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Cu composites
Materials Letters 57 (2003) 4583 – 4591 www.elsevier.com/locate/matlet
The effect of interfacial modifying on the mechanical and wear properties of SiCp/Cu composites Yongzhong Zhan *, Guoding Zhang State Key Laboratory of Metal Matrix Composites, Shanghai Jiaotong University, 1954 Huashan Road, Shanghai, People’s Republic of China Received 26 February 2003; accepted 24 April 2003
Abstract Copper matrix composites reinforced with 10 vol.% SiC particles with or without nickel coating were fabricated by powder metallurgy plus hot extrusion method. The results show that the densification, bulk hardness, ultimate flexural strength and ductility were improved by the electroless plating on the reinforcement, while the electrical conductivity remains at a fairly high level as the uncoated one. The nickel coating modifies the interface structural model and is effective in passing load between the matrix and the reinforcement. By lessening the extent of interfacial debonding, the coating changes the fracture and wear mechanism of the composites, therefore improving the mechanical properties and wear resistance. D 2003 Elsevier Science B.V. All rights reserved. Keywords: Interfacial modifying; SiCp/Cu composite; Mechanical property; Wear resistance
1. Introduction Due to the high electrical and thermal conductivity, good corrosion resistance and high melting point, copper is widely used in thermal and electronic applications, i.e. electronic packaging, electrical contacts and resistance welding electrodes. Nevertheless, the low mechanical property at both room and high temperatures limits the extensive application of pure copper. The room temperature mechanical strength can be improved dramatically by addition of small amount of elements such as Cr, Zr, Ag or Fe, etc. to form precipitation-hardened alloys. However, these
* Corresponding author. Tel.: +86-21-62933106; fax: +86-2162822012. E-mail address: [email protected] (Y. Zhan).
copper alloys lose their high strength at higher temperature (usually 500 jC) generally related to the structural instability caused by the coarsening of precipitation particles [1,2]. Furthermore, low wear resistance and conductivity are the other two limitations of these alloys. Discontinuously reinforced metal matrix composites are a class of materials that exhibit blending properties of the reinforcement and the matrix [3,4]. Their distinctive properties of high stiffness, high strength, good resistance and low coefficient of thermal expansion that could not be found in monolithic materials have promoted a number of applications for them. The incorporation of ceramic particulate reinforcement can improve the high-temperature mechanical property and wear resistance significantly, without severe deterioration of thermal and electrical conductivity of the matrix. In other words, particulate-rein-
0167-577X/03/$ - see front matter D 2003 Elsevier Science B.V. All rights reserved. doi:10.1016/S0167-577X(03)00365-3
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forced copper matrix composites may have many prominent advantages that the copper alloys do not possess. Therefore, these kinds of materials are considered to be promising candidates for applications where high conductivity, high mechanical property and good wear resistance are required [5]. The fabrication and performance of metal matrix composites are strongly influenced by the reinforcement –matrix interface. Proper bonding at the interface can attain good load transfer between phases. In some composites, the intrinsic lack of wetting between the matrix and the reinforcement causes difficulties in production and even debonding of interface during the service life. In this case, adhesion promoters are needed to modify the interface structural model. Studies of the Cu –SiC system indicate poor wettability between them [6]. A value of h = 140j was reported for the wetting of SiC by liquid Cu at 1100 jC [7]. Interfacial reactions were reported in many cases when the temperature was higher than 1173 K [8,9], although the formation of different copper silicides was reported by different authors. However, all in all, the interfacial bonding strength of the Cu – SiC system is considered to be weak for the practical application of composites. In this investigation, the effects of nickel coating on the mechanical and wear properties of silicon carbide particle-reinforced copper matrix composites were studied. A comparison of the fracture mechanism and wear mechanism of the composites reinforced with coated and uncoated SiC particles has been made based on a microstructural analysis.
2. Experimental procedure The metal matrix composites studied in this work were based on pure copper reinforced with 10 vol.% SiC particles. The starting material for the matrix was pure electrolytic copper powder having an average particle size of 48 Am. SiC particles with a diameter of 14 Am were irregular and angular in shape. Before composite fabrication, the nickel-coated SiC powder was prepared by an electroless plating process. At first, the SiC particles were surface cleaned by immersion in acetone and then in HNO3 (aq) for 20 min. The cleaned SiC was sensitized in an aqueous solution of SnCl2H2O and HCl for 20 min and then activated in a
solution containing 0.25 g/l of PdCl2 and 0.25 ml/l of 35% HCl for 25 min. After drying at 70 jC, the SiC powder was immersed in the electroless plating solution with 35 g/l of NiCl26H2O, 10 g/l of NaOOCCH2CH2COONa6H2O, 10 g/l of H2NCH2COOH, 20 g/l of NaH2PO2H2O and 2.5 mg/l of Pb(NO3)2. The pH was 4.5 and the temperature of the bath solution was 70 jC. The solution along with the SiC powder was stirred at 250 rpm and the plating time was 20 min. The weight of coating was estimated from the difference in weight of SiC particles before and after electroless coating, which was considered during composite processing. The average thickness of nickel coating was about 0.54 Am by converting the weight gain into particulate volumetric augmentation. Both the SiC particles with and without nickel coating were mixed with the calculated amount of copper powder, cold compacted and finally sintered in dissociated ammonia gas at 820 jC for 3 h. The sintered compacts were further hot extruded at a ratio of 10. The bulk densities of the hot-extruded composites were measured by a method based on Archimedes’ law and compared with theoretical densities to obtain various degree of densification. The porosity can be determined by the equation fp ¼ 1 q=q0 where fp is the pore volume fraction, q is the measured density and q0 is the theoretical density. Bulk hardness was measured on a Brinell hardness tester. For measurement of the volume electrical conductivity, the four-probe method was used. Bending tests were performed on a Shimadsu AG-100 kNA testing machine. The load –extension curve obtained from the bending test was converted to a true stress(r)– true stain(e) curve. Dry sliding wear tests were carried out using a block-on-ring type wear machine within a load range of 20– 100 N and at a constant sliding velocity of 0.42 m/s. The samples were machined into rectangular blocks of 6 7 16 mm. The faces of 7 16 mm were put in contact with the slider rings. The counterface ring (outer diameter: 40 mm) was made of GCr15-type bearing steel and had a bulk hardness of HRC62 F 2. Prior to wear testing, both contact surfaces were polished, cleaned in acetone in an ultrasonic
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cleaner and dried. The sliding distance for each test was normally 1.5 103 m for all normal loads. The wear losses of the specimens were measured by the width of the wear track, and were then converted to volume losses using the method of GB12444.2-90 [10]. Results from the specimens with wear track geometry that were not compatible with the specification of GB12444.2-90 were discarded. The fracture surfaces and worn surfaces were observed by scanning electron microscope (SEM). Specimens for transmission electron microscope (TEM) were cut from the bending samples and then prepared by an ion thinning method. TEM observation was performed with a JEOL JEM-100CX instrument.
3. Results and discussion 3.1. Microstructure and basic properties Fig. 1a shows the surface morphology of electroless nickel plated SiC particles. It was found that a uniform layer of nickel film formed on the SiC particles. The microstructure of the SiCp(Ni)/Cu composite parallel to the extrusion direction is illustrated in Fig. 1b. It can be observed that SiC particles were dispersed uniformly in the copper matrix. Micropore and particle fracture can hardly be found after hot extrusion. Table 1 lists the basic properties of both composites. The porosity of SiCp(Ni)/Cu composite is
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Table 1 Basic properties of nickel coated and uncoated SiC particle reinforced composites Material
Porosity (%)
Hardness (HB)
Electrical conductivity (% IACS)
SiCp/Cu SiCp(Ni)/Cu
1.4 0.9
78.1 84.6
82.63 80.17
lower than that of the SiCp/Cu composite. Therefore, it can be deduced that the bonding strength between copper matrix and SiC particle was increased by interfacial modifying. Nickel-coated SiC particles reinforced copper matrix composite exhibits higher hardness than the uncoated SiC reinforced one, while the electrical conductivity is not obviously deteriorated (Table 1). Hardness is a parameter that reflects the local deformation resistance of a material. As the nickel coating reduces the micropores near the interface and increases the bonding strength, the composite can resist local deformation more effectively. The effect of coating on the mechanical properties will be discussed in the following part. Though the coating introduces a new element in the matrix, which can increase the electrical resistivity of pure copper, increment of densification improves the electrical conductivity on the other hand. Consequently, the SiCp(Ni)/Cu composite maintains a fairly high electrical conductivity as the uncoated one.
Fig. 1. Micrographs of (a) SiC particles after electroless nickel coating and (b) microstructure of copper matrix composite reinforced with 10 vol.% SiC particles with nickel coating.
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Fig. 2. Stress – strain curves of SiCp/Cu and SiCp(Ni)/Cu composites.
3.2. Mechanical property Typical bending stress –strain curves of the nickel coated and uncoated SiC particle composites are shown in Fig. 2. It is clear that the SiCp(Ni)/Cu composite exhibits a better combination of flexural strength and ductility than the SiC/Cu composite. The ultimate flexural strength and corresponding strain of the coated SiC composite are 445.9 MPa and 27.3%, while in the case of the uncoated SiC reinforced composite are 381.5 MPa and 24.6%, respectively. The fracture surface of the two composites after the bending test is shown in Fig. 3. It can be found that the fracture mechanism of the SiCp/Cu composite
includes the ductile rupture at copper matrix and the debonding of the SiC – Cu interface (as shown in Fig. 3a). Pulling out of SiC particles and big dimples with and without particles in them can be clearly seen on the fracture surface. It may be concluded that the poor bonding strength of the SiC – Cu interface is the primary factor responsible for the relatively low flexural strength and ductility. The fracture appearance of the SiCp(Ni)/Cu composite is much different from the former one. It is not as even as that of the SiCp/Cu composite in both low and high magnification views. Debonding of the SiC –Cu interface is lessened, and particle pulling out from the copper matrix is not obvious on the fracture surface. Severe ductile deformation in the matrix is the main fracture mechanism of this composite. Fig. 4 presents the typical interface of the two composites. Impurities can be clearly seen on the SiC – Cu interface in the case of uncoated SiC-reinforced composite (Fig. 4a). SEM and EDAX analysis of the original surface of the SiC particle show that a small quantity of dissociative Fe and O can be found. These impurities may come from the environment during the producing process of SiC powder. Furthermore, at some places of the interface, interspace between SiC and the matrix can be clearly seen (Fig. 4b). The interface of SiCp(Ni)/Cu composite, however, is smooth and compact. Copper matrix integrates with the reinforcement closely, and impurities can hardly be found (Fig. 4c). During the electroless plating process of the SiC powder, impurities on
Fig. 3. SEM photographs showing the fracture surfaces the composites. (a) SiCp/Cu composite and (b) SiCp(Ni)/Cu composites.
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gram, it is indicated that an infinite solid solution with ultimate tensile strength which is much higher than that of the pure copper is formed by the two metals [11]. Since the preparation temperature of the composites (820 jC) is high enough for the Cu and Ni atoms to diffuse to form a film of Cu – Ni alloy at the interface, the component of this film changes continuously from the copper side to the nickel side. The bonding strength of this SiC –Ni – Cu interface is high enough to pass load from the matrix to the reinforcement effectively. As a result, in the bending test, debonding of the interface is more difficult to take place than in the uncoated SiC-reinforced composite. In other words, the nickel coating on the SiC reinforcement improves the strength and ductility of the composite by lessening the extent of interfacial debonding. 3.3. Wear property The wear rate of the nickel coated and uncoated SiC particle composites is plotted as a function of the normal load in Fig. 5. An increase in wear rate with increasing normal load is observed for both the materials. However, the nickel-coated SiC-reinforced composite exhibits lower wear rate than the uncoated one at all the load levels. With the increment of normal load, the wear rate of the SiCp(Ni)/Cu composite increases slowly and keeps a rather low value. Therefore, nickel coating on SiC particle surface
Fig. 4. TEM micrographs showing the characteristics of the interfaces of the composites. (a) Impurities and (b) interspace on the SiC – Cu interface; (c) SiC – Ni – Cu interface.
the SiC particle surface were eliminated by the activation and sensitization processes to get a high bonding strength between the nickel coating and the reinforcement. From the Cu –Ni binary phase dia-
Fig. 5. The variation of wear rates with the applied load for SiCp/Cu and SiCp(Ni)/Cu composites.
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Fig. 6. SEM photographs of worn surfaces of the composites (100 N). (a) Low magnification view of the SiCp/Cu composite; (b) high magnification of the SiCp/Cu composite; (c) low magnification view of the SiCp(Ni)/Cu composite.
Fig. 7. Morphology of wear debris generated from the composites tested at 100 N: (a) SiCp/Cu composite; (b) SiCp(Ni)/Cu composite.
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improves the wear resistance of the composite, especially at high load levels. Fig. 6 presents the surface morphologies of the samples worn at 100 N. It can be clearly seen that the worn surface of the uncoated SiC-reinforced composite is largely composed of deep grooves parallel to the sliding direction and low-lying patches (Fig. 6a). These patches are essentially composed of shear dimples, which indicate that they are the places where the wear debris detaches from the worn surface. High magnification view of a groove (Fig. 6b) shows the presence of black pits and microcrack perpendicular to the sliding direction. A pit should be the original place for a SiC particle and is formed when a particle is pulled out from the matrix during the wear process.
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The worn surface of the coated SiC-reinforced composite is relatively smooth, and some discontinuousand-slim grooves can be found, as is shown in Fig. 6c. Many little shear dimples on it suggest severe deformation in the subsurface region. In Fig. 7, the morphologies of the wear debris generated from the two composites tested at 100 N are compared. In the case of the SiCp/Cu composite (Fig. 7a), the size of some of the debris is about 40 –50 Am, approximates that of the low-lying patches on the worn surface. The debris is large and flake-like, indicating that large pieces of material had delaminated from the contact surface under this load level. The wear debris of the SiCp(Ni)/Cu composite is small (15 –20 Am) and shows an equiaxed morphology(Fig. 7b).
Fig. 8. Longitudinal cross-section perpendicular to the worn surface showing subsurface region of the composites and EDAX result of the MML layer (tested at 100 N). (a) Low magnification view showing the cracks beneath the worn surface of SiCp/Cu composite; (b) subsurface region of the SiCp(Ni)/Cu composite; (c) EDAX spectrum of the MML layer in SiCp/Cu composite.
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Longitudinal cross-sections perpendicular to worn surface showing the subsurface region are presented in Fig. 8a and b to analyse the wear mechanism of the composites. For both the composites, it is clearly seen that a tribolayer which is not uniform in thickness covers the worn surface as a ‘‘protective layer’’. This kind of layer has been observed on the worn surface of aluminum matrix composites and is called the mechanically mixed layer (MML)[12]. EDAX analysis of the MML layer on the SiC/Cu composite indicates that it contains such elements as Fe, Cu, C, Si and O (Fig. 8c). Fe element is rich in the layer, and it must come from the surface of the counterface when the test sample slid against the steel ring. The presence of O implies that Fe and Cu must have been oxidized, for the flash temperature on the contact surface is often high enough to reach the oxidation temperature of the fine wear debris [13]. Subsurface cracks can be clearly seen just beneath the MML layer in the SiCp/Cu composites (Fig. 8a). These cracks were formed around the SiC particles and are apt to extend to the worn surface along the sliding direction. In the wear test, the worn surface of the sample was rubbed by counterface repeatedly. Thereupon, large strains and strain gradients were generated in the subsurface region. As mentioned above, the SiC– Cu interface of this composite is so weak that interfacial debonding takes place when the stress attains the critical bonding strength. In other words, the interface is a preferential place for the crack nucleation and propagation. Crack propagation around a SiC particle which is proud on the worn surface can lead to the direct scaling off (Fig. 6b). However, if the crack is formed in the subsurface region which is far away from the contact surface, it will propagate to the nearby matrix and connect with the other cracks (Fig. 8a). When the two sides of the crack finally extend to reach the contact surface, a piece of composite debris delaminates from the bulk material. It should be noted that the delamination of the MML layer is also an important mechanism for this material. In the case of the SiCp(Ni)/Cu composite, the nickel coating improves the bonding strength of the interface. Interfacial debonding and microcrack propagation are not so easy to take place under the same stress condition (Fig. 8b). Therefore, the SiC particles act as load-bearing components and lessen the strain and
strain gradient in the subsurface region, which mitigates the wear loss of the composite effectively. The main wear mechanism, however, is the microcutting effect of the counterface asperity on the worn surface, and adhesion of the composite to steel induced by the ductile deformation in the subsurface region. In a word, interfacial modifying is effective to improve the wear resistance of copper matrix composites by preventing microcrack nucleation in the interface. 4. Conclusions The major conclusions resulting from the work presented in this paper can be listed as follows. 1. A uniform nickel surface can be formed on the SiC powders with electroless plating process. Compact composites with coated or uncoated SiC particles dispersed uniformly in the matrix can be fabricated successfully by the powder metallurgy plus hot extrusion method. 2. Relative density and bulk hardness of nickel-coated SiC-reinforced composite are higher than that of the uncoated one, while electrical conductivity is not significantly deteriorated by the introduction of coating. 3. The SiCp(Ni)/Cu composite exhibits better combination of flexural strength and ductility than the SiCp/Cu composite. The nickel coating on the SiC reinforcement improves the mechanical properties by lessening the extent of interfacial debonding. 4. The nickel coating modifies the interface structural model and changes the wear mechanism. It is in favor of the load-bearing effect of the reinforcement, thus improves the wear resistance of the composite. References [1] J.B. Correia, H.A. Davies, C.M. Sellars, Acta Mater. 45 (1997) 177. [2] M.A. Morris, D.G. Morris, Acta Metall. 35 (1987) 2511. [3] W.C. Harrigan Jr., Mater. Sci. Eng. A244 (1998) 75. [4] Y.B. Liu, S.C. Lim, L. Lu, M.O. Lai, J. Mater. Sci. 29 (1994) 1999. [5] R. Murakami, K. Matsui, Wear 201 (1996) 193. [6] P. Nikolopoulos, S. Agathopoulos, G.N. Angelopoulos, J. Mater. Sci. 27 (1992) 139.
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R. Warren, C.-H. Andersson, Composites 15 (1984) 101. H.-K. Lee, J.-K. Lee, J. Mater. Sci. Lett. 11 (1992) 550. C.-D. Qin, B. Derby, Br. Ceram., Trans. J. 90 (1991) 124. China National Standards, GB12444.2-90, Standards Press of China, Beijing, 1994.
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[11] S.S. Tan, Nonferrous Materials, Metallurgy Industry Press, Beijing, 1993, p. 51. [12] B. Venkataraman, G. Sundararajan, Acta Mater. 44 (1996) 461. [13] S.C. Lim, M.F. Ashby, Acta Metall. 35 (1987) 8.
The effect of interfacial modifying on the mechanical and wear properties of SiCp/Cu composites Yongzhong Zhan *, Guoding Zhang State Key Laboratory of Metal Matrix Composites, Shanghai Jiaotong University, 1954 Huashan Road, Shanghai, People’s Republic of China Received 26 February 2003; accepted 24 April 2003
Abstract Copper matrix composites reinforced with 10 vol.% SiC particles with or without nickel coating were fabricated by powder metallurgy plus hot extrusion method. The results show that the densification, bulk hardness, ultimate flexural strength and ductility were improved by the electroless plating on the reinforcement, while the electrical conductivity remains at a fairly high level as the uncoated one. The nickel coating modifies the interface structural model and is effective in passing load between the matrix and the reinforcement. By lessening the extent of interfacial debonding, the coating changes the fracture and wear mechanism of the composites, therefore improving the mechanical properties and wear resistance. D 2003 Elsevier Science B.V. All rights reserved. Keywords: Interfacial modifying; SiCp/Cu composite; Mechanical property; Wear resistance
1. Introduction Due to the high electrical and thermal conductivity, good corrosion resistance and high melting point, copper is widely used in thermal and electronic applications, i.e. electronic packaging, electrical contacts and resistance welding electrodes. Nevertheless, the low mechanical property at both room and high temperatures limits the extensive application of pure copper. The room temperature mechanical strength can be improved dramatically by addition of small amount of elements such as Cr, Zr, Ag or Fe, etc. to form precipitation-hardened alloys. However, these
* Corresponding author. Tel.: +86-21-62933106; fax: +86-2162822012. E-mail address: [email protected] (Y. Zhan).
copper alloys lose their high strength at higher temperature (usually 500 jC) generally related to the structural instability caused by the coarsening of precipitation particles [1,2]. Furthermore, low wear resistance and conductivity are the other two limitations of these alloys. Discontinuously reinforced metal matrix composites are a class of materials that exhibit blending properties of the reinforcement and the matrix [3,4]. Their distinctive properties of high stiffness, high strength, good resistance and low coefficient of thermal expansion that could not be found in monolithic materials have promoted a number of applications for them. The incorporation of ceramic particulate reinforcement can improve the high-temperature mechanical property and wear resistance significantly, without severe deterioration of thermal and electrical conductivity of the matrix. In other words, particulate-rein-
0167-577X/03/$ - see front matter D 2003 Elsevier Science B.V. All rights reserved. doi:10.1016/S0167-577X(03)00365-3
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forced copper matrix composites may have many prominent advantages that the copper alloys do not possess. Therefore, these kinds of materials are considered to be promising candidates for applications where high conductivity, high mechanical property and good wear resistance are required [5]. The fabrication and performance of metal matrix composites are strongly influenced by the reinforcement –matrix interface. Proper bonding at the interface can attain good load transfer between phases. In some composites, the intrinsic lack of wetting between the matrix and the reinforcement causes difficulties in production and even debonding of interface during the service life. In this case, adhesion promoters are needed to modify the interface structural model. Studies of the Cu –SiC system indicate poor wettability between them [6]. A value of h = 140j was reported for the wetting of SiC by liquid Cu at 1100 jC [7]. Interfacial reactions were reported in many cases when the temperature was higher than 1173 K [8,9], although the formation of different copper silicides was reported by different authors. However, all in all, the interfacial bonding strength of the Cu – SiC system is considered to be weak for the practical application of composites. In this investigation, the effects of nickel coating on the mechanical and wear properties of silicon carbide particle-reinforced copper matrix composites were studied. A comparison of the fracture mechanism and wear mechanism of the composites reinforced with coated and uncoated SiC particles has been made based on a microstructural analysis.
2. Experimental procedure The metal matrix composites studied in this work were based on pure copper reinforced with 10 vol.% SiC particles. The starting material for the matrix was pure electrolytic copper powder having an average particle size of 48 Am. SiC particles with a diameter of 14 Am were irregular and angular in shape. Before composite fabrication, the nickel-coated SiC powder was prepared by an electroless plating process. At first, the SiC particles were surface cleaned by immersion in acetone and then in HNO3 (aq) for 20 min. The cleaned SiC was sensitized in an aqueous solution of SnCl2H2O and HCl for 20 min and then activated in a
solution containing 0.25 g/l of PdCl2 and 0.25 ml/l of 35% HCl for 25 min. After drying at 70 jC, the SiC powder was immersed in the electroless plating solution with 35 g/l of NiCl26H2O, 10 g/l of NaOOCCH2CH2COONa6H2O, 10 g/l of H2NCH2COOH, 20 g/l of NaH2PO2H2O and 2.5 mg/l of Pb(NO3)2. The pH was 4.5 and the temperature of the bath solution was 70 jC. The solution along with the SiC powder was stirred at 250 rpm and the plating time was 20 min. The weight of coating was estimated from the difference in weight of SiC particles before and after electroless coating, which was considered during composite processing. The average thickness of nickel coating was about 0.54 Am by converting the weight gain into particulate volumetric augmentation. Both the SiC particles with and without nickel coating were mixed with the calculated amount of copper powder, cold compacted and finally sintered in dissociated ammonia gas at 820 jC for 3 h. The sintered compacts were further hot extruded at a ratio of 10. The bulk densities of the hot-extruded composites were measured by a method based on Archimedes’ law and compared with theoretical densities to obtain various degree of densification. The porosity can be determined by the equation fp ¼ 1 q=q0 where fp is the pore volume fraction, q is the measured density and q0 is the theoretical density. Bulk hardness was measured on a Brinell hardness tester. For measurement of the volume electrical conductivity, the four-probe method was used. Bending tests were performed on a Shimadsu AG-100 kNA testing machine. The load –extension curve obtained from the bending test was converted to a true stress(r)– true stain(e) curve. Dry sliding wear tests were carried out using a block-on-ring type wear machine within a load range of 20– 100 N and at a constant sliding velocity of 0.42 m/s. The samples were machined into rectangular blocks of 6 7 16 mm. The faces of 7 16 mm were put in contact with the slider rings. The counterface ring (outer diameter: 40 mm) was made of GCr15-type bearing steel and had a bulk hardness of HRC62 F 2. Prior to wear testing, both contact surfaces were polished, cleaned in acetone in an ultrasonic
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cleaner and dried. The sliding distance for each test was normally 1.5 103 m for all normal loads. The wear losses of the specimens were measured by the width of the wear track, and were then converted to volume losses using the method of GB12444.2-90 [10]. Results from the specimens with wear track geometry that were not compatible with the specification of GB12444.2-90 were discarded. The fracture surfaces and worn surfaces were observed by scanning electron microscope (SEM). Specimens for transmission electron microscope (TEM) were cut from the bending samples and then prepared by an ion thinning method. TEM observation was performed with a JEOL JEM-100CX instrument.
3. Results and discussion 3.1. Microstructure and basic properties Fig. 1a shows the surface morphology of electroless nickel plated SiC particles. It was found that a uniform layer of nickel film formed on the SiC particles. The microstructure of the SiCp(Ni)/Cu composite parallel to the extrusion direction is illustrated in Fig. 1b. It can be observed that SiC particles were dispersed uniformly in the copper matrix. Micropore and particle fracture can hardly be found after hot extrusion. Table 1 lists the basic properties of both composites. The porosity of SiCp(Ni)/Cu composite is
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Table 1 Basic properties of nickel coated and uncoated SiC particle reinforced composites Material
Porosity (%)
Hardness (HB)
Electrical conductivity (% IACS)
SiCp/Cu SiCp(Ni)/Cu
1.4 0.9
78.1 84.6
82.63 80.17
lower than that of the SiCp/Cu composite. Therefore, it can be deduced that the bonding strength between copper matrix and SiC particle was increased by interfacial modifying. Nickel-coated SiC particles reinforced copper matrix composite exhibits higher hardness than the uncoated SiC reinforced one, while the electrical conductivity is not obviously deteriorated (Table 1). Hardness is a parameter that reflects the local deformation resistance of a material. As the nickel coating reduces the micropores near the interface and increases the bonding strength, the composite can resist local deformation more effectively. The effect of coating on the mechanical properties will be discussed in the following part. Though the coating introduces a new element in the matrix, which can increase the electrical resistivity of pure copper, increment of densification improves the electrical conductivity on the other hand. Consequently, the SiCp(Ni)/Cu composite maintains a fairly high electrical conductivity as the uncoated one.
Fig. 1. Micrographs of (a) SiC particles after electroless nickel coating and (b) microstructure of copper matrix composite reinforced with 10 vol.% SiC particles with nickel coating.
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Fig. 2. Stress – strain curves of SiCp/Cu and SiCp(Ni)/Cu composites.
3.2. Mechanical property Typical bending stress –strain curves of the nickel coated and uncoated SiC particle composites are shown in Fig. 2. It is clear that the SiCp(Ni)/Cu composite exhibits a better combination of flexural strength and ductility than the SiC/Cu composite. The ultimate flexural strength and corresponding strain of the coated SiC composite are 445.9 MPa and 27.3%, while in the case of the uncoated SiC reinforced composite are 381.5 MPa and 24.6%, respectively. The fracture surface of the two composites after the bending test is shown in Fig. 3. It can be found that the fracture mechanism of the SiCp/Cu composite
includes the ductile rupture at copper matrix and the debonding of the SiC – Cu interface (as shown in Fig. 3a). Pulling out of SiC particles and big dimples with and without particles in them can be clearly seen on the fracture surface. It may be concluded that the poor bonding strength of the SiC – Cu interface is the primary factor responsible for the relatively low flexural strength and ductility. The fracture appearance of the SiCp(Ni)/Cu composite is much different from the former one. It is not as even as that of the SiCp/Cu composite in both low and high magnification views. Debonding of the SiC –Cu interface is lessened, and particle pulling out from the copper matrix is not obvious on the fracture surface. Severe ductile deformation in the matrix is the main fracture mechanism of this composite. Fig. 4 presents the typical interface of the two composites. Impurities can be clearly seen on the SiC – Cu interface in the case of uncoated SiC-reinforced composite (Fig. 4a). SEM and EDAX analysis of the original surface of the SiC particle show that a small quantity of dissociative Fe and O can be found. These impurities may come from the environment during the producing process of SiC powder. Furthermore, at some places of the interface, interspace between SiC and the matrix can be clearly seen (Fig. 4b). The interface of SiCp(Ni)/Cu composite, however, is smooth and compact. Copper matrix integrates with the reinforcement closely, and impurities can hardly be found (Fig. 4c). During the electroless plating process of the SiC powder, impurities on
Fig. 3. SEM photographs showing the fracture surfaces the composites. (a) SiCp/Cu composite and (b) SiCp(Ni)/Cu composites.
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gram, it is indicated that an infinite solid solution with ultimate tensile strength which is much higher than that of the pure copper is formed by the two metals [11]. Since the preparation temperature of the composites (820 jC) is high enough for the Cu and Ni atoms to diffuse to form a film of Cu – Ni alloy at the interface, the component of this film changes continuously from the copper side to the nickel side. The bonding strength of this SiC –Ni – Cu interface is high enough to pass load from the matrix to the reinforcement effectively. As a result, in the bending test, debonding of the interface is more difficult to take place than in the uncoated SiC-reinforced composite. In other words, the nickel coating on the SiC reinforcement improves the strength and ductility of the composite by lessening the extent of interfacial debonding. 3.3. Wear property The wear rate of the nickel coated and uncoated SiC particle composites is plotted as a function of the normal load in Fig. 5. An increase in wear rate with increasing normal load is observed for both the materials. However, the nickel-coated SiC-reinforced composite exhibits lower wear rate than the uncoated one at all the load levels. With the increment of normal load, the wear rate of the SiCp(Ni)/Cu composite increases slowly and keeps a rather low value. Therefore, nickel coating on SiC particle surface
Fig. 4. TEM micrographs showing the characteristics of the interfaces of the composites. (a) Impurities and (b) interspace on the SiC – Cu interface; (c) SiC – Ni – Cu interface.
the SiC particle surface were eliminated by the activation and sensitization processes to get a high bonding strength between the nickel coating and the reinforcement. From the Cu –Ni binary phase dia-
Fig. 5. The variation of wear rates with the applied load for SiCp/Cu and SiCp(Ni)/Cu composites.
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Fig. 6. SEM photographs of worn surfaces of the composites (100 N). (a) Low magnification view of the SiCp/Cu composite; (b) high magnification of the SiCp/Cu composite; (c) low magnification view of the SiCp(Ni)/Cu composite.
Fig. 7. Morphology of wear debris generated from the composites tested at 100 N: (a) SiCp/Cu composite; (b) SiCp(Ni)/Cu composite.
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improves the wear resistance of the composite, especially at high load levels. Fig. 6 presents the surface morphologies of the samples worn at 100 N. It can be clearly seen that the worn surface of the uncoated SiC-reinforced composite is largely composed of deep grooves parallel to the sliding direction and low-lying patches (Fig. 6a). These patches are essentially composed of shear dimples, which indicate that they are the places where the wear debris detaches from the worn surface. High magnification view of a groove (Fig. 6b) shows the presence of black pits and microcrack perpendicular to the sliding direction. A pit should be the original place for a SiC particle and is formed when a particle is pulled out from the matrix during the wear process.
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The worn surface of the coated SiC-reinforced composite is relatively smooth, and some discontinuousand-slim grooves can be found, as is shown in Fig. 6c. Many little shear dimples on it suggest severe deformation in the subsurface region. In Fig. 7, the morphologies of the wear debris generated from the two composites tested at 100 N are compared. In the case of the SiCp/Cu composite (Fig. 7a), the size of some of the debris is about 40 –50 Am, approximates that of the low-lying patches on the worn surface. The debris is large and flake-like, indicating that large pieces of material had delaminated from the contact surface under this load level. The wear debris of the SiCp(Ni)/Cu composite is small (15 –20 Am) and shows an equiaxed morphology(Fig. 7b).
Fig. 8. Longitudinal cross-section perpendicular to the worn surface showing subsurface region of the composites and EDAX result of the MML layer (tested at 100 N). (a) Low magnification view showing the cracks beneath the worn surface of SiCp/Cu composite; (b) subsurface region of the SiCp(Ni)/Cu composite; (c) EDAX spectrum of the MML layer in SiCp/Cu composite.
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Longitudinal cross-sections perpendicular to worn surface showing the subsurface region are presented in Fig. 8a and b to analyse the wear mechanism of the composites. For both the composites, it is clearly seen that a tribolayer which is not uniform in thickness covers the worn surface as a ‘‘protective layer’’. This kind of layer has been observed on the worn surface of aluminum matrix composites and is called the mechanically mixed layer (MML)[12]. EDAX analysis of the MML layer on the SiC/Cu composite indicates that it contains such elements as Fe, Cu, C, Si and O (Fig. 8c). Fe element is rich in the layer, and it must come from the surface of the counterface when the test sample slid against the steel ring. The presence of O implies that Fe and Cu must have been oxidized, for the flash temperature on the contact surface is often high enough to reach the oxidation temperature of the fine wear debris [13]. Subsurface cracks can be clearly seen just beneath the MML layer in the SiCp/Cu composites (Fig. 8a). These cracks were formed around the SiC particles and are apt to extend to the worn surface along the sliding direction. In the wear test, the worn surface of the sample was rubbed by counterface repeatedly. Thereupon, large strains and strain gradients were generated in the subsurface region. As mentioned above, the SiC– Cu interface of this composite is so weak that interfacial debonding takes place when the stress attains the critical bonding strength. In other words, the interface is a preferential place for the crack nucleation and propagation. Crack propagation around a SiC particle which is proud on the worn surface can lead to the direct scaling off (Fig. 6b). However, if the crack is formed in the subsurface region which is far away from the contact surface, it will propagate to the nearby matrix and connect with the other cracks (Fig. 8a). When the two sides of the crack finally extend to reach the contact surface, a piece of composite debris delaminates from the bulk material. It should be noted that the delamination of the MML layer is also an important mechanism for this material. In the case of the SiCp(Ni)/Cu composite, the nickel coating improves the bonding strength of the interface. Interfacial debonding and microcrack propagation are not so easy to take place under the same stress condition (Fig. 8b). Therefore, the SiC particles act as load-bearing components and lessen the strain and
strain gradient in the subsurface region, which mitigates the wear loss of the composite effectively. The main wear mechanism, however, is the microcutting effect of the counterface asperity on the worn surface, and adhesion of the composite to steel induced by the ductile deformation in the subsurface region. In a word, interfacial modifying is effective to improve the wear resistance of copper matrix composites by preventing microcrack nucleation in the interface. 4. Conclusions The major conclusions resulting from the work presented in this paper can be listed as follows. 1. A uniform nickel surface can be formed on the SiC powders with electroless plating process. Compact composites with coated or uncoated SiC particles dispersed uniformly in the matrix can be fabricated successfully by the powder metallurgy plus hot extrusion method. 2. Relative density and bulk hardness of nickel-coated SiC-reinforced composite are higher than that of the uncoated one, while electrical conductivity is not significantly deteriorated by the introduction of coating. 3. The SiCp(Ni)/Cu composite exhibits better combination of flexural strength and ductility than the SiCp/Cu composite. The nickel coating on the SiC reinforcement improves the mechanical properties by lessening the extent of interfacial debonding. 4. The nickel coating modifies the interface structural model and changes the wear mechanism. It is in favor of the load-bearing effect of the reinforcement, thus improves the wear resistance of the composite. References [1] J.B. Correia, H.A. Davies, C.M. Sellars, Acta Mater. 45 (1997) 177. [2] M.A. Morris, D.G. Morris, Acta Metall. 35 (1987) 2511. [3] W.C. Harrigan Jr., Mater. Sci. Eng. A244 (1998) 75. [4] Y.B. Liu, S.C. Lim, L. Lu, M.O. Lai, J. Mater. Sci. 29 (1994) 1999. [5] R. Murakami, K. Matsui, Wear 201 (1996) 193. [6] P. Nikolopoulos, S. Agathopoulos, G.N. Angelopoulos, J. Mater. Sci. 27 (1992) 139.
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R. Warren, C.-H. Andersson, Composites 15 (1984) 101. H.-K. Lee, J.-K. Lee, J. Mater. Sci. Lett. 11 (1992) 550. C.-D. Qin, B. Derby, Br. Ceram., Trans. J. 90 (1991) 124. China National Standards, GB12444.2-90, Standards Press of China, Beijing, 1994.
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[11] S.S. Tan, Nonferrous Materials, Metallurgy Industry Press, Beijing, 1993, p. 51. [12] B. Venkataraman, G. Sundararajan, Acta Mater. 44 (1996) 461. [13] S.C. Lim, M.F. Ashby, Acta Metall. 35 (1987) 8.