Effect of carbide interlayers on the microstructure and properties of graphene-nanoplatelet-reinforced copper matrix composites

Materials Science & Engineering A 708 (2017) 311–318

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Effect of carbide interlayers on the microstructure and properties of graphene-nanoplatelet-reinforced copper matrix composites

MARK

Xiaoyang Sia,b,1, Mian Lia,c,1, Fanyan Chena, Per Eklundd, Jianming Xuee, Feng Huanga, ⁎ Shiyu Dua, Qing Huanga, a Engineering Laboratory of Specialty Fibers and Nuclear Energy Materials (FiNE), Ningbo Institute of Industrial Technology, Chinese Academy of Sciences, Ningbo, Zhejiang 315201, PR China b School of Materials Science and Engineering, Shanghai University, Shanghai 200072, PR China c University of Chinese Academy of Sciences, 19A Yuquan Rd, Shijingshan District, Beijing 10049, PR China d Thin Film Physics Division, Dept. of Physics, Chemistry, and Biology (IFM), Linköping University, 581 83 Linköping, Sweden e State Key Laboratory of Nuclear Physics and Technology, CAPT and IFSA Collaborative Innovation Center of MoE, Peking University, Beijing 100871, PR China

A R T I C L E I N F O

A B S T R A C T

Keywords: Graphene-nanoplatelets Interlayers Interfacial bonding Metal matrix composites

Copper matrix composites reinforced with carbide-coated graphene nanoplatelets (GNPs) were investigated in order to understand the role of the interlayers on the thermal, electrical, mechanical and electro-tribological properties of the composites. The TiC or VC coatings were formed in situ on the two sides of GNPs through a controllable reaction in molten salts. Compared with bare GNPs composites, the bonding between the GNPs and copper was improved. Accordingly, the tensile strength and the fracture elongation of Cu/GNPs composites with an interlayer were enhanced by strengthened interfacial bonding. Furthermore, the wear resistance of Cu/GNPs composites was remarkably improved.

1. Introduction Graphene has properties such as extraordinary intrinsic mobility [1], high thermal conductivity [2], and high Young's modulus and fracture strength [3,4]. Graphene as a reinforcement, or filler phase, in composites has been extensively studied in polymer and ceramic matrix composites to enhance thermal, electrical and mechanical properties [5–10]. However, it is challenging to incorporate graphene in metal matrix composites, especially in copper matrix composites. In addition, the heat conduction and mechanical properties of graphene-reinforced copper matrix composites (Cu/Gr) are typically not satisfactory [11,12]. According to our previous research, with increasing concentration of graphene-nanoplatelets (GNPs) in copper matrix composites, the thermal conductivity distinctly decreased, while the tensile strength increased for low graphene content but decreased drastically with graphene content of more than 0.6 vol% [13]. One of the key issues that caused the deteriorated properties of Cu/ Gr composites was the poor interfacial properties. For metal matrix composites (MMCs), good interfacial bonding is needed to ensure heat conduction and load transfer. Nevertheless, carbon materials generally exhibit poor wettability with Cu; for example, the contact angle

between Cu and graphite is typically 140° [14]. Along with the large difference in the coefficient of thermal expansion (CTE) between graphene and copper, these composites often exhibit weak interfacial bonding. On one hand, the weak interface may not withstand the load transfer from matrix to filler, resulting in poor mechanical properties. On the other hand, the poor interface bonding results in high thermal resistance, i.e., a lower thermal conductivity of the composites. In order to improve the interfacial bonding of MMCs reinforced by carbon materials, a feasible way is to add an interlayer between reinforcement and matrix. The early transition metal carbides, with their mixture of metallic, covalent and ionic bonding, are appropriate candidates for the interlayer materials. Chu et al. [15] prepared Cr3C2 as the interlayer of carbon-nanotubes (CNT)-reinforced copper matrix composites. The results showed that a Cr3C2 interlayer could significantly enhance the interfacial bonding and improve mechanical properties of the composites. Liu et al. [16,17] reported copper matrix composites reinforced with TiC- and Mo2C- coated graphite fibers, resulting in improved interfacial bonding and thermal conductivity. Zhang et al. [18] fabricated copper matrix composites reinforced by TiC-coated diamond particles. Similarly, the thermal conductivity of the composites reinforced with coated diamond was higher than that of the

Abbreviations: Gr, graphene; GNPs, graphene-nanoplatelets; MMCs, metal matrix composites; CTE, coefficient of thermal expansion ⁎ Corresponding author. E-mail address: [email protected] (Q. Huang). 1 These authors contributed equally. http://dx.doi.org/10.1016/j.msea.2017.10.015 Received 25 August 2017; Received in revised form 3 October 2017; Accepted 4 October 2017 Available online 05 October 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.

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2. Experimental section

mixed solution to form a Cu(OH)2/GNPs slurry; (4) keeping the slurry at 60 °C for 4 h to sufficiently precipitate GNPs/CuO2 composites; (5) washing, drying and reducing the composites under an H2 atmosphere to obtain Cu/GNPs composites. The Cu/coated GNPs composites were prepared by the same method. The bulk composites were fabricated by Spark plasma sintering (SPS, HP D25/1, FCT Systeme GmbH, Rauenstein, Germany) with the following steps. Firstly, the powder composites were put into a cylindrical graphite die with an inner diameter of 13 mm (for thermal and electrical test), 20 mm (for tensile tests), and 40 mm (for electrical friction and wear tests), respectively. Then, the powder composites were pressed by an initial pressure of 5 MPa to form a disc. Subsequently, the composites were sintered at 700 °C for 5 min (heating rate of 50 °C/min). Before cooling, a pressure of 35 MPa was kept. The pressure was increased to 40 MPa during the whole cooling stage (cooling rate of 50 °C/min).

2.1. Materials

2.4. Characterization

GNPs from Ningbo Morsh Technology Co. Ltd were used as raw carbon source materials. Titanium powder (particle size: ~ 1 µm, purity: 99.9 wt%, Shanghai Pantian Nano Materials Co. Ltd, China) and vanadium powder (particle size: ~ 45 µm, purity: 99.9 wt%, General Research Institute for Nonferrous Metals, China) were used as metal powders. Other chemicals were all purchased from Aladdin Industrial Corporation at analytical grade.

An X-ray diffractometer (XRD, Bruker D8 Advance Davinci) in θ–2θ geometry was utilized to analyze the phases of the coating with Cu Kα radiation, wavelength 1.54 Å, with a step of 0.02° and a collection time per step of 1 s. Field emission scanning electron microscope (SEM, FEI Quanta FEG 250) and transmission electron microscope (TEM, JEOL2100 HR) were used to characterize morphology and microstructure of samples. The TEM samples of the powders were prepared by dropping diluted GNPs@Ti and GNPs@V solution onto a carboncoated copper grid and drying in air. The TEM sample of bulk composites with a diameter of 3 mm and a thickness of 20–30 µm were prepared by mechanical stamping and polishing. Subsequently, these samples were thinned by an argon ion milling (Gatan 691) to obtain a transparent thin area. The change of GNPs after molten salt treatment was examined by Raman spectroscopy (Renishaw inVia Reflex) with an excitation wavelength of 532 nm. Tensile tests were conducted on the universal material tester (Instron 5569A). The bulk density of the composites was measured by Archimedes' method. The theoretical densities of pure copper (8.96 g/cm3) and GNPs (2.2 g/cm3, adopting the theoretical density of graphite) were used to calculate the relative density of the samples. The thermal diffusivity was performed on a Netzsch 457 laser thermal analyzer. The electrical conductivity was conducted by an eddy current conductivity meter (FIRST, FD-102). To investigate the electro-tribological properties, electrical friction and wear tests were conducted on brush-rotor type friction and wear test equipment in air [19]. The test environment was relative humidity of 40–60% and temperature of 15–25 °C. The rotor rings used in the experiment was pure Cu. All samples were set as the positive brush materials with an AC current density of 5 × 104 A/m2 and a pressure of 2.5 × 10−2 MPa during test. The sliding time was 5 h (continuous) at a speed of 5 m/s. The contact voltage drop of samples was measured by a four-point contact system. Based on the energy conservation, the friction coefficient is calculated by

composites reinforced with uncoated diamond. Based on the above research, we propose that synthesizing coatings of early transition metal carbides on graphene could be an effective way to optimize the interfacial properties of graphene-reinforced copper matrix composites. Graphene sheets are prone to agglomerate due to the van der Waals forces. Therefore, it is difficult to fabricate uniform and continuous coatings on graphene. So far, there is no report on producing early transition metal carbides coatings on graphene. In the present work, TiC or VC coatings were in-situ synthesized on graphenenanoplatelets (GNPs) through molten-salt treatment. To demonstrate the effect of the as-obtained coating, copper matrix composites reinforced with bare GNPs and coated GNPs were fabricated and the microstructure, thermal, electrical, mechanical and electro-tribological properties of the bulk composites were investigated.

2.2. Preparation of GNP coatings The coating processing was performed as follows. Firstly, 1 g GNPs were dispersed in 200 mL ethanol in an ultrasonic bath for 1 h. Then, the suspension slurry was mixed with NaCl-KCl salts (1:1, mol%) or LiCl-KCl salts (1:1, mol%) through high-energy planetary ball milling. The high-energy planetary ball milling ran at the speed of 250 r/min for 12 h. NaCl-KCl molten salts were applied to produce TiC-coated GNPs. LiCl-KCl molten salts were used to produce VC-coated GNPs. After drying and grinding, a uniform mixture was achieved. Subsequently, these materials were mixed with pure metal powders and placed in an alumina crucible. The whole mixture was heated for 1 h at 850 °C and for 6 h at 750 °C to obtain TiC coated GNPs and VC coated GNPs, respectively. All reactions were performed in a quartz tube furnace with flowing argon gas. The reaction products were washed repeatedly in warm deionized water to remove the residual salts, and then dried for 24 h at 45 °C in a vacuum drying oven. The molar ratio of titanium and GNPs of the mixture was 1/5, 1/10 and 1/20, respectively. The molar ratio of vanadium and GNPs was 1/10. The samples were assigned sample IDs as in Table 1. 2.3. Preparation of Cu composites The process for synthesis of Cu composites powders is based on our previous research [13] and summarized as follows: (1) dispersing raw GNPs in an alcoholic solution of Cu ion in an ultrasonic bath for 1 h; (2) mixing C6H12O6 as reductant in the suspending solution in an ultrasonic bath for 0.5 h; (3) slowly stirring while adding NaOH solution into the

W-W0 =f⋅S↔(P-P0)⋅t=NB⋅μ⋅Fr⋅ω⋅rc⋅t μ=

GNPs@Ti-5 GNPs@Ti-10 GNPs@Ti-20 GNPs@V-10

Composition (molar ratio)

Reaction condition

NaCl/LiCl

KCl

GNPs

Ti/V

1 1 1 1

1 1 1 1

1 1 1 1

1/5 1/10 1/20 1/10

(2)

where P is the actual power of the motor during the friction test; P0 is the light run power of the motor; NB is the number of brushes; Fr is the normal load on the brush; rc is the radius of the ring; ω is the angular speed of ring. The wear rate of the composites is given by

Table 1 The sample IDs and the corresponding reactant composition. Sample ID

P-P0 NB⋅Fr⋅ω⋅rc

(1)

∆M M0 -M = ρ⋅Fr⋅S ρ⋅Fr⋅S

850 °C, 1 h

WV=

750 °C, 6 h

where WV is the wear rate of brush; ∆M is the mass loss after the wear test; M0 is the mass of the brush before the wear test; M is the mass of 312

(3)

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Fig. 1. XRD patterns and Raman spectra of the reaction products; (a), (b) products of GNPs@Ti at 850 °C for 1 h in the NaCl-KCl molten-salt system, (c), (d) products of GNPs@V at 750 °C for 6 h in the LiCl-KCl molten-salt system.

this phenomenon probably is related to the nucleation and growth behavior of carbide. The amounts of defects decreased with increasing amount of TiC or VC crystals. It is worth noting that the G and 2D peak of GNPs@Ti displayed a red-shift that became larger with the increasing Ti/C ratio. As can be seen in Fig. 2c, the G and 2D peak of GNPs@Ti-5 showed the largest red-shifts by 11.56 cm−1 and 26.49 cm−1, respectively. The shift could be attributed to stress arising from the lattice mismatch of interfaces. The peak shift is usually stronger for the 2D peak than the G peak in Raman spectroscopy [26–28]. Tensile stress can be introduced by deposited a thin silicon layer on the graphene [22], giving a large lattice mismatch resulting in internal tensile stress causing the elongation of the carbon-carbon bonds. As a result, the vibration frequency of carbon-carbon bonds decreased and the G and 2D peaks of GNPs@Ti showed a red-shift. Fig. 2 shows SEM images, TEM images, and SAED patterns of the composites. As shown in Fig. 2a, the raw GNPs with clean and smooth surfaces exhibit a flexible and wrinkled structure. The VC coating (Fig. 2b) was discontinuous with most of the particles distributed on the edges of the GNPs. For GNPs@Ti-20, the graphene sheets were still flexible and transparent with some particles on the surface (Fig. 2c). For GNPs@Ti-10, due to the formation of continuous and thicker coating, the surface of GNPs became relatively rough (Fig. 2d). For GNPs@Ti-5, the surface of the samples is rougher, and the coating appeared to be fully continuous and dense. The samples were further investigated by TEM. Fig. 2f–h are typical TEM images of GNPs@V-10, GNPs@Ti-20 and GNPs@Ti-5, respectively. It is clear that the coating is composed of many nanocrystalline grains. The inset SAED patterns display diffraction rings indexed to the (111) (200), (220) and (222) planes of cubic VC or TiC, which further verified the nanocrystalline structure of the VC or TiC coatings. However, VC grains in the vicinity of defects were much larger, as shown in

the brush after the wear test; ρ is the density of the brush; and S is the sliding distance of the brush. 3. Results and discussion 3.1. Microstructure of carbide coating on GNPs Fig. 1a and b are the XRD patterns of GNPs@Ti and GNPs@V, respectively. By comparing raw GNPs and as-obtained GNPs, it can be seen that the diffraction peaks of TiC or VC phases appeared after the molten-salt reaction. The peaks corresponding to TiC increased in intensity with increasing Ti/C ratio of the starting powders, indicating higher amount of TiC. The graphite peaks of all as-obtained GNPs were strong and sharp. No residual metal was detected. The prominent features of graphene in the Raman spectra are the socalled G, D and 2D peaks, which are located at around 1580 cm−1, 1350 cm−1 and 2700 cm−1, respectively [20]. The Raman spectra in Fig. 1c and d show the presence of graphene in the as-obtained GNPs. The G peak originates from the bonding stretching vibration of the sp2 carbon atoms. The D peak originates from the double resonance Raman process and is activated by defects (such as substitutional hetero-atoms, vacancies, dislocation, and dangling bonds) [21–23]. Therefore, the D peak could not be observed in the Raman spectrum of defect-free graphene. The intensity ratio of D peak to G peak (ID/IG) is usually used to quantify the disorder of graphene [24]. The 2D peak is the second order of the D peak and closely related to the band structure of graphene layers [25]. It can be observed that the ratio of ID/IG decreased for GNPs@Ti and GNPa@V, which indicates that the number of defects in GNPs was largely reduced. The ratio of ID/IG declined from 0.62 to 0.14 as the Ti/C ratio increased to 1/5 and the ratio reduced to 0.22 for GNPs@V-10. Combined with the XRD analysis, it can be concluded that 313

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Fig. 2. SEM images of (a) raw GNPs, (b) GNPs@V-10, (c) GNPs@Ti-20, (d) GNPs@Ti-10, (e) GNPs@Ti-5; TEM images and SAED of (a) GNPs@V-10, (b) GNPs@Ti-20 and (c) GNPs@Ti-5.

3.2. Microstructure of composites

Fig. 2f. Comparing Fig. 2g and h, it is clear that the coating became thicker with increasing Ti content. The formation of the carbide coating can be described as a dissolution-diffusion-reaction mechanism. When the ambient temperature exceeds the melting point of the eutectic salts, the salts melt and the metal powders are dissolved in the molten salts. Afterwards, driven by the chemical potential, the metal cations diffuse to the surface of the GNPs and react with the GNPs. Taking advantage of the liquid-ion medium, the diffusion speed of the metal cations is high. Thus the reaction can take place uniformly at the surface of GNPs. As a result, uniform and nanocrystalline carbide coating are formed (Scheme 1).

Interfacial debonding is a common phenomenon in Cu/Gr composites owning to weak adhesion [29]. In order to determine the effect of the introduced carbide interlayer on interfacial properties, raw GNPs, GNPs@Ti-5 and GNPs@V-10 were used to produce composites. The GNPs content of all composites was 0.5 wt%. For GNPs@Ti-5 and GNPs@V-10, the contents of carbide and residual GNPs were qualified by TG/DTA and TG/DSC, respectively. For the morphology of powder composites, the distribution of GNPs with and without coating is no evident difference in the composites. Fig. 3 shows SEM images of the fracture morphology of composites. The three kinds of composites exhibit different fracture characteristics. Dimples can be seen only to a limited extent in the Cu/GNPs composite Scheme 1. Schematic illustration of GNPs with carbide coating; (a) the ‘sandwiched’ structure of GNPs with carbide coating, (b) coating variation with increased of transition metal element content.

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Fig. 3. SEM images of the fracture morphology of (a) Cu/GNPs, (b) Cu/GNPs@TiC, and (c) Cu/GNPs@VC; (d) (e) and (f) magnified images of (a) (b) and (c), respectively.

interface to below 90° [31]. Moreover, some TiC particles were embedded into the copper matrix, resulting in a strong mechanical bonding between TiC interlayer and copper matrix. Fig. 4b is a TEM image of the interfacial morphology of Cu/GNPs@VC composite. Similar to TiC, the VC coating was bonded with GNPs and Cu, which was beneficial to improve the interfacial bonding of Cu/GNPs composites. The pore at the interface between GNPs and Cu may originate from the preparation process of the sample. Thus, the TiC or VC interlayers changed the “GNPs-gap-Cu” interface structure to “GNPs-carbide-Cu” interface structure, which contributed to enhance the performance of Cu/GNPs composites.

(Fig. 3a), while both the amount and the size of the dimples are much larger in the coated composites (Fig. 3b and c). It can be seen that most of GNPs@Ti-5 and GNPs@V-10 adhere tightly to the copper matrix. There were few pores in the Cu/GNPs@TiC composite while there were many pores in the Cu/GNPs@VC composite. From Fig. 3d, it can be seen that there are much clearer gaps and pores between GNPs and copper matrix. The bonding was weak between raw GNPs and the copper matrix in the Cu/GNPs composite. In contrast, the fringes of GNPs@Ti-5 and GNPs@V-10 (Fig. 3e and f) were embedded in the copper matrix, i.e., the interfacial bonding was strengthened. The GNPs showed different fracture behaviors, as confirmed by the higher-magnification images. In Fig. 3d, the Cu/GNPs composites present a loose interface structure and the GNPs are intact without apparent tear. However, fragments and curled graphene sheets can be found on the surface of GNPs@Ti-5 and GNPs@V-10, indicating that GNPs@Ti-5 and GNPs@V-10 were torn during tension. It is worth noting that a smooth surface of GNPs without coating appeared in the fracture morphology. Comparing with Fig. 2h, it can be concluded that the tearing probably originated from the interlayer of graphene sheets. Besides the interaction of graphene sheets, there were Cu-TiC and TiC-GNPs interfaces in the Cu/GNPs@TiC composites. However, these results alone cannot show whether the tearing originated from the interlayer of graphene sheets, TiC-GNPs interface or Cu-TiC interface, requiring further investigations by TEM. Further interfacial structures of Cu/GNPs@TiC and Cu/GNPs@VC composites were investigated by TEM (Fig. 4a and b). In Fig. 4a, striplike graphene layers and clear interfaces without pores can be observed. Nano-sized TiC particles arranged closely form a continuous interlayer with an average thickness of 20 nm. Both sides of the GNPs were coated by TiC, showing that a TiC-GNPs-TiC sandwich structure was formed during the molten salt reaction processing. The TiC interlayer tightly adhered to both GNPs and Cu matrix. The TiC coating was in situ synthesized on the surface of GNPs. Thus, a strong bonding was formed at the GNPs-TiC interface. In addition, good wetting was achieved between the interface of TiC and Cu. One of the factors determining this wetting behavior could be the very low oxygen partial pressure at the interface in the SPS processing [29,30]. Another possible factor is titanium transfer from the carbide to the Cu matrix due to the slight TiC dissolution, which could decreased the contact angle of TiC-Cu

3.3. Mechanical, thermal and electrical properties Fig. 5a presents stress-strain curves of pure copper and composites. Pure copper and Cu/GNPs@VC composites displayed apparently yield in the stress-strain curves, showing the fracture mode of those was plastic fracture. However, for Cu/GNPs and Cu/GNPs@TiC composites, yield-free behavior was observed from the stress-strain curves. This observation indicates that the fracture mode changed from plastic fracture to brittle fracture. The tensile strength of three composites was higher than pure copper. For Cu/GNPs@TiC composites, the tensile strength was increased to 470 ± 7.2 MPa, which was 40% higher than that of Cu/GNPs. Moreover, based on the analysis of the variation tendency at the elastic deformation stage of the stress-strain curves, the elastic modulus of Cu/GNPs@VC composites was lower than other composites. Grain size refinement, dislocations strengthening and load transfer are usually involved strengthening mechanisms for graphene reinforced metal matrix composites. According to our previous research [13], the grain size of copper matrix was decreased by the addition of GNPs. Due to the thermal expansion mismatch of copper matrix and GNPs, a high dislocation density was generated at the interface of composites. The dislocation movement was impeded by the increased grain boundaries. According to the shear-lag model, load transfer from the matrix to reinforcement depends largely on the interface bonding [32]. For Cu/GNPs, the interfacial bonding strength was weak. Thus, the load could not transfer from the matrix to the reinforcement phase. In contrast, strong adhesion and a mechanical interlock effect were achieved at the GNPs-TiC and TiC-Cu interface in Fig. 4a, which ensure 315

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Fig. 4. Microstructure of the interface of Cu/GNPs @TiC composites (a) and Cu/GNPs @VC composites (b).

3.4. Electro-tribological properties

the load transfer. Thus, the high strength of GNPs was fully utilized and the Cu/GNPs@TiC composite exhibited higher tensile strength. Fig. 5b and c show the thermal diffusivity and the electrical conductivity of pure copper and different composites. Compared with pure copper, the thermal diffusivity and electrical conductivity of composites are lower. The thermal diffusivity of Cu/GNPs@TiC and Cu/GNPs@VC composites was slightly higher than that of Cu/GNPs composites. The electrical conductivity of three composites does not differ substantially. Therefore, these two interlayers had no apparent influence on the electrical conductivity of composites when the GNPs content was the same in the composites. The decrease of the thermal diffusivity is related to the interfacial structures. As is well known, the thermal transport mode of graphene is phonon conduction while the mode of pure copper is electron conduction. Therefore, electron-phonon coupling plays a determining role in the interfacial heat conduction. The Cu/GNPs had a loose interfacial structure, thus the electron-phonon coupling was poor and the interfacial thermal resistance was high. As a result, the thermal diffusivity of Cu/GNPs was decreased. For Cu/ GNPs@TiC and Cu/GNPs@VC, the interface was continuous and tightly bonded. The TiC and VC interlayers, with a mixture of metallic and covalent bonding, can reduce the energy of electron-phonon coupling and by this decrease the interfacial thermal resistance. It had been proven in the Cu/graphite system that a thin and integrated carbide interlayer could greatly decrease the interfacial thermal resistance from 8.60 × 10−6 m2 K W−1 to 0.11 × 10−6 m2 K W−1 [33]. For these reasons, the thermal diffusivity of Cu/GNPs@TiC and Cu/GNPs@VC was higher than that of Cu/GNPs. The thermal diffusivity of Cu/ GNPs@TiC and Cu/GNPs@VC was not improved greatly because the thermal conductivity of TiC and VC is relatively low [18]. When the graphene content of is low, the electrical conductivity of the copper matrix dominates. In this case, the change of interface structure has a limited effect on the electrical conductivity. Therefore, we expect that interlayers of early transition metal carbides with high thermal diffusivity could further improve the thermal performance of Cu/GNPs composites.

Fig. 6 shows the influence of carbide interlayer on the electrical sliding friction and wear properties. The contact voltage drop of composites with and without carbide interlayer versus sliding time is presented in Fig. 6a. The contact voltage drop gradually increased with sliding time. The trend of the contact voltage drop versus sliding time with carbide interlayer is similar to that without interlayer. The Cu/ GNPs@VC composites kept relatively stable and low contact voltage drop during the whole test, varying from 0.17V to 0.24 V. In the steady state, the contact voltage drop of Cu/GNPs@TiC composites was 1.16 V. From Fig. 6b, it could be concluded that the friction coefficient vs pure copper of Cu/GNPs@TiC and Cu/GNPs@VC composites in the steady state was higher than that in the initial state. Compared with Cu/ GNPs@VC composites, Cu/GNPs@TiC composites possessed lower friction coefficient. Fig. 6c shows the wear rate of Cu/GNPs@TiC and Cu/GNPs@VC composites with the sliding time. The wear rate of Cu/ GNPs@VC composites gradually decreased with prolonged sliding time. After 2 h, the wear rate of Cu/GNPs@VC composites tended to be stable and kept near 0.45 × 10−4 mm3/(N m). However, the wear rate of Cu/ GNPs@TiC composites decreased initially and then increased. At 5 h, the wear rate of Cu/GNPs@TiC composites reached 4.37 × 10−4 mm3/ (N m). For the Cu/GNPs composites, the friction and wear test was terminated at 1 h because of excessive wear. Graphene is a good solid lubricant and is used to reduce friction and wear of composites [34–37]. Due to the chemical inertness of carbon to copper, it was difficult for raw GNPs to form lubricating films on the exposed surfaces. However, the carbide interlayer strengthens the interfacial bonding between GNPs and Cu, which decreases the exfoliation of GNPs. Previous studies also reported that good interfacial bonding contributed to obtain low friction and wear [38,39]. Accordingly, the Cu/GNPs@VC and Cu/ GNPs@TiC composites showed lower wear rate than Cu/GNPs composites. Fig. 7 shows SEM images of the morphologies of the worn surfaces of the composites with and without carbide interlayer. There are

Fig. 5. (a) Stress-strain curves, (b) thermal diffusivity and (c) electrical conductivity of Cu/GNPs, Cu/GNPs@TiC and Cu/GNPs@VC composites.

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Fig. 6. (a) Contact voltage drop, (b) friction coefficient and (c) wear rate of Cu/GNPs, Cu/GNPs@TiC and Cu/GNPs@VC composites as functions of sliding time.

materials could be improved by the addition of V, which forms selflubricant oxides [40,41]. The Cu/GNPs@VC composites had low contact resistance and low wear rate.

continuous grooves on the surface of all composites, attributed to ploughing in abrasive wear. Owing to serious plastic deformation, Cu/ GNPs composites were stripped in the form of blocks. For Cu/ GNPs@TiC composites, numerous pits appeared because of particles cracking and separating from the matrix. No obvious deformation was observed from the fringes of the pits. Thus, large brittle fracture took place in the Cu/GNPs@TiC composites, which intensified the wear loss. However, no plastic deformation and spalling can be detected on the surface of Cu/GNPs@TiC composites. In Fig. 7c, the worn surface of Cu/GNPs@TiC composites is relatively smooth. From Fig. 7d–f, it can be concluded that GNPs became debris due to the poor adhesion of graphene to copper while GNPs@TiC and GNPs@VC were still welladhered to the copper matrix. The GNPs@VC film formed on the surface of GNPs@VC composites can act as lubricant to reduce wear [36]. Therefore, in terms of the worn morphologies, the Cu/GNPs@VC composites exhibit better wear resistance. For Cu/GNPs composites, the matrix materials were deformed, indicating low hardness. The roughness was high, which also aggravates the wear of Cu/GNPs composite. Furthermore, the effects of electrical and frictional heat can induce high temperature at the interface. Thus, massive stress occurs because of the mismatch of CTE between GNPs and Cu. Combining Figs. 5 and 7, it can deduced that the Cu/GNPs@TiC composite was brittle, so that the matrix easily cracked and flaked off. Hence, the wear rate of Cu/ GNPs@TiC composite was high. The wear resistance of coating

4. Conclusions Transition metal carbide (TiC and VC) coatings were synthesized on graphene nanoplatelets (GNPs) through molten-salt treatment to improve the interfacial properties of Cu/GNPs composites. The modified GNPs with a “carbide-GNPs-carbide” structure strengthened the interfacial bonding by replacing the “GNPs-gap-Cu” interface structure with a “GNPs-carbide-Cu” interface structure. In contrast to Cu/GNPs composites without interlayer, the tensile strength of Cu/GNPs composites with a TiC interlayer increased by 40% and the fracture elongation of Cu/GNPs composites with a VC interlayer increased by 133%. In addition, carbide interlayers were beneficial to the thermal diffusivity of Cu/GNPs composites, while it had little evident effect on the electrical conductivity. For the electro-tribological properties, the addition of a VC interlayer improved the wear resistance of Cu/GNPs composites. Supporting information TG/DTA spectra for GNPs@Ti-5 and TG/DSC spectra for GNPs@V10; the contents summary of carbide and residual GNPs for GNPs@Ti-5

Fig. 7. SEM morphologies of the wear tracks: (a), (d) Cu/GNPs; (b), (e) Cu/GNPs@TiC; (c), (f) Cu/GNPs@VC.

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and GNPs@V-10; SEM images of composites powders. [15]

Acknowledgment [16]

The authors sincerely acknowledge the National Natural Science Foundation of China (91426304) and the Ningbo Municipal Key Project (2014S10001) for the financial support, Preferred Foundation of Postdoctoral Scientific Research Project of Zhejiang Province (BSH1502159). P. E. also acknowledges support from the Swedish Foundation for Strategic Research (SSF) through the Future Research Leaders 5 program and the Synergy Grant FUNCASE.

[17]

[18]

[19]

Author contributions

[20]

X. S. and M. L. performed the experiment. F. C conceived the GNPs coated by the TiC or VC coating through the molten salts treatment. F. H, J. X and S. D participated discussing. Q. H supervised and conducted the whole experiment. P. E, X. S., and M. L. participated revising the manuscript.

[21] [22]

[23] [24]

Notes The authors declare no competing financial interests.

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