alloys. Nano-sized TiC p was surrounded by a layer of Cu 3 Ti, which played a role as heterogeneous nuclei of Cu. The microstructure of Cu-Cr-Zr composites changed from dendritic grain to equiaxed crystal with the addition of TiC p /Cu master alloy. The grain size decreased from 82 to 28 µm with the increasing TiC p content. Compared with Cu-Cr-Zr alloys, the ultimate compressive strength (σ UCS ), yield strength (σ 0.2 ) and hardness (HV) of nano-sized TiCp-reinforced Cu-Cr-Zr composites were improved. The wear resistance of Cu-Cr-Zr alloys and TiCp-reinforced Cu-Cr-Zr composites increased with the increasing content of TiCp. The refinement of grain, nano-sized particle strengthening and the good interface bonding improved the strength, hardness and wear resistance of Cu-Cr-Zr alloys and nano-sized TiCp-reinforced Cu-Cr-Zr composites with different particle content.