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carbon-Cu composites with Cr3C2 coatings by Cr-solution immersion method
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Enhanced wetting and properties of carbon/carbon-Cu composites with Cr3 C2 coatings by Cr-solution immersion method Bo Kong a , Jinming Ru b , Hongdi Zhang a , Tongxiang Fan a,∗ a b
State Key Laboratory for Metal Matrix Composite, Shanghai Jiaotong University, Shanghai 200204, China Institute of Advanced Manufacturing and Modern Equipment Technology, Jiangsu University, Zhenjiang 212013, China
a r t i c l e
i n f o
Article history: Received 1 August 2016 Received in revised form 8 December 2016 Accepted 20 December 2016 Available online xxx
a b s t r a c t A facile ammonium-dichromate solution immersion method was introduced to synthesize the copperwettable Cr3 C2 coating on and inside the carbon-carbon (C/C) preform. The formation mechanism and the microstructures of the Cr3 C2 coatings were studied. The contact angle between molten copper and the C/C decreased from 140◦ to 60◦ , demonstrating the significant improvement in the wettability. The Cr3 C2 coated C/C-Cu composite with only 4.2% porosity and 3.69 g cm−3 density was manufactured through copper infiltration. As a result, the thermal and electrical conductivity of the modified C/C-Cu increased significantly due to the infiltrated copper. Also the mechanical properties of the composites including both the flexural and compressive strengths were enhanced by over 100%. The modified C/C-Cu composite exhibited lower friction coefficients and wear rates for different load levels than those of the commercial C/Cu composite. These results demonstrate the potential of the modified C/C-Cu material for use in electrical contacts. © 2017 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.
1. Introduction Carbon-carbon (C/C) composites are attracting growing attention due to their high specific strength, good conductivity, low density and thermal expansion, as well as the excellent ablative and wear resistances [1–4]. Considering the porous structures of C/C composites despite some modified densification processes were introduced [5,6], a new type of carbon-copper (C-Cu) composite material can be manufactured by infiltrating the porous C/C materials with molten copper (C/C-Cu). Combining the excellent electrical conductivity of copper and the good wear resistance of C/C thanks to the lubrication of the deposited graphite, the C/C-Cu composite is an ideal candidate material for sliding electrical contacts, which require a low wear rate and friction coefficient as well as high electrical and thermal conductivity [7,8]. Moreover, the C/CCu composites can provide better mechanical performances than that of the C-Cu composites widely used as sliding electrical contact materials [9–11]. Due to the poor wetting between carbon and molten copper, the infiltration process requires a high pressure to overcome the resistance generated by small pores [12] inside the C/C compos-
∗ Corresponding author. E-mail address: [email protected] (T. Fan).
ites. Considering the high temperature required (>1356 K, which is the melting point of copper), the application of excessive pressure is unstable and unsafe. Moreover, the non-wetted condition causes weak interfacial bonding, which will undermine the conductivities and mechanical properties of the composite. Thus, improving the wetting between the C/C composites and copper is the main factor to enhance the infiltration and obtain composites with better properties. In previous studies, Ran et al. [13] reported that mixing titanium into the copper powder could effectively improve the wetting of the carbon by molten copper and promote the infiltration of the molten copper into the C/C. However, the addition of an alloying element such as Ti would sharply decrease the conductivity of copper. Song et al. [14] and Zhou et al. [15] used a coating method via molybdenum carbide to enhance the infiltration of copper into porous carbon materials. However, the procedure could possibly leave residues inside the porous carbon. The early research by Mortimer and Nicholas [16] proved that utilizing the copper-chromium alloys to form the chromium carbides could improve the wettability of carbon by molten copper. In recent studies, methods of forming the copper-wettable chromium carbides above carbon materials via molten salt [17] or slurry [18–20] were preferred for the surface modification. Chromium carbides in the phase of Cr3 C2 , Cr7 C3 and Cr23 C6 were synthesized and proved to improve the wetting between carbon materials and molten copper effectively. Never-
http://dx.doi.org/10.1016/j.jmst.2017.01.028 1005-0302/© 2017 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.
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Fig. 1. Schematic diagram of procedure to prepare coated-C/C composites.
theless, these procedures were relatively complicated and did not generate a thorough coating inside the porous C/C because of the relatively high viscosity of the molten salt or slurry. In this work, copper-wettable Cr3 C2 coatings were synthesized on the surface and on the inner pore-walls of C/C through a simple ammonium dichromate solution immersion method. The effect of the Cr3 C2 coatings on promoting the wetting behavior between molten copper and the C/C preform was observed using a hightemperature wettability experiment. The C/C-Cu composites were manufactured by infiltrating the Cr3 C2 -coated and uncoated C/C preforms with molten copper under a gas pressure. Microstructures, thermal and electrical conductivities of the C/C and the C/C-Cu composites, as well as their mechanical and frictional performance, were investigated to determine the improvement resulting from the enhanced copper infiltration. 2. Experimental 2.1. Materials and preparation Fig. 1 shows the procedure to prepare the coated C/C composites. The C/C composites used in this work were fabricated using chemical vapor infiltration (C.V.I) along with high temperature heat-treatment repeatedly to densify the original carbon felt from 0.3 g cm−3 to 1.46 g cm−3 . The final C/C composites have a total porosity of 30.2% and a closed porosity 2.4%. As also illustrated in Fig. 1, all the C/C samples were cleaned with ethanol in an ultrasonic bath to remove the impurities and attached powder. To achieve a complete and thorough immersion, the porous C/C were evacuated before the immersion in a 2 mol L−1 ammonium dichromate solution using deionized water as the solvent. The samples were then heated to 1473 K with a rate 300 K h−1 in an argon atmosphere to complete the coating treatment. Copper with 99.9% purity was used in this study. A pressureinfiltration apparatus similar to the hot-wall furnace described by Senkov et al. [21] was applied. As illustrated in Fig. 2, the C/C preform was placed under the copper bulk in an air-tight pressure vessel. The pressure vessel was placed in a 500 well-type furnace utilizing silicon carbide rods as the heating elements. The materials were heated with 300 K h−1 to 1473 K and maintained at the peak temperature for 1 h to make the copper melt completely. The infiltration was conducted under argon gas (99.9% purity) fed through a 10 mm internal-diameter high-pressure hose. The pressure applied for the infiltration was 1 MPa, and the pressure-holding time was 0.5 h. The C/Cu composite utilized in this study for comparison was supplied by Morgan Advanced Materials and was used as an electric brush in real applications. The C/Cu composite was fabricated via powder metallurgy, and its density was 5.8 g cm−3 .
Fig. 2. Schematic drawing of infiltration procedure to manufacture C/C-Cu composites.
2.2. Wettability measurement The contact angle (CA) was measured under a high-vacuum environment, similar to the conditions of the infiltration procedure. Substrate samples with dimension of 35 mm × 35 mm × 5 mm were cut from larger C/C cylinders. They were placed in a corundum tube (38 mm) horizontally underneath copper samples shaped in a semisphere (6 mm). The tube was evacuated to 10−4 Pa to eliminate the effect of ambient gas when measuring the contact angle. The chamber was heated with 300 K h−1 –1523 K in this high vacuum state. The continually changing morphologies of the copper melt were recorded by an infrared camera until the end of the heating process. The camera was controlled using Interface Data software. The contact angles () were measured utilizing the sessile drop method and were calculated as Young’s equation [22]: cos =
SV − SL , LV
(1)
where SV , SL , and LV are the surface free energies of the solid, solid/liquid interfaces and liquid, respectively. The contact angle at any instant was calculated by taking the mean values of left and right angles measured from the still images obtained from recorded videos. 2.3. Characterization Thermogravimetric–Differential thermogravimetric (TG–DTG, Netzsch STA449-F3-Jupiter) analysis was used to examine the change in mass of the immersed C/C sample from room temper-
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Fig. 3. (a) XRD patterns and (b) TG-DTG curves of the ammonium-dichromate-immersed C/C at different temperatures during modification procedure.
Fig. 4. (a) Surface and (b) cross-section microstructures of Cr3 C2 -coated C/C substrates and (c, d) linear EDS analysis on the cross-section of the coated C/C.
Table 1 Density, porosity, conductivities and mechanical properties of the different composites (C/C-Cu #1 is uncoated C/C infiltrated by copper, C/C-Cu #2 is coated C/C infiltrated by copper). Materials
Density (g cm−3 )
Porosity (%)
Thermal conductivity (W m−1 K−1 )
Electrical resistivity ( m)
Flexural strength (MPa)
Compressive strength (MPa)
C/C C/C-Cu #1 C/C-Cu #2
1.45 2.29 3.69
30.2 18.9 4.2
39.6 61.1 168.6
40.1010 11.6070 0.0755
53.4 69.2 129.5
74.1 106.2 240.4
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ature to 1673 K. X-ray diffraction (XRD, Cu-K˛ , 35 kV, 200 mA) was used for the phase analysis on the C/C samples. The porosity (p) of the materials was examined via the draining method using kerosene and was calculated as follows: p=
m2 −m1 L
VS
,
(2)
where m1 and m2 are the masses of the sample before and after immersion in the kerosene, respectively. The variable L is the density of the kerosene solution, and VS is the volume of the sample. The density of the samples was calculated using a densimeter (DH120 M) via the mass/volume method. The thermal diffusivities were measured on wafer samples (12.7 mm × 4 mm) at 298 K using the laser flash method (Netzsch LFA447), and the specific heat capacity was calculated by the ‘rule of mixtures’ based on the weight ratio [4,5]. The thermal conductivity was calculated as follows: = C · · ,
(3)
where C and are the specific heat and thermal diffusivity obtained from the measurement, respectively, and is the density of the material. The electric resistances were tested at room temperature on bulk samples (10 mm × 10 mm × 4 mm) by a DC resistance probe tester (ST2258C) according to ASTM B193-2000. The flexural strength was measured using a three-point bending test with a 30 mm span and a 0.5 mm min−1 crosshead speed according to JB/T 8133.7-1999. The compressive strength was tested on specimens measuring 10 mm × 10 mm at a loading speed of 1 mm min−1 . Both tests were performed on a universal testing machine (Zwick/Roell Z020). The dry friction tests of the composites were performed on bulk samples (30 mm × 7 mm × 6 mm) using the wear test machine (MM-T2000). The test time was 1 h, and the velocity of the wearing ring (40 mm, 45 steel with a hardness of 42-45 HRC) was 200 rpm. The samples were tested under loads of 50 N, 100 N and 200 N. The wear volume of the material was calculated from the mass loss of the material divided by the density of the material. The wear rate was evaluated from the wear volume divided by the test distance before dividing by the test load. Scanning electron microscopy (SEM, Hitachi S-4800 operated at 10 kV) and energy-dispersive spectroscopy (EDS, operated at 15 kV)
Fig. 5. Contact angle curves as a function of time and morphologies of copper melt on (a) Cr3 C2 -coated C/C and (b) uncoated C/C.
were used to characterize the microstructures and elemental distributions in the specimens. 3. Results and discussion 3.1. Synthesis and microstructure of chromium carbide layers on the C/C Fig. 3(a) shows the results of phase analysis on the ammoniumdichromate immersed C/C sample after different heating temperatures. According to the XRD patterns shown in Fig. 3(a), chromium oxides (Cr2 O3 ) formed on the sample after heating at 573 K, and chromium carbides were synthesized as Cr3 C2 on the C/C after increasing the heating temperature to 1473 K. The formed Cr3 C2 did not replace the C/C surface entirely. However, the presence of the Cr3 C2 coating affected the (002) peak of the graphite, indicating that the layer thicknesses were adequate. Compared to the preferred molten salt method [14,15], no potential residue possibly affect-
Fig. 6. SEM micrographs of interface between copper and modified C/C with (a) low, (b) medium, and (c) high magnification and (d) elemental mapping of carbon, chromium and copper.
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Fig. 7. Morphology (a), optical micrographs at low (b) and high (c) magnification of modified C/C-Cu and SEM micrographs (d, f) with linear EDS analysis (e, g) of modified C/C-Cu composite (d, e) and unmodified C/C-Cu (f, g).
Table 2 Friction coefficients and wear rates of different composite materials under different loads. Sample
C-C Modified C/C-Cu C/Cu
Wear rate (10−7 mm3 m−1 N−1 )
Friction coefficient 50 N
100 N
200 N
50 N
100 N
200 N
0.10 0.09 0.12
0.14 0.12 0.2
0.11 0.10 0.2
– 32.4 136.0
– 35.9 152.5
– 43.16 171.6
ing the infiltration procedure and properties of the composites remained on the C/C, which was confirmed by the XRD results. This represents one of the major advantages of this proposed method. Fig. 3(b) shows the TG-DSC curves of the immersed C/C sample. Water evaporation contributed to the slight drop in mass from room temperature to approximately 400 K. A significant decrease in mass occurred at 440.2 K, indicating the decomposition of ammonium dichromate and the formation of Cr2 O3 , which was detected from the XRD results shown in Fig. 3(a). The decomposition reaction of the ammonium dichromate can be expressed as follows:
(NH4 )2 Cr2 O7 →Cr2 O3 + 4H2 O + N2
(4)
Cr2 O3 remained until the temperature reached 1478.2 K when the other major mass loss event occurred, as shown by the curves in Fig. 3(b). Cr3 C2 was the product of the corresponding reaction. Utilizing the previously formed Cr2 O3 and the elemental carbon
from the surface and pore walls of the C/C, the Cr3 C2 coatings were synthesized according to the following reaction: 1368K
3Cr2 O3 + 13C → 2Cr3 C2 + 9CO G = 1076.81 − 0.7875T(G1473 = −83.1775J mol−1 ),
(5)
The Gibbs free energy (G) from Eq. (5) at 1473 K was calculated to verify the feasibility of the reaction in Eq. (5). A minus value of G confirmed that this reaction could occur during the heat treatment. Fig. 4(a) and (b) shows the microstructures of both the surface and the cross-section of the Cr3 C2 -coated C/C composite. The surface of C/C was coated with the chromium carbides, and good bulk mechanical properties of the C/C could be maintained with the carbon fibers intact. Cr3 C2 coatings were formed above the pyrolytic carbon, which was deposited around the carbon fibers, as demonstrated in the linear EDS results shown in Fig. 4(c) and (d). The synthesized chromium carbides formed stronger bonding with the C/C samples than those carbides formed using the molybdenum or tungsten powders via slurry method, which was introduced by
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Fig. 8. Comparison of modified C/C-Cu and C/C-Cu composites with previous studies [13,15,24]: (a) thermal conductivity; (b) specific thermal conductivity; (c) bending strength improvement associated with infiltration; (d) electrical resistivity.
Casalegno et al. [18,20]. The Coulombic forces between the positively charged chromium ions and the negatively charged carbon samples contributed to the bond strength. In addition, covalent bonds were generated through the reaction described in Eq. (5). The thorough immersion and infiltration inside the C/C due to the low viscosity of the solution allowed the formation of chromium carbides on the pyrolytic carbon pore walls, which favored the infiltration of the copper melt.
Additionally, the value of the contact angle decreased more sharply after the heating time exceeded 6000 s when the temperature was higher than 1473 K. As the temperature increased, the diffusion of the molten copper into the chromium carbides was enhanced due to the more active movement of atoms in the melt. The enhanced diffusion led to an increase in the work of adhesion, which is the typical parameter used to characterize the capillary action of a liquid infiltrant in a solid phase. The work of adhesion Wa was determined using Young’s equation:
3.2. Wetting behavior of the copper on the C/C substrate
Wa = LV (1 + cos)
Fig. 5 illustrates the contact angle as a function of time along with the changing morphologies of the molten copper on the Cr3 C2 coated and uncoated C/C substrates. For the semispherical copper samples used in the wettability experiment, it was easier to distinguish a good wetting situation from a poor one. The CA of the solid copper prior to melting on the coated substrates was 90◦ and decreased over time after the copper melted, as shown in Fig. 5(a). The CA was 61◦ at the end of the heating procedure, indicating good wetting between the copper and the substrate. The CA of the molten copper on the uncoated C/C increased as the heating procedure progressed, even though the CA and shape of the copper were almost the same before melting. The CA reached 145◦ and ended at 138◦ , indicating a non-wetted situation.
where LV is the surface free energy of the liquid (melt in this case) and is the contact angle between the molten copper and the solid substrate. It can be concluded from Eq. (6) that the increase in the work of adhesion would lead to the reduction in the CA indicating an improvement in the wettability. Fig. 6 shows micrographs of the interface between the solidified copper and the modified C/C substrate. As seen in Fig. 6(a) and (b), the copper follows the irregular C/C surface appropriately and partially infiltrates into the pores close to the interfacial boundary, which confirms the good wettability. The generation of the Cr3 C2 layers on the C/C substrate acted as lubricants for the molten copper and promoted the wetting behavior between the copper and the C/C. Fig. 6(c) shows that no interfacial gap appeared as in the
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(6)
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Fig. 9. Microstructures of worn surface of modified C/C-Cu composite under loads of (a) 50 N, (b) 100 N and (c) 200 N with (d) the EDS linear analysis, (e) low and (f) high magnification of worn surface of C/C composite under load of 200 N, and modified C/C-Cu composite under loads of (g) 50 N, (h) 100 N and (i) 200 N.
unwetted situation, and the Cr3 C2 coating turned into the interface between the copper and the C/C substrate as demonstrated in the elemental mapping results (Fig. 6(d)), which led to the tight bonding. 3.3. Microstructure and properties of the C/C-Cu composites Fig. 7(a) shows the sectional morphology of the modified C/CCu composite manufactured by infiltrating molten copper into the Cr3 C2 -coated C/C. The dimensions of the modified C/C-Cu sample were 40 mm × 58 mm. Fig. 7(b) and (c) shows the optical micrographs of the modified C/C-Cu composites. The improved wettability due to the Cr3 C2 coating resulted in a thorough infiltration of the C/C by molten copper. No obvious open-pores are observed in the composites. There was an adhesive connection between the infiltrated copper and the C/C perform of the modified C/C-Cu composite, unlike the non-wetting situation of the unmodified C/C-Cu shown in Fig. 7(f). The thickness of the chromium carbide interface is approximately 1 m, according to the linear EDS results shown in Fig. 7(d) and (e). Table 1 summarizes the properties of the C/C composite and the C/C-Cu composites. With enhanced infiltration due to the improved wettability, the porosity of the C/C-Cu composite decreased markedly from 18.9% to 4.2%, and the flexural strength of the composite increased from 53.4 MPa to 129.5 MPa, which is significantly higher than that of the commonly used C/Cu strip (42 MPa) [23]. Additionally, the compressive strength of the composites increased from 74.1 MPa to 240.4 MPa. Copper filled over 90% of the open pores (open porosity decreased from 28.1% to 2.1%)
and played a role in absorbing additional energy under load-bearing conditions by decreasing the composite porosity. The low remaining porosity of the C/C-Cu composite resulted from the thorough infiltration of the copper, which served to enhance the conductivities of the composites. Table 1 shows that the thermal conductivity of the composite increased from 39.6 W m−1 K−1 to 61.1 W m−1 K−1 and even further to 168.6 W m−1 K−1 , whereas the electrical resistivity decreased from 40.1010 m to 11.6070 m, and even further to 0.0755 m, which is less than an order of magnitude from that of pure copper (0.017 m). Moreover, the interconnected structure of the modified C/C-Cu composite leads to a unique electrical conduction mechanism, known as network conduction [23] in which the electrical conduction occurs through the copper matrix and is not affected by the presence of the carbon phase. Comparisons of the C/C-Cu composite from this work and the C/C-Cu composite from previous studies are illustrated in Fig. 8. The modified C/C-Cu in this work exhibited better thermal conductivity, specific thermal conductivity and bending strength improvement (Fig. 8(a)–(c)) and reached the same order of magnitude on electrical resistivity as that of C/C-Cu from Ran’s research [18] (Fig. 8(d)) and even that of pure copper (0.017 m), demonstrating the enhanced properties due to the thorough copper infiltration with improved wettability. 3.4. Friction performance of the C/C-Cu composites The friction coefficient and wear rates of the C/C, the modified C/C-Cu and the commercial C/Cu composites are listed in Table 2. Due to the ‘lubrication effect’ of the graphite phase under wearing
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conditions, the friction coefficient of the C/C is low (approximately 0.1), and it was reduced further with the addition of copper. Moreover, the modified C/C-Cu composite exhibited both a lower friction coefficient and wear rate than those of the commercial C/Cu under different loads, indicating that better friction properties are possessed by the modified C/C-Cu composite. The infiltrated copper was interconnected and bonded more strongly to the C/C preform because of the improved wettability. Therefore, the copper phase of the modified C/C-Cu composite could not easily spall or be removed during the wear process, avoiding forming the surface unevenness that occurs for the C/Cu composite manufactured via powder metallurgy method. This helped avoiding the increase in the frictional resistances for the wearing ring and obtaining lower friction coefficients as well as lower wear-rates. In addition, the fiber structure of the C/C preform was helpful for gathering the wear debris and forming compacted layers. These layers on the worn surface were effective for distributing contact stresses, inhibiting surface spalling and decreasing the wear loss [24]. The microstructures of the worn surfaces of the modified C/CCu specimens are shown in Fig. 9(a)–(d). For the load of 50 N, an integral and continuous layer of wear debris covered the surface of the sample, avoiding further adhesive wear. As the load increased to 100 N, the layer was shattered into large pieces and became rough, as illustrated in Fig. 9(b), leading to an increase in both the friction coefficient and wear rate. As the load increased to 200 N, the large pieces were shattered into smaller ones (Fig. 9(c)), resulting in a higher wear rate. Meanwhile, a relatively smoother layer was formed by the smaller particles or pieces, leading to a lower frictional resistance and friction coefficient. From the linear EDS analysis illustrated in Fig. 9(d), the previous layer consisted of carbon particles and copper oxide. The formation of copper oxide helped to protect the material surface from the adhesive wear due to metal-metal contact. Compared to the worn surface of the C/C under load of 200 N presented in Fig. 9(e) and (f), the copper oxide formed in the interlayer region between the deposited carbons acted as a smoother to reduce the extra frictional resistance caused by the unevenness between the carbon fiber and the deposited carbon layer. This procedure explains the decrease in the friction coefficient of the C/C with the infiltrated copper. The C/C-Cu composite without the modification exhibited the similar morphology change to the modified C/C-Cu under the load of 50 N and 100 N as illustrated in Fig. 9(g) and (h). As the load increased to 200 N (Fig. 9(i)), a smoother surface consists of large pieces of copper and debris from the C/C matrix was formed while small plots of copper were removed from the surface due to the lack of adhesion between the copper and the C/C preform. This led to a tribological characteristic more similar to the C/C composite for the unmodified C/C-Cu composite. It can be concluded that the C/C-Cu composite manufactured through enhanced copper infiltration due to improved wettability exhibited a lower friction coefficient and wear rate in addition to better conductivity and mechanical properties, demonstrating that it is a candidate for the use as sliding electrical contact materials. Furthermore, the facile method via solution immersion proposed in this study can be utilized for the surface modification of other carbon/copper systems.
(2) The wettability between the molten copper and the C/C was improved markedly as the CA between them decreased from 140◦ to 60◦ . As a result, the Cr3 C2 -coated C/C-Cu composite with density 3.69 g cm−3 and porosity 4.2% was manufactured through copper infiltration while the C/C-Cu with density only 2.29 g cm−3 and porosity 18.9% was obtained without the coatings. (3) With more copper infiltrated, the thermal conductivity and electrical resistance of the C/C-Cu was improved to 168.6 W m−1 K−1 and 0.0755 m, respectively. Also the flexural and compressive strengths of the C/C-Cu composites were increased to 129.5 MPa and 240.4 MPa, respectively. (4) Both the friction coefficients and the wear rates of the modified C/C-Cu composite under different loads were lower than the corresponding values of the commercial C/Cu composite, demonstrating that the modified C/C-Cu composite is a candidate for the use in electrical contacts. Acknowledgement The authors acknowledge the financial support from of the National Basic Research Program of China (Nos. 2012CB619600 and 2011CB012803). References [1] B. Venkataraman, G. Sundararajan, Acta Mater. 50 (2002) 1153–1163. [2] Y.Q. Qin, Z.S. Yu, Mater. Charact. 61 (2010) 635–639. [3] K.Z. Li, X.T. Shen, H.J. Li, S.Y. Zhang, T. Feng, L.L. Zhang, Carbon 49 (2011) 1208–1215. [4] Q.G. Fu, J.P. Zhang, H.J. Li, Carbon 93 (2015) 1081–1084. [5] P. Delhaèsa, M. Trinquecostea, J.F. Linesb, A. Cosculluelab, J.M. Goyhénècheb, M. Couzic, Carbon 43 (2005) 681–691. [6] K.Z. Li, J. Xie, Q.G. Fu, H.J. Li, L.J. Guo, Carbon 57 (2013) 161–168. [7] B.L. Lei, L.L. He, M.Z. Yi, L.P. Ran, H.J. Xu, Y.C. Ge, K. Peng, Carbon 49 (2011) 4554–4562. [8] C. Guiderdonia, E. Pavlenkob, V. Turqa, A. Weibela, P. Puechb, C. Estournèsc, A. Peigneya, W. Bacsab, C. Laurenta, Carbon 58 (2013) 185–197. ˇ Emmerb, J. Bielekc, L. Keleˇsic, Wear 265 (2008) 421–517. [9] J. Kováˇcik, S. [10] T. Futamia, M. Ohira, H. Muto, M. Sakai, Carbon 47 (2009) 2742–2751. [11] H. Sarmadi, A.H. Kokabi, S.M. Seyed-Reihani, Wear 304 (2013) 1–12. [12] A. Mortensen, T. Wong, Metall. Trans. A 21 (1990) 2257–2263. [13] L.P. Ran, K. Peng, M.Z. Yi, L. Yang, Mater. Lett. 65 (2011) 2076–2078. [14] J.L. Song, Q.G. Guo, X.Q. Gao, Z.H. Tao, J.L. Shi, L. Liu, Carbon 49 (2011) 3165–3170. [15] W.Y. Zhou, M.Z. Yi, K. Peng, L.P. Ran, Y.C. Ge, Mater. Lett. 145 (2015) 264–268. [16] D. Mortimer, M. Nicholas, J. Mater. Sci. 8 (1973) 640–648. [17] Q.P. Kang, X.B. He, S.B. Ren, L. Zhang, M. Wu, C.Y. Guo, W. Cui, X.H. Qu, Appl. Therm. Eng. 60 (2013) 423–429. [18] V. Casalegno, M. Salvo, M. Ferraris, Carbon 50 (2012) 2296–2306. [19] P. Appendino, M. Ferraris, V. Casalegno, J. Nucl. Mater. 329–333 (2004) 1563–1566. [20] P. Appendino, M. Ferraris, V. Casalegno, J. Nucl. Mater. 348 (2006) 102–107. [21] Metallic Materials with High Structural Efficiency, in: O. Senkov, D. Miracle, S. Firstov (Eds.), Kluwer Academic Publishers, New York, 2004, pp. 380–382. [22] N. Eustathopoulos, M. Nicholas, B. Drevet, Wettability at High Temperature, 1st ed., Pergamon, New York, 1999, p. 7. [23] Y.P. Tang, H.Z. Liu, H.J. Zhao, L. Liu, Y.T. Wu, Mater. Des. 29 (2008) 257–261. [24] L. Yang, L.P. Ran, M.Z. Yi, Mater. Des. 32 (2011) 2365–2369.
4. Conclusions (1) The copper-wettable Cr3 C2 coatings were synthesized on the surfaces and within the pore walls of carbon-carbon composites through the facile ammonium-dichromate solution immersion method.
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Enhanced wetting and properties of carbon/carbon-Cu composites with Cr3 C2 coatings by Cr-solution immersion method Bo Kong a , Jinming Ru b , Hongdi Zhang a , Tongxiang Fan a,∗ a b
State Key Laboratory for Metal Matrix Composite, Shanghai Jiaotong University, Shanghai 200204, China Institute of Advanced Manufacturing and Modern Equipment Technology, Jiangsu University, Zhenjiang 212013, China
a r t i c l e
i n f o
Article history: Received 1 August 2016 Received in revised form 8 December 2016 Accepted 20 December 2016 Available online xxx
a b s t r a c t A facile ammonium-dichromate solution immersion method was introduced to synthesize the copperwettable Cr3 C2 coating on and inside the carbon-carbon (C/C) preform. The formation mechanism and the microstructures of the Cr3 C2 coatings were studied. The contact angle between molten copper and the C/C decreased from 140◦ to 60◦ , demonstrating the significant improvement in the wettability. The Cr3 C2 coated C/C-Cu composite with only 4.2% porosity and 3.69 g cm−3 density was manufactured through copper infiltration. As a result, the thermal and electrical conductivity of the modified C/C-Cu increased significantly due to the infiltrated copper. Also the mechanical properties of the composites including both the flexural and compressive strengths were enhanced by over 100%. The modified C/C-Cu composite exhibited lower friction coefficients and wear rates for different load levels than those of the commercial C/Cu composite. These results demonstrate the potential of the modified C/C-Cu material for use in electrical contacts. © 2017 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.
1. Introduction Carbon-carbon (C/C) composites are attracting growing attention due to their high specific strength, good conductivity, low density and thermal expansion, as well as the excellent ablative and wear resistances [1–4]. Considering the porous structures of C/C composites despite some modified densification processes were introduced [5,6], a new type of carbon-copper (C-Cu) composite material can be manufactured by infiltrating the porous C/C materials with molten copper (C/C-Cu). Combining the excellent electrical conductivity of copper and the good wear resistance of C/C thanks to the lubrication of the deposited graphite, the C/C-Cu composite is an ideal candidate material for sliding electrical contacts, which require a low wear rate and friction coefficient as well as high electrical and thermal conductivity [7,8]. Moreover, the C/CCu composites can provide better mechanical performances than that of the C-Cu composites widely used as sliding electrical contact materials [9–11]. Due to the poor wetting between carbon and molten copper, the infiltration process requires a high pressure to overcome the resistance generated by small pores [12] inside the C/C compos-
∗ Corresponding author. E-mail address: [email protected] (T. Fan).
ites. Considering the high temperature required (>1356 K, which is the melting point of copper), the application of excessive pressure is unstable and unsafe. Moreover, the non-wetted condition causes weak interfacial bonding, which will undermine the conductivities and mechanical properties of the composite. Thus, improving the wetting between the C/C composites and copper is the main factor to enhance the infiltration and obtain composites with better properties. In previous studies, Ran et al. [13] reported that mixing titanium into the copper powder could effectively improve the wetting of the carbon by molten copper and promote the infiltration of the molten copper into the C/C. However, the addition of an alloying element such as Ti would sharply decrease the conductivity of copper. Song et al. [14] and Zhou et al. [15] used a coating method via molybdenum carbide to enhance the infiltration of copper into porous carbon materials. However, the procedure could possibly leave residues inside the porous carbon. The early research by Mortimer and Nicholas [16] proved that utilizing the copper-chromium alloys to form the chromium carbides could improve the wettability of carbon by molten copper. In recent studies, methods of forming the copper-wettable chromium carbides above carbon materials via molten salt [17] or slurry [18–20] were preferred for the surface modification. Chromium carbides in the phase of Cr3 C2 , Cr7 C3 and Cr23 C6 were synthesized and proved to improve the wetting between carbon materials and molten copper effectively. Never-
http://dx.doi.org/10.1016/j.jmst.2017.01.028 1005-0302/© 2017 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.
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Fig. 1. Schematic diagram of procedure to prepare coated-C/C composites.
theless, these procedures were relatively complicated and did not generate a thorough coating inside the porous C/C because of the relatively high viscosity of the molten salt or slurry. In this work, copper-wettable Cr3 C2 coatings were synthesized on the surface and on the inner pore-walls of C/C through a simple ammonium dichromate solution immersion method. The effect of the Cr3 C2 coatings on promoting the wetting behavior between molten copper and the C/C preform was observed using a hightemperature wettability experiment. The C/C-Cu composites were manufactured by infiltrating the Cr3 C2 -coated and uncoated C/C preforms with molten copper under a gas pressure. Microstructures, thermal and electrical conductivities of the C/C and the C/C-Cu composites, as well as their mechanical and frictional performance, were investigated to determine the improvement resulting from the enhanced copper infiltration. 2. Experimental 2.1. Materials and preparation Fig. 1 shows the procedure to prepare the coated C/C composites. The C/C composites used in this work were fabricated using chemical vapor infiltration (C.V.I) along with high temperature heat-treatment repeatedly to densify the original carbon felt from 0.3 g cm−3 to 1.46 g cm−3 . The final C/C composites have a total porosity of 30.2% and a closed porosity 2.4%. As also illustrated in Fig. 1, all the C/C samples were cleaned with ethanol in an ultrasonic bath to remove the impurities and attached powder. To achieve a complete and thorough immersion, the porous C/C were evacuated before the immersion in a 2 mol L−1 ammonium dichromate solution using deionized water as the solvent. The samples were then heated to 1473 K with a rate 300 K h−1 in an argon atmosphere to complete the coating treatment. Copper with 99.9% purity was used in this study. A pressureinfiltration apparatus similar to the hot-wall furnace described by Senkov et al. [21] was applied. As illustrated in Fig. 2, the C/C preform was placed under the copper bulk in an air-tight pressure vessel. The pressure vessel was placed in a 500 well-type furnace utilizing silicon carbide rods as the heating elements. The materials were heated with 300 K h−1 to 1473 K and maintained at the peak temperature for 1 h to make the copper melt completely. The infiltration was conducted under argon gas (99.9% purity) fed through a 10 mm internal-diameter high-pressure hose. The pressure applied for the infiltration was 1 MPa, and the pressure-holding time was 0.5 h. The C/Cu composite utilized in this study for comparison was supplied by Morgan Advanced Materials and was used as an electric brush in real applications. The C/Cu composite was fabricated via powder metallurgy, and its density was 5.8 g cm−3 .
Fig. 2. Schematic drawing of infiltration procedure to manufacture C/C-Cu composites.
2.2. Wettability measurement The contact angle (CA) was measured under a high-vacuum environment, similar to the conditions of the infiltration procedure. Substrate samples with dimension of 35 mm × 35 mm × 5 mm were cut from larger C/C cylinders. They were placed in a corundum tube (38 mm) horizontally underneath copper samples shaped in a semisphere (6 mm). The tube was evacuated to 10−4 Pa to eliminate the effect of ambient gas when measuring the contact angle. The chamber was heated with 300 K h−1 –1523 K in this high vacuum state. The continually changing morphologies of the copper melt were recorded by an infrared camera until the end of the heating process. The camera was controlled using Interface Data software. The contact angles () were measured utilizing the sessile drop method and were calculated as Young’s equation [22]: cos =
SV − SL , LV
(1)
where SV , SL , and LV are the surface free energies of the solid, solid/liquid interfaces and liquid, respectively. The contact angle at any instant was calculated by taking the mean values of left and right angles measured from the still images obtained from recorded videos. 2.3. Characterization Thermogravimetric–Differential thermogravimetric (TG–DTG, Netzsch STA449-F3-Jupiter) analysis was used to examine the change in mass of the immersed C/C sample from room temper-
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Fig. 3. (a) XRD patterns and (b) TG-DTG curves of the ammonium-dichromate-immersed C/C at different temperatures during modification procedure.
Fig. 4. (a) Surface and (b) cross-section microstructures of Cr3 C2 -coated C/C substrates and (c, d) linear EDS analysis on the cross-section of the coated C/C.
Table 1 Density, porosity, conductivities and mechanical properties of the different composites (C/C-Cu #1 is uncoated C/C infiltrated by copper, C/C-Cu #2 is coated C/C infiltrated by copper). Materials
Density (g cm−3 )
Porosity (%)
Thermal conductivity (W m−1 K−1 )
Electrical resistivity ( m)
Flexural strength (MPa)
Compressive strength (MPa)
C/C C/C-Cu #1 C/C-Cu #2
1.45 2.29 3.69
30.2 18.9 4.2
39.6 61.1 168.6
40.1010 11.6070 0.0755
53.4 69.2 129.5
74.1 106.2 240.4
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ature to 1673 K. X-ray diffraction (XRD, Cu-K˛ , 35 kV, 200 mA) was used for the phase analysis on the C/C samples. The porosity (p) of the materials was examined via the draining method using kerosene and was calculated as follows: p=
m2 −m1 L
VS
,
(2)
where m1 and m2 are the masses of the sample before and after immersion in the kerosene, respectively. The variable L is the density of the kerosene solution, and VS is the volume of the sample. The density of the samples was calculated using a densimeter (DH120 M) via the mass/volume method. The thermal diffusivities were measured on wafer samples (12.7 mm × 4 mm) at 298 K using the laser flash method (Netzsch LFA447), and the specific heat capacity was calculated by the ‘rule of mixtures’ based on the weight ratio [4,5]. The thermal conductivity was calculated as follows: = C · · ,
(3)
where C and are the specific heat and thermal diffusivity obtained from the measurement, respectively, and is the density of the material. The electric resistances were tested at room temperature on bulk samples (10 mm × 10 mm × 4 mm) by a DC resistance probe tester (ST2258C) according to ASTM B193-2000. The flexural strength was measured using a three-point bending test with a 30 mm span and a 0.5 mm min−1 crosshead speed according to JB/T 8133.7-1999. The compressive strength was tested on specimens measuring 10 mm × 10 mm at a loading speed of 1 mm min−1 . Both tests were performed on a universal testing machine (Zwick/Roell Z020). The dry friction tests of the composites were performed on bulk samples (30 mm × 7 mm × 6 mm) using the wear test machine (MM-T2000). The test time was 1 h, and the velocity of the wearing ring (40 mm, 45 steel with a hardness of 42-45 HRC) was 200 rpm. The samples were tested under loads of 50 N, 100 N and 200 N. The wear volume of the material was calculated from the mass loss of the material divided by the density of the material. The wear rate was evaluated from the wear volume divided by the test distance before dividing by the test load. Scanning electron microscopy (SEM, Hitachi S-4800 operated at 10 kV) and energy-dispersive spectroscopy (EDS, operated at 15 kV)
Fig. 5. Contact angle curves as a function of time and morphologies of copper melt on (a) Cr3 C2 -coated C/C and (b) uncoated C/C.
were used to characterize the microstructures and elemental distributions in the specimens. 3. Results and discussion 3.1. Synthesis and microstructure of chromium carbide layers on the C/C Fig. 3(a) shows the results of phase analysis on the ammoniumdichromate immersed C/C sample after different heating temperatures. According to the XRD patterns shown in Fig. 3(a), chromium oxides (Cr2 O3 ) formed on the sample after heating at 573 K, and chromium carbides were synthesized as Cr3 C2 on the C/C after increasing the heating temperature to 1473 K. The formed Cr3 C2 did not replace the C/C surface entirely. However, the presence of the Cr3 C2 coating affected the (002) peak of the graphite, indicating that the layer thicknesses were adequate. Compared to the preferred molten salt method [14,15], no potential residue possibly affect-
Fig. 6. SEM micrographs of interface between copper and modified C/C with (a) low, (b) medium, and (c) high magnification and (d) elemental mapping of carbon, chromium and copper.
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Fig. 7. Morphology (a), optical micrographs at low (b) and high (c) magnification of modified C/C-Cu and SEM micrographs (d, f) with linear EDS analysis (e, g) of modified C/C-Cu composite (d, e) and unmodified C/C-Cu (f, g).
Table 2 Friction coefficients and wear rates of different composite materials under different loads. Sample
C-C Modified C/C-Cu C/Cu
Wear rate (10−7 mm3 m−1 N−1 )
Friction coefficient 50 N
100 N
200 N
50 N
100 N
200 N
0.10 0.09 0.12
0.14 0.12 0.2
0.11 0.10 0.2
– 32.4 136.0
– 35.9 152.5
– 43.16 171.6
ing the infiltration procedure and properties of the composites remained on the C/C, which was confirmed by the XRD results. This represents one of the major advantages of this proposed method. Fig. 3(b) shows the TG-DSC curves of the immersed C/C sample. Water evaporation contributed to the slight drop in mass from room temperature to approximately 400 K. A significant decrease in mass occurred at 440.2 K, indicating the decomposition of ammonium dichromate and the formation of Cr2 O3 , which was detected from the XRD results shown in Fig. 3(a). The decomposition reaction of the ammonium dichromate can be expressed as follows:
(NH4 )2 Cr2 O7 →Cr2 O3 + 4H2 O + N2
(4)
Cr2 O3 remained until the temperature reached 1478.2 K when the other major mass loss event occurred, as shown by the curves in Fig. 3(b). Cr3 C2 was the product of the corresponding reaction. Utilizing the previously formed Cr2 O3 and the elemental carbon
from the surface and pore walls of the C/C, the Cr3 C2 coatings were synthesized according to the following reaction: 1368K
3Cr2 O3 + 13C → 2Cr3 C2 + 9CO G = 1076.81 − 0.7875T(G1473 = −83.1775J mol−1 ),
(5)
The Gibbs free energy (G) from Eq. (5) at 1473 K was calculated to verify the feasibility of the reaction in Eq. (5). A minus value of G confirmed that this reaction could occur during the heat treatment. Fig. 4(a) and (b) shows the microstructures of both the surface and the cross-section of the Cr3 C2 -coated C/C composite. The surface of C/C was coated with the chromium carbides, and good bulk mechanical properties of the C/C could be maintained with the carbon fibers intact. Cr3 C2 coatings were formed above the pyrolytic carbon, which was deposited around the carbon fibers, as demonstrated in the linear EDS results shown in Fig. 4(c) and (d). The synthesized chromium carbides formed stronger bonding with the C/C samples than those carbides formed using the molybdenum or tungsten powders via slurry method, which was introduced by
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Fig. 8. Comparison of modified C/C-Cu and C/C-Cu composites with previous studies [13,15,24]: (a) thermal conductivity; (b) specific thermal conductivity; (c) bending strength improvement associated with infiltration; (d) electrical resistivity.
Casalegno et al. [18,20]. The Coulombic forces between the positively charged chromium ions and the negatively charged carbon samples contributed to the bond strength. In addition, covalent bonds were generated through the reaction described in Eq. (5). The thorough immersion and infiltration inside the C/C due to the low viscosity of the solution allowed the formation of chromium carbides on the pyrolytic carbon pore walls, which favored the infiltration of the copper melt.
Additionally, the value of the contact angle decreased more sharply after the heating time exceeded 6000 s when the temperature was higher than 1473 K. As the temperature increased, the diffusion of the molten copper into the chromium carbides was enhanced due to the more active movement of atoms in the melt. The enhanced diffusion led to an increase in the work of adhesion, which is the typical parameter used to characterize the capillary action of a liquid infiltrant in a solid phase. The work of adhesion Wa was determined using Young’s equation:
3.2. Wetting behavior of the copper on the C/C substrate
Wa = LV (1 + cos)
Fig. 5 illustrates the contact angle as a function of time along with the changing morphologies of the molten copper on the Cr3 C2 coated and uncoated C/C substrates. For the semispherical copper samples used in the wettability experiment, it was easier to distinguish a good wetting situation from a poor one. The CA of the solid copper prior to melting on the coated substrates was 90◦ and decreased over time after the copper melted, as shown in Fig. 5(a). The CA was 61◦ at the end of the heating procedure, indicating good wetting between the copper and the substrate. The CA of the molten copper on the uncoated C/C increased as the heating procedure progressed, even though the CA and shape of the copper were almost the same before melting. The CA reached 145◦ and ended at 138◦ , indicating a non-wetted situation.
where LV is the surface free energy of the liquid (melt in this case) and is the contact angle between the molten copper and the solid substrate. It can be concluded from Eq. (6) that the increase in the work of adhesion would lead to the reduction in the CA indicating an improvement in the wettability. Fig. 6 shows micrographs of the interface between the solidified copper and the modified C/C substrate. As seen in Fig. 6(a) and (b), the copper follows the irregular C/C surface appropriately and partially infiltrates into the pores close to the interfacial boundary, which confirms the good wettability. The generation of the Cr3 C2 layers on the C/C substrate acted as lubricants for the molten copper and promoted the wetting behavior between the copper and the C/C. Fig. 6(c) shows that no interfacial gap appeared as in the
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(6)
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Fig. 9. Microstructures of worn surface of modified C/C-Cu composite under loads of (a) 50 N, (b) 100 N and (c) 200 N with (d) the EDS linear analysis, (e) low and (f) high magnification of worn surface of C/C composite under load of 200 N, and modified C/C-Cu composite under loads of (g) 50 N, (h) 100 N and (i) 200 N.
unwetted situation, and the Cr3 C2 coating turned into the interface between the copper and the C/C substrate as demonstrated in the elemental mapping results (Fig. 6(d)), which led to the tight bonding. 3.3. Microstructure and properties of the C/C-Cu composites Fig. 7(a) shows the sectional morphology of the modified C/CCu composite manufactured by infiltrating molten copper into the Cr3 C2 -coated C/C. The dimensions of the modified C/C-Cu sample were 40 mm × 58 mm. Fig. 7(b) and (c) shows the optical micrographs of the modified C/C-Cu composites. The improved wettability due to the Cr3 C2 coating resulted in a thorough infiltration of the C/C by molten copper. No obvious open-pores are observed in the composites. There was an adhesive connection between the infiltrated copper and the C/C perform of the modified C/C-Cu composite, unlike the non-wetting situation of the unmodified C/C-Cu shown in Fig. 7(f). The thickness of the chromium carbide interface is approximately 1 m, according to the linear EDS results shown in Fig. 7(d) and (e). Table 1 summarizes the properties of the C/C composite and the C/C-Cu composites. With enhanced infiltration due to the improved wettability, the porosity of the C/C-Cu composite decreased markedly from 18.9% to 4.2%, and the flexural strength of the composite increased from 53.4 MPa to 129.5 MPa, which is significantly higher than that of the commonly used C/Cu strip (42 MPa) [23]. Additionally, the compressive strength of the composites increased from 74.1 MPa to 240.4 MPa. Copper filled over 90% of the open pores (open porosity decreased from 28.1% to 2.1%)
and played a role in absorbing additional energy under load-bearing conditions by decreasing the composite porosity. The low remaining porosity of the C/C-Cu composite resulted from the thorough infiltration of the copper, which served to enhance the conductivities of the composites. Table 1 shows that the thermal conductivity of the composite increased from 39.6 W m−1 K−1 to 61.1 W m−1 K−1 and even further to 168.6 W m−1 K−1 , whereas the electrical resistivity decreased from 40.1010 m to 11.6070 m, and even further to 0.0755 m, which is less than an order of magnitude from that of pure copper (0.017 m). Moreover, the interconnected structure of the modified C/C-Cu composite leads to a unique electrical conduction mechanism, known as network conduction [23] in which the electrical conduction occurs through the copper matrix and is not affected by the presence of the carbon phase. Comparisons of the C/C-Cu composite from this work and the C/C-Cu composite from previous studies are illustrated in Fig. 8. The modified C/C-Cu in this work exhibited better thermal conductivity, specific thermal conductivity and bending strength improvement (Fig. 8(a)–(c)) and reached the same order of magnitude on electrical resistivity as that of C/C-Cu from Ran’s research [18] (Fig. 8(d)) and even that of pure copper (0.017 m), demonstrating the enhanced properties due to the thorough copper infiltration with improved wettability. 3.4. Friction performance of the C/C-Cu composites The friction coefficient and wear rates of the C/C, the modified C/C-Cu and the commercial C/Cu composites are listed in Table 2. Due to the ‘lubrication effect’ of the graphite phase under wearing
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conditions, the friction coefficient of the C/C is low (approximately 0.1), and it was reduced further with the addition of copper. Moreover, the modified C/C-Cu composite exhibited both a lower friction coefficient and wear rate than those of the commercial C/Cu under different loads, indicating that better friction properties are possessed by the modified C/C-Cu composite. The infiltrated copper was interconnected and bonded more strongly to the C/C preform because of the improved wettability. Therefore, the copper phase of the modified C/C-Cu composite could not easily spall or be removed during the wear process, avoiding forming the surface unevenness that occurs for the C/Cu composite manufactured via powder metallurgy method. This helped avoiding the increase in the frictional resistances for the wearing ring and obtaining lower friction coefficients as well as lower wear-rates. In addition, the fiber structure of the C/C preform was helpful for gathering the wear debris and forming compacted layers. These layers on the worn surface were effective for distributing contact stresses, inhibiting surface spalling and decreasing the wear loss [24]. The microstructures of the worn surfaces of the modified C/CCu specimens are shown in Fig. 9(a)–(d). For the load of 50 N, an integral and continuous layer of wear debris covered the surface of the sample, avoiding further adhesive wear. As the load increased to 100 N, the layer was shattered into large pieces and became rough, as illustrated in Fig. 9(b), leading to an increase in both the friction coefficient and wear rate. As the load increased to 200 N, the large pieces were shattered into smaller ones (Fig. 9(c)), resulting in a higher wear rate. Meanwhile, a relatively smoother layer was formed by the smaller particles or pieces, leading to a lower frictional resistance and friction coefficient. From the linear EDS analysis illustrated in Fig. 9(d), the previous layer consisted of carbon particles and copper oxide. The formation of copper oxide helped to protect the material surface from the adhesive wear due to metal-metal contact. Compared to the worn surface of the C/C under load of 200 N presented in Fig. 9(e) and (f), the copper oxide formed in the interlayer region between the deposited carbons acted as a smoother to reduce the extra frictional resistance caused by the unevenness between the carbon fiber and the deposited carbon layer. This procedure explains the decrease in the friction coefficient of the C/C with the infiltrated copper. The C/C-Cu composite without the modification exhibited the similar morphology change to the modified C/C-Cu under the load of 50 N and 100 N as illustrated in Fig. 9(g) and (h). As the load increased to 200 N (Fig. 9(i)), a smoother surface consists of large pieces of copper and debris from the C/C matrix was formed while small plots of copper were removed from the surface due to the lack of adhesion between the copper and the C/C preform. This led to a tribological characteristic more similar to the C/C composite for the unmodified C/C-Cu composite. It can be concluded that the C/C-Cu composite manufactured through enhanced copper infiltration due to improved wettability exhibited a lower friction coefficient and wear rate in addition to better conductivity and mechanical properties, demonstrating that it is a candidate for the use as sliding electrical contact materials. Furthermore, the facile method via solution immersion proposed in this study can be utilized for the surface modification of other carbon/copper systems.
(2) The wettability between the molten copper and the C/C was improved markedly as the CA between them decreased from 140◦ to 60◦ . As a result, the Cr3 C2 -coated C/C-Cu composite with density 3.69 g cm−3 and porosity 4.2% was manufactured through copper infiltration while the C/C-Cu with density only 2.29 g cm−3 and porosity 18.9% was obtained without the coatings. (3) With more copper infiltrated, the thermal conductivity and electrical resistance of the C/C-Cu was improved to 168.6 W m−1 K−1 and 0.0755 m, respectively. Also the flexural and compressive strengths of the C/C-Cu composites were increased to 129.5 MPa and 240.4 MPa, respectively. (4) Both the friction coefficients and the wear rates of the modified C/C-Cu composite under different loads were lower than the corresponding values of the commercial C/Cu composite, demonstrating that the modified C/C-Cu composite is a candidate for the use in electrical contacts. Acknowledgement The authors acknowledge the financial support from of the National Basic Research Program of China (Nos. 2012CB619600 and 2011CB012803). References [1] B. Venkataraman, G. Sundararajan, Acta Mater. 50 (2002) 1153–1163. [2] Y.Q. Qin, Z.S. Yu, Mater. Charact. 61 (2010) 635–639. [3] K.Z. Li, X.T. Shen, H.J. Li, S.Y. Zhang, T. Feng, L.L. Zhang, Carbon 49 (2011) 1208–1215. [4] Q.G. Fu, J.P. Zhang, H.J. Li, Carbon 93 (2015) 1081–1084. [5] P. Delhaèsa, M. Trinquecostea, J.F. Linesb, A. Cosculluelab, J.M. Goyhénècheb, M. Couzic, Carbon 43 (2005) 681–691. [6] K.Z. Li, J. Xie, Q.G. Fu, H.J. Li, L.J. Guo, Carbon 57 (2013) 161–168. [7] B.L. Lei, L.L. He, M.Z. Yi, L.P. Ran, H.J. Xu, Y.C. Ge, K. Peng, Carbon 49 (2011) 4554–4562. [8] C. Guiderdonia, E. Pavlenkob, V. Turqa, A. Weibela, P. Puechb, C. Estournèsc, A. Peigneya, W. Bacsab, C. Laurenta, Carbon 58 (2013) 185–197. ˇ Emmerb, J. Bielekc, L. Keleˇsic, Wear 265 (2008) 421–517. [9] J. Kováˇcik, S. [10] T. Futamia, M. Ohira, H. Muto, M. Sakai, Carbon 47 (2009) 2742–2751. [11] H. Sarmadi, A.H. Kokabi, S.M. Seyed-Reihani, Wear 304 (2013) 1–12. [12] A. Mortensen, T. Wong, Metall. Trans. A 21 (1990) 2257–2263. [13] L.P. Ran, K. Peng, M.Z. Yi, L. Yang, Mater. Lett. 65 (2011) 2076–2078. [14] J.L. Song, Q.G. Guo, X.Q. Gao, Z.H. Tao, J.L. Shi, L. Liu, Carbon 49 (2011) 3165–3170. [15] W.Y. Zhou, M.Z. Yi, K. Peng, L.P. Ran, Y.C. Ge, Mater. Lett. 145 (2015) 264–268. [16] D. Mortimer, M. Nicholas, J. Mater. Sci. 8 (1973) 640–648. [17] Q.P. Kang, X.B. He, S.B. Ren, L. Zhang, M. Wu, C.Y. Guo, W. Cui, X.H. Qu, Appl. Therm. Eng. 60 (2013) 423–429. [18] V. Casalegno, M. Salvo, M. Ferraris, Carbon 50 (2012) 2296–2306. [19] P. Appendino, M. Ferraris, V. Casalegno, J. Nucl. Mater. 329–333 (2004) 1563–1566. [20] P. Appendino, M. Ferraris, V. Casalegno, J. Nucl. Mater. 348 (2006) 102–107. [21] Metallic Materials with High Structural Efficiency, in: O. Senkov, D. Miracle, S. Firstov (Eds.), Kluwer Academic Publishers, New York, 2004, pp. 380–382. [22] N. Eustathopoulos, M. Nicholas, B. Drevet, Wettability at High Temperature, 1st ed., Pergamon, New York, 1999, p. 7. [23] Y.P. Tang, H.Z. Liu, H.J. Zhao, L. Liu, Y.T. Wu, Mater. Des. 29 (2008) 257–261. [24] L. Yang, L.P. Ran, M.Z. Yi, Mater. Des. 32 (2011) 2365–2369.
4. Conclusions (1) The copper-wettable Cr3 C2 coatings were synthesized on the surfaces and within the pore walls of carbon-carbon composites through the facile ammonium-dichromate solution immersion method.
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