- Email: [email protected]
Cu–Al composite
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 1 9 7 ( 2 0 0 8 ) 43–48
journal homepage: www.elsevier.com/locate/jmatprotec
Characterization of a powder metallurgy SiC/Cu–Al composite Hailong Wang a,b , Rui Zhang a,∗ , Xing Hu a , Chang-An Wang b , Yong Huang b a b
Laboratory of Material Physics of the Ministry of Education of China, Zhengzhou University, Henan 450002, PR China The State Key Laboratory of New Ceramics and Fine Processing, Tsinghua University, Beijing 100084, PR China
a r t i c l e
i n f o
a b s t r a c t
Article history:
SiC particulate-reinforced Al composites were prepared by powder metallurgy (PM) method
Received 1 July 2006
and conventional atmospheric sintering. Scanning electron microscope (SEM), X-ray diffrac-
Received in revised form
tion (XRD) techniques were used to characterize the sintered composites. The effect of
9 February 2007
temperature on the density, hardness, strength, and microstructure of composites was
Accepted 2 June 2007
investigated. Detailed failure behavior was analyzed. It was found that the segregation of SiC appeared at higher temperature. The highest micro-hardness of 80 MPa occurred at 700 ◦ C. The strength tended to increase with the increasing temperature due to the formation of
Keywords:
Al2 Cu. Both ductile and brittle fracture features were observed.
SiC/Cu–Al
© 2007 Published by Elsevier B.V.
Composite Microstructure Density Powder metallurgy
1.
Introduction
SiC/Al composites have recently received particular interests due to their high specific modulus, high strength, and high thermal stability. They can be widely used in the aerospace, automobiles industry such as electronic heat sinks, automotive drive shafts, ground vehicle brake rotors, jet fighter aircraft fins or explosion engine components, etc. (Degischer et al., 2001). The physical and chemical compatibility between SiC particle and Al-matrix is the main concern in the preparation of SiC/Al composites (Zhang et al., 2003). Si, Mg, Nb, Ti, Cu, and some rare earth elements cannot only enhance the bonding between SiC and Al but also increase the matrix strength (Gu et al., 1999; Starink et al., 1999; Lee et al., 2001; Lima et al., 1999). The oxidation layer of SiO2 on the surface of SiC particles was reported to facilitate the wetting between SiC and Al at the interface (Narushima et al., 2003; Tang et al., 2001).
∗
Corresponding author. Tel.: +86 371 67781596. E-mail address: [email protected] (R. Zhang). 0924-0136/$ – see front matter © 2007 Published by Elsevier B.V. doi:10.1016/j.jmatprotec.2007.06.002
Coating technique, with which a thin layer of metal is coated on the surface of SiC, also performs to improve the interfacial behavior (Zhang et al., 2004a). Chemical vapor deposition (CVD) and physical vapor deposition (PVD) were used to realize coating process. However, they are too costly and complicated to be used in practical applications (Varadarajan et al., 2001). The particle reinforced metal matrix composites can be synthesized by such methods as standard ingot metallurgy (IM), powder metallurgy (PM), disintegrated melt deposition (DMD) technique, spray atomization and co-deposition approach. Different method results in different properties. The PM processing route is generally preferred since it shows a number of product advantages. The uniform distribution of ceramic particle reinforcements is readily realized. On the other hand, the solid-state process minimizes the reactions between the metal matrix and the ceramic reinforcement, and thus enhances the bonding between the reinforcement
44
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 1 9 7 ( 2 0 0 8 ) 43–48
and the matrix. However, the coefficient of thermal expansion (CTE) mismatch between the reinforcement and the matrix will give rise to high residual stress, which leads to the low tensile ductility of the composite (Xu et al., 1999; Liu and Shang, 1996; Manoharan and Gupta, 1999; Hong and Chung, 1995; Lee and Lee, 2000; Kim et al., 2001; Zhang and Wang, 2005). In this study, Cu coated SiC particles are used to reinforce Al matrix. A PM method is carried out to prepare SiC/Al composites. Some properties of the SiC/Al composites are characterized.
2.
Experimental procedures
-SiC particles (4–10 m in diameter; China White Dove Group) and Al particles (100 m in diameter; China Great-Wall Aluminum Group) were commercially available. The coated SiC/Cu particles were prepared by an electroless coating process. Cu2 O was detected in the coated SiC/Cu particles due to the instantaneous oxidation of Cu during the coating process (Zhang et al., 2004b). Details of the coating process were previously reported in our early work (Zhang et al., 2004a). The aluminum powder and the 50 vol.% SiC–50 vol.% Cu coated particles were mixed by ball-milling. The amount of the coated SiC/Cu particles was 10 wt.%. The mixture was cold pressed at 180 MPa. The pressed samples were sintered in flowing argon atmosphere at 650, 700, 750, and 800 ◦ C for 2 h, respectively. The heating rate was about 200 ◦ C/h. The sintered compacts were ground and polished. The morphology of the sintered composites was observed using an optical microscope (OLYMPUS, BHT-M, Japan) and field-emission scanning electron microscope (SEM; JEOL JSM- 6700F). Phases in the samples were determined by the X-ray diffraction analysis (XRD; D/MAX-2550V). The density measurement was carried out with the Archimedes method. The micro-hardness of the composites was measured using a diamond Vickers hardness tester (AVK-A Japan). The flexural strength was measured by a three-point bending test with a Zwick/Roell multifunction machine (Z030). The span was 30 mm and a crosshead speed of 0.5 mm/min. The samples were 3 mm × 4 mm × 36 mm in dimension with the tensile surface polished with diamond paste.
3.
Fig. 1 – XRD patterns of the composite sintered at different temperatures.
of Al with Cu lowers the densification of the SiC/Cu–Al composites. The Vickers’ hardness of the composites is illustrated in Fig. 3. The hardness of the SiC/Cu–Al composites increases with the sintering temperature. A maximum hardness of 80 MPa appears at 700 ◦ C. This might be associated with the changes in the microstructure with the sintering temperature. As shown in Fig. 4, the morphologies of the sintered SiC/Cu–Al composites are different at different sintering temperature. At 650 ◦ C discrete SiC grains primarily locate at the grainboundary. A transient layer can be distinguished between SiC and Al matrix under an optical microscope according to the contrast. The transient layer shows brown color and might be the coating layer of Cu on SiC grains. The existence of Cu2 O facilitates the wetting between SiC and Cu due to the formation of SiO2 –Cu2 O eutectic with SiO2 on the surface of SiC particles. Meanwhile, Cu2 O offers barrier for the diffusion of
Results and discussion
Fig. 1 shows XRD patterns of the composites sintered at different temperatures. Cu and Cu2 O are detected at 650 ◦ C. Cu2 O is formed during coating process as the oxidation of copper in water (Zhang et al., 2004b). This implies that no chemical change occurred in the sintered SiC/Cu–Al compacts at lower temperature. As temperature increases, Cu2 O decomposes (Zhang et al., 2003). Al melts at above 660 ◦ C and will flow to cover the SiC–Cu coated particles. At 800 ◦ C, reaction between Cu and the molten Al occurs. As a result, Al2 Cu is detected in the composites. Fig. 2 shows the changes in density of SiC/Cu–Al composites sintered at different temperatures. No obvious difference is observed at different sintering temperature. However, it shows a weak decreasing trend with increasing temperature. This implies that the decomposition of Cu2 O or the reaction
Fig. 2 – Density of SiC/Cu–Al composites vs. sintering temperature.
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 1 9 7 ( 2 0 0 8 ) 43–48
Fig. 3 – Microhardness of the SiC/Cu–Al composites vs. sintering temperature.
Al into Cu and limits the reaction of Al with Cu. The densification proceeds via physical wetting. Compared to the original Al particles, grain growth phenomenon is obviously observed (∼400 m). Since Al cannot melt at such low temperature of 650 ◦ C, the grain growth might proceed via solid-state atomic diffusion. No intra-granular SiC grain is found in Fig. 4(a). Correspondingly, the composites shows relatively low densification (Fig. 2). At high temperatures, Al melts and diffuses
45
along the interface between SiC and Al. Fig. 4(b) shows the morphology of SiC/Cu–Al sintered at 700 ◦ C. It is found that very thin grain boundary appears between two adjacent Al grains. SiC grains disperse homogeneously in the boundary region or inside the grown Al grains. The relatively homogeneous microstructure leads to the highest hardness in Fig. 3. Fig. 4(c) shows the morphology of SiC/Cu–Al at 750 ◦ C. SiC particles are observed both at the grain boundary and in the inner of Al grains. The width of grain boundary is wider than that in Fig. 4(b). The reason for this is that Cu2 O decomposes to suppress the eutectic of SiO2 –Cu2 O formation. The molten Al at 750 ◦ C contacts with Cu. Due to the fairly high surface tension, the molten Al tends to agglomerate with SiC grains segregating around in the grain boundary region. The rapid diffusion of the molten Al leads to the formation of intra-granular SiC/Cu grains inside Al grains. At 800 ◦ C, the surface tension of the molten Al markedly decreases. High-temperature viscous flow is enhanced. Some SiC grains will stay in situ inside the Al matrix to form intra-granular grains. However, most SiC would separate from the melting Al or Cu due to the poor wetting ability between each other, and tends to segregate at the tri-granular region, resulting in the heterogeneous microstructure. Additionally, reaction between Al and Cu occurs and results in the formation of Al2 Cu, as indicated in Fig. 1. Even though heterogeneous microstructure appears, the hardness of the composites increases. This abnormal change might be ascribed to the substantial intra-granular SiC grains. Fig. 5 shows the changes in the flexural strength versus the sintering temperature. The strength tends to increase
Fig. 4 – Morphology of SiC/Cu–Al composite sintered at different temperatures: (a) 650 ◦ C; (b) 700 ◦ C; (c) 750 ◦ C; (d) 800 ◦ C.
46
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 1 9 7 ( 2 0 0 8 ) 43–48
Fig. 5 – Bending strength of the SiC/Cu–Al compacts vs. sintering temperature.
with increasing temperature. It is only 33 MPa at 650 ◦ C, but increases to 140 MPa at 800 ◦ C. Normally, the failure of metal matrix composites corresponds to three possible processes, e.g., particles fracture, interface de-bonding and matrix material exhaustion (Park and Mohamed, 1995). They are related not only to the concentrations of the stress and strain components, but also to the transferring of stress from the matrix to particles, as well as several properties of each phase (Mohamed, 1997; Xu and Watt, 1996). The noticeable difference of fracture behavior is related to the microstructure and corresponds to the failure mechanisms. Cracks initiates in local area, and the peak stress or strain levels develop in the composite will control the failure process. Fig. 6 exhibits typical fracture surfaces of the SiC/Cu–Al composites sintered at different temperatures. It is noticed that the fracture behavior of the composites sintered at 650 ◦ C is inter-granular fracture. Fracture occurs at interface between the reinforcement particles and the matrix. This implies the weak bonding between
the reinforcement particles and the matrix. The crack tends to proceed along the weak-bonding interface. As a result, the sample is damaged by interface de-bonding. Fig. 6(b) shows the failure behavior of the composites sintered at 750 ◦ C. Flocculation phenomenon appears. The failure behavior has reminiscence of both locally ductile and brittle mechanism, indicating a strong interfacial bonding (Srivatsan and Prakash, 1995). When an elastic inclusion of rigid SiC particle is embedded in a plastic matrix, tensile stress arises at the interface. Meanwhile, the difference in thermal expansion coefficient between SiC and Al matrix also contributes to the interfacial stress. Therefore, dislocation occurs at the interface, leading to the strengthening of the composites (Samuel et al., 1995). At higher temperature of 800 ◦ C, many dimples are observed in the fractured surface, as shown in Fig. 7(a), which indicates a ductile failure process. The dimples are formed by the pulling out of the reinforcement clusters. As for the SiC segregation region, both ductile and brittle features are revealed, as shown in Fig. 7(b). Broken SiC particle is also observed in Fig. 7(c). According to Clyne and Withers (Clyne and Withers, 1993), two distinct processes of the reinforcement clusters are probably present during plastic deformation. First, the clusters can deform collectively as a group somewhat like a single hard particle, so that the deformation within the matrix at the heart of the cluster is much less than in the composite as a whole. Such process leads to the formation of dimples. On the other hand, the tensile hydrostatic stresses in the matrix proposed by Prangnell et al. (Prangnell et al., 1996) will be relaxed by diffusion and void nucleation, resulting in the ductile fracture. While in the second process, the particles behave independently, so that the deformation within the cluster is much greater than in the composite overall. Higher stresses might be expected in both the matrix and particles in the clustered regions (Corbin and Wilsons, 1994). The particles in the clusters will bear higher triaxial stress than the matrix (Watt et al., 1996). Then the SiC particles will be cracked by the abnormal interfacial shear stress despite of its very high intrinsic strength.
Fig. 6 – Morphology of fracture surface of the compacts sintered at (a) 650 ◦ C; (b) 750 ◦ C.
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 1 9 7 ( 2 0 0 8 ) 43–48
47
Fig. 7 – Fractographs of the compacts sintered at 800 ◦ C.
4.
Conclusion
SiC/Cu–Al composites can be prepared by using the conventional pressureless sintering. The segregation of SiC appears at higher temperature. Reaction between Al and Cu is detected at above 750 ◦ C to form Al2 Cu. The highest microhardness of 80 MPa occurs at 700 ◦ C. The strength tends to increase with the increasing temperature due to the formation of Al2 Cu. Both ductile and brittle fracture features were observed.
Acknowledgements The authors would thank the Laboratory of Material Physics of the Ministry of Education of China, Zhengzhou University for their help, and also thank the State Key Laboratory of New Ceramics and Fine Processing, Tsinghua University for their support.
references
Clyne, T.W., Withers, P.J., 1993. An Introduction to Metal Matrix Composites. Cambridge University Press, Cambridge. Corbin, S.F., Wilsons, D.S., 1994. The influence of particle distribution on the mechanical response of a particulate metal matrix composite. Acta Metall. Mater. 42, 1311–1318. Degischer, H.P., Peter, P., Christopher, S.M., 2001. Assessment of metal matrix composites for innovations—intermediate report of a European Thematic Network. Compos.: Part A 32, 1161–1166. Gu, M.Y., Zhi, M., Jin, Y.P., 1999. Structure and amorphization of the oxide on the silicon carbide surface in an SiCp/Al composite. Scripta Materialia 40, 985–991. Hong, S.H., Chung, K.H., 1995. Effects of vacuum hot pressing parameters on the tensile properties and microstructures of SiC–2124 Al composites. Mater. Sci. Eng. A 194, 165–170. Kim, B.G., Dong, S.L., Park, S.D., 2001. Effects of thermal processing on thermal expansion coefficient of a 50 vol.% SiCp/Al composite. Mater. Chem. Phys. 72, 42–47.
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Lee, T.W., Lee, C.H., 2000. Statistical evaluation of strength for die-cast SiCp/Al alloy composites. Mater. Lett. 46, 93–97. Lee, H.S., Yeo, J.S., Hong, S.H., 2001. The fabrication process and mechanical properties of SiCp/Al–Si metal matrix composites for automobile air-conditioner compressor pistons. J. Mater. Process. Technol. 113, 202–208. Lima, S.C., Gupta, M., Ren, L., 1999. The tribological properties of Al–Cu/SiCp metal–matrix composites fabricated using the rheocasting technique. J. Mater. Process. Technol. 89–90, 591–596. Liu, G., Shang, J.K., 1996. Fatigue crack tip opening behavior in particulate reinforced Al-alloy composites. Acta Mater. 44, 79–91. Manoharan, M., Gupta, M., 1999. Effect of silicon carbide volume fraction on the work hardening behaviour of thermomechanically processed aluminium-based metal–matrix composites. Compos.: Part B 30, 107–112. Mohamed, F.A., 1997. On the creep strengthening of SiC particulates in SiC–Al composites. Metall. Mater. Trans. A 28, 2780–2782. Narushima, T., Goto, T., Maruyama, M., 2003. Oxidation of boron carbide–silicon carbide composite at 1073 to 1773 K. Mater. Trans. 44, 401–406. Park, K.T., Mohamed, F.A., 1995. Creep strengthening in a discontinuous SiC–Al composite. Metall. Mater. Trans. A 26, 3119–3129. Prangnell, P.B., Barnes, S.J., Roberts, S.M., 1996. The effect of particle distribution on damage formation in particulate reinforced metal matrix composites deformed in compression. Mater. Sci. Eng. A 220, 41–56. Samuel, A.M., Gotmare, A., Samuel, F.H., 1995. Effect of solidification rate and metal feedability on porosity and SiC/Al2 O3 particle distribution in an Al–Si–Mg (359) alloy. Compos. Sci. Technol. 53, 301–315.
Srivatsan, T.S., Prakash, A., 1995. The quasi-static fracture behavior of an aluminum alloy metal–matrix composite. Compos. Sci. Technol. 53, 307–315. Starink, M.J., Wang, P., Sinclair, I., 1999. Microstructure and strengthening of Al–Li–Cu–Mg alloys and MMCs. I. Analysis and modelling of microstructural changes. Acta Mater. 47, 3841–3853. Tang, W.M., Zheng, Z.X., Ding, H.F., 2001. Oxidation of SiC and its effect on the interface stability of SiC/Fe. J. Inorg. Mater. (in Chinese) 16, 289–296. Varadarajan, S., Pattanaik, A.K., Sarin, V.K., 2001. Mullite interfacial coatings for SiC fibers. Surf. Coat. Technol. 139, 153–160. Watt, D.F., Xu, X.Q., Lloyd, D.J., 1996. Effects of particle morphology and spacing on the strain fields in a plastically deforming matrix. Acta Mater. 44, 789–799. Xu, Q., Hayes, R.W., Huntjr, W.H., 1999. Mechanical properties and fracture behavior of layered 6061/SiCp composites produced by spray atomization and co-deposition. Acta Mater. 47, 43–53. Xu, X.Q., Watt, D.F., 1996. A numerical analysis of the effects of reinforcement content on strength and ductility in Al/(SiC)p MMCs. Acta Mater. 44, 4501–4511. Zhang, R., Gao, L., Guo, J.K., 2003. Thermodynamic behavior of copper-coated silicon carbide particles during conventional heating and spark plasma sintering. J. Am. Ceram. 86, 1446–1448. Zhang, R., Gao, L., Guo, J.K., 2004a. Preparation and characterization of coated nanoscale Cu/SiCp composite particles. Ceram. Int. 30, 401–404. Zhang, R., Gao, L., Guo, J.K., 2004b. Effect of Cu2 O on the fabrication of Cu/SiCp nanocomposites using coated particles and conventional sintering. Compos.: Part A 35, 1301–1305. Zhang, R., Wang, H.L., 2005. Spheroid growth during sintering of copper coated silicon carbide particles in the fabrication of nanocomposite. Key Eng. Mater. 280–283, 1275–1278.
journal homepage: www.elsevier.com/locate/jmatprotec
Characterization of a powder metallurgy SiC/Cu–Al composite Hailong Wang a,b , Rui Zhang a,∗ , Xing Hu a , Chang-An Wang b , Yong Huang b a b
Laboratory of Material Physics of the Ministry of Education of China, Zhengzhou University, Henan 450002, PR China The State Key Laboratory of New Ceramics and Fine Processing, Tsinghua University, Beijing 100084, PR China
a r t i c l e
i n f o
a b s t r a c t
Article history:
SiC particulate-reinforced Al composites were prepared by powder metallurgy (PM) method
Received 1 July 2006
and conventional atmospheric sintering. Scanning electron microscope (SEM), X-ray diffrac-
Received in revised form
tion (XRD) techniques were used to characterize the sintered composites. The effect of
9 February 2007
temperature on the density, hardness, strength, and microstructure of composites was
Accepted 2 June 2007
investigated. Detailed failure behavior was analyzed. It was found that the segregation of SiC appeared at higher temperature. The highest micro-hardness of 80 MPa occurred at 700 ◦ C. The strength tended to increase with the increasing temperature due to the formation of
Keywords:
Al2 Cu. Both ductile and brittle fracture features were observed.
SiC/Cu–Al
© 2007 Published by Elsevier B.V.
Composite Microstructure Density Powder metallurgy
1.
Introduction
SiC/Al composites have recently received particular interests due to their high specific modulus, high strength, and high thermal stability. They can be widely used in the aerospace, automobiles industry such as electronic heat sinks, automotive drive shafts, ground vehicle brake rotors, jet fighter aircraft fins or explosion engine components, etc. (Degischer et al., 2001). The physical and chemical compatibility between SiC particle and Al-matrix is the main concern in the preparation of SiC/Al composites (Zhang et al., 2003). Si, Mg, Nb, Ti, Cu, and some rare earth elements cannot only enhance the bonding between SiC and Al but also increase the matrix strength (Gu et al., 1999; Starink et al., 1999; Lee et al., 2001; Lima et al., 1999). The oxidation layer of SiO2 on the surface of SiC particles was reported to facilitate the wetting between SiC and Al at the interface (Narushima et al., 2003; Tang et al., 2001).
∗
Corresponding author. Tel.: +86 371 67781596. E-mail address: [email protected] (R. Zhang). 0924-0136/$ – see front matter © 2007 Published by Elsevier B.V. doi:10.1016/j.jmatprotec.2007.06.002
Coating technique, with which a thin layer of metal is coated on the surface of SiC, also performs to improve the interfacial behavior (Zhang et al., 2004a). Chemical vapor deposition (CVD) and physical vapor deposition (PVD) were used to realize coating process. However, they are too costly and complicated to be used in practical applications (Varadarajan et al., 2001). The particle reinforced metal matrix composites can be synthesized by such methods as standard ingot metallurgy (IM), powder metallurgy (PM), disintegrated melt deposition (DMD) technique, spray atomization and co-deposition approach. Different method results in different properties. The PM processing route is generally preferred since it shows a number of product advantages. The uniform distribution of ceramic particle reinforcements is readily realized. On the other hand, the solid-state process minimizes the reactions between the metal matrix and the ceramic reinforcement, and thus enhances the bonding between the reinforcement
44
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 1 9 7 ( 2 0 0 8 ) 43–48
and the matrix. However, the coefficient of thermal expansion (CTE) mismatch between the reinforcement and the matrix will give rise to high residual stress, which leads to the low tensile ductility of the composite (Xu et al., 1999; Liu and Shang, 1996; Manoharan and Gupta, 1999; Hong and Chung, 1995; Lee and Lee, 2000; Kim et al., 2001; Zhang and Wang, 2005). In this study, Cu coated SiC particles are used to reinforce Al matrix. A PM method is carried out to prepare SiC/Al composites. Some properties of the SiC/Al composites are characterized.
2.
Experimental procedures
-SiC particles (4–10 m in diameter; China White Dove Group) and Al particles (100 m in diameter; China Great-Wall Aluminum Group) were commercially available. The coated SiC/Cu particles were prepared by an electroless coating process. Cu2 O was detected in the coated SiC/Cu particles due to the instantaneous oxidation of Cu during the coating process (Zhang et al., 2004b). Details of the coating process were previously reported in our early work (Zhang et al., 2004a). The aluminum powder and the 50 vol.% SiC–50 vol.% Cu coated particles were mixed by ball-milling. The amount of the coated SiC/Cu particles was 10 wt.%. The mixture was cold pressed at 180 MPa. The pressed samples were sintered in flowing argon atmosphere at 650, 700, 750, and 800 ◦ C for 2 h, respectively. The heating rate was about 200 ◦ C/h. The sintered compacts were ground and polished. The morphology of the sintered composites was observed using an optical microscope (OLYMPUS, BHT-M, Japan) and field-emission scanning electron microscope (SEM; JEOL JSM- 6700F). Phases in the samples were determined by the X-ray diffraction analysis (XRD; D/MAX-2550V). The density measurement was carried out with the Archimedes method. The micro-hardness of the composites was measured using a diamond Vickers hardness tester (AVK-A Japan). The flexural strength was measured by a three-point bending test with a Zwick/Roell multifunction machine (Z030). The span was 30 mm and a crosshead speed of 0.5 mm/min. The samples were 3 mm × 4 mm × 36 mm in dimension with the tensile surface polished with diamond paste.
3.
Fig. 1 – XRD patterns of the composite sintered at different temperatures.
of Al with Cu lowers the densification of the SiC/Cu–Al composites. The Vickers’ hardness of the composites is illustrated in Fig. 3. The hardness of the SiC/Cu–Al composites increases with the sintering temperature. A maximum hardness of 80 MPa appears at 700 ◦ C. This might be associated with the changes in the microstructure with the sintering temperature. As shown in Fig. 4, the morphologies of the sintered SiC/Cu–Al composites are different at different sintering temperature. At 650 ◦ C discrete SiC grains primarily locate at the grainboundary. A transient layer can be distinguished between SiC and Al matrix under an optical microscope according to the contrast. The transient layer shows brown color and might be the coating layer of Cu on SiC grains. The existence of Cu2 O facilitates the wetting between SiC and Cu due to the formation of SiO2 –Cu2 O eutectic with SiO2 on the surface of SiC particles. Meanwhile, Cu2 O offers barrier for the diffusion of
Results and discussion
Fig. 1 shows XRD patterns of the composites sintered at different temperatures. Cu and Cu2 O are detected at 650 ◦ C. Cu2 O is formed during coating process as the oxidation of copper in water (Zhang et al., 2004b). This implies that no chemical change occurred in the sintered SiC/Cu–Al compacts at lower temperature. As temperature increases, Cu2 O decomposes (Zhang et al., 2003). Al melts at above 660 ◦ C and will flow to cover the SiC–Cu coated particles. At 800 ◦ C, reaction between Cu and the molten Al occurs. As a result, Al2 Cu is detected in the composites. Fig. 2 shows the changes in density of SiC/Cu–Al composites sintered at different temperatures. No obvious difference is observed at different sintering temperature. However, it shows a weak decreasing trend with increasing temperature. This implies that the decomposition of Cu2 O or the reaction
Fig. 2 – Density of SiC/Cu–Al composites vs. sintering temperature.
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 1 9 7 ( 2 0 0 8 ) 43–48
Fig. 3 – Microhardness of the SiC/Cu–Al composites vs. sintering temperature.
Al into Cu and limits the reaction of Al with Cu. The densification proceeds via physical wetting. Compared to the original Al particles, grain growth phenomenon is obviously observed (∼400 m). Since Al cannot melt at such low temperature of 650 ◦ C, the grain growth might proceed via solid-state atomic diffusion. No intra-granular SiC grain is found in Fig. 4(a). Correspondingly, the composites shows relatively low densification (Fig. 2). At high temperatures, Al melts and diffuses
45
along the interface between SiC and Al. Fig. 4(b) shows the morphology of SiC/Cu–Al sintered at 700 ◦ C. It is found that very thin grain boundary appears between two adjacent Al grains. SiC grains disperse homogeneously in the boundary region or inside the grown Al grains. The relatively homogeneous microstructure leads to the highest hardness in Fig. 3. Fig. 4(c) shows the morphology of SiC/Cu–Al at 750 ◦ C. SiC particles are observed both at the grain boundary and in the inner of Al grains. The width of grain boundary is wider than that in Fig. 4(b). The reason for this is that Cu2 O decomposes to suppress the eutectic of SiO2 –Cu2 O formation. The molten Al at 750 ◦ C contacts with Cu. Due to the fairly high surface tension, the molten Al tends to agglomerate with SiC grains segregating around in the grain boundary region. The rapid diffusion of the molten Al leads to the formation of intra-granular SiC/Cu grains inside Al grains. At 800 ◦ C, the surface tension of the molten Al markedly decreases. High-temperature viscous flow is enhanced. Some SiC grains will stay in situ inside the Al matrix to form intra-granular grains. However, most SiC would separate from the melting Al or Cu due to the poor wetting ability between each other, and tends to segregate at the tri-granular region, resulting in the heterogeneous microstructure. Additionally, reaction between Al and Cu occurs and results in the formation of Al2 Cu, as indicated in Fig. 1. Even though heterogeneous microstructure appears, the hardness of the composites increases. This abnormal change might be ascribed to the substantial intra-granular SiC grains. Fig. 5 shows the changes in the flexural strength versus the sintering temperature. The strength tends to increase
Fig. 4 – Morphology of SiC/Cu–Al composite sintered at different temperatures: (a) 650 ◦ C; (b) 700 ◦ C; (c) 750 ◦ C; (d) 800 ◦ C.
46
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 1 9 7 ( 2 0 0 8 ) 43–48
Fig. 5 – Bending strength of the SiC/Cu–Al compacts vs. sintering temperature.
with increasing temperature. It is only 33 MPa at 650 ◦ C, but increases to 140 MPa at 800 ◦ C. Normally, the failure of metal matrix composites corresponds to three possible processes, e.g., particles fracture, interface de-bonding and matrix material exhaustion (Park and Mohamed, 1995). They are related not only to the concentrations of the stress and strain components, but also to the transferring of stress from the matrix to particles, as well as several properties of each phase (Mohamed, 1997; Xu and Watt, 1996). The noticeable difference of fracture behavior is related to the microstructure and corresponds to the failure mechanisms. Cracks initiates in local area, and the peak stress or strain levels develop in the composite will control the failure process. Fig. 6 exhibits typical fracture surfaces of the SiC/Cu–Al composites sintered at different temperatures. It is noticed that the fracture behavior of the composites sintered at 650 ◦ C is inter-granular fracture. Fracture occurs at interface between the reinforcement particles and the matrix. This implies the weak bonding between
the reinforcement particles and the matrix. The crack tends to proceed along the weak-bonding interface. As a result, the sample is damaged by interface de-bonding. Fig. 6(b) shows the failure behavior of the composites sintered at 750 ◦ C. Flocculation phenomenon appears. The failure behavior has reminiscence of both locally ductile and brittle mechanism, indicating a strong interfacial bonding (Srivatsan and Prakash, 1995). When an elastic inclusion of rigid SiC particle is embedded in a plastic matrix, tensile stress arises at the interface. Meanwhile, the difference in thermal expansion coefficient between SiC and Al matrix also contributes to the interfacial stress. Therefore, dislocation occurs at the interface, leading to the strengthening of the composites (Samuel et al., 1995). At higher temperature of 800 ◦ C, many dimples are observed in the fractured surface, as shown in Fig. 7(a), which indicates a ductile failure process. The dimples are formed by the pulling out of the reinforcement clusters. As for the SiC segregation region, both ductile and brittle features are revealed, as shown in Fig. 7(b). Broken SiC particle is also observed in Fig. 7(c). According to Clyne and Withers (Clyne and Withers, 1993), two distinct processes of the reinforcement clusters are probably present during plastic deformation. First, the clusters can deform collectively as a group somewhat like a single hard particle, so that the deformation within the matrix at the heart of the cluster is much less than in the composite as a whole. Such process leads to the formation of dimples. On the other hand, the tensile hydrostatic stresses in the matrix proposed by Prangnell et al. (Prangnell et al., 1996) will be relaxed by diffusion and void nucleation, resulting in the ductile fracture. While in the second process, the particles behave independently, so that the deformation within the cluster is much greater than in the composite overall. Higher stresses might be expected in both the matrix and particles in the clustered regions (Corbin and Wilsons, 1994). The particles in the clusters will bear higher triaxial stress than the matrix (Watt et al., 1996). Then the SiC particles will be cracked by the abnormal interfacial shear stress despite of its very high intrinsic strength.
Fig. 6 – Morphology of fracture surface of the compacts sintered at (a) 650 ◦ C; (b) 750 ◦ C.
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 1 9 7 ( 2 0 0 8 ) 43–48
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Fig. 7 – Fractographs of the compacts sintered at 800 ◦ C.
4.
Conclusion
SiC/Cu–Al composites can be prepared by using the conventional pressureless sintering. The segregation of SiC appears at higher temperature. Reaction between Al and Cu is detected at above 750 ◦ C to form Al2 Cu. The highest microhardness of 80 MPa occurs at 700 ◦ C. The strength tends to increase with the increasing temperature due to the formation of Al2 Cu. Both ductile and brittle fracture features were observed.
Acknowledgements The authors would thank the Laboratory of Material Physics of the Ministry of Education of China, Zhengzhou University for their help, and also thank the State Key Laboratory of New Ceramics and Fine Processing, Tsinghua University for their support.
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