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The effect of SiC particles on the microstructure and mechanical properties of friction stir welded pure copper joints
Materials Science and Engineering A 528 (2011) 5470–5475
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
The effect of SiC particles on the microstructure and mechanical properties of friction stir welded pure copper joints Y.F. Sun, H. Fujii ∗ Joining and Welding Research Institute, Osaka University, 11-1 Mihogaoka, Ibaraki 567-0047, Japan
a r t i c l e
i n f o
Article history: Received 5 December 2010 Received in revised form 10 March 2011 Accepted 19 March 2011 Available online 25 March 2011 Keywords: Friction stir welding Copper SiC Dynamic recrystallization Nucleation site
a b s t r a c t The SiC particles with mean size of 5 m were introduced into the pure Cu joints during the friction stir welding process. After one pass processing, the distribution of SiC particle is not uniform and some pores formed around the aggregation of SiC particles. However, the particle distribution became rather uniform after two passes processing and the previous pores can be refilled. In addition, banded structure consisting of particle-rich and particle-free region can be observed in the stir zone. The particle-rich region has an average grain size of less than 2 m, much smaller than that of about 8 m in the particlefree region. Microstructural observation confirmed that the SiC particles can act as the heterogeneous nucleation site in the dynamic recrystallization of Cu grains. The SiC dispersed Cu joints exhibit a Vickers hardness of 110 HV, much higher than 70 HV in the stir zone without SiC particles. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Friction stir welding (FSW) is a kind of solid-state joining technique, which was invented at The Welding Institute (TWI) of UK in 1991 and was originally applied to Al alloys. With the rapid development of this technique and the application of high strength and durable rotation tools, the use of FSW has been expanded to many other materials including Mg, Cu, Ti, steels and Ni alloys, etc. [1–6]. Recently, metal matrix composites (MMCs) containing various kinds of particles have been successfully fabricated by FSW technique. During the FSW process, the particles are first put into the groove in the metallic materials. With the traveling of the high rotation tools, the particles can be completely scattered and merged into the plastic flow of the metallic materials. Since the entire fabrication process is performed under solid state, the problems that often occur in casting process can be avoided, such as the chemical reaction and poor wettability between the ceramic particles and the metal melts, and the aggregation of the particles in the metal melts. For example, Morisada et al. fabricated SiC reinforced AZ31 and carbon nanotube reinforced AZ31 alloy composites by FSW technique. It was found that the joints properties can be remarkably improved due to the microstructural refinements and the presence of the particles can greatly retard the grain growth of the Mg alloy matrix at higher temperature [7,8]. Similarly, Mahmoud et al. produced SiC contained Al alloys, which show higher
∗ Corresponding author. Tel.: +81 6 68798663; fax: +81 6 68798663. E-mail address: [email protected] (H. Fujii). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.03.077
hardness than the base metal (BM) and much improved wearresistant properties [9,10]. Recently, it was found that the particles can also be used to refill the gap or pores in the workpiece by a novel technique called friction stir powder process (FSPP), which was developed based on the principle of FSW. The relationship between the width of the pre-formed gap and the mechanical properties of the sample refilled by FSPP technique was also investigated [11]. In this study, the micro-scaled SiC particles were introduced into the FSW processed Cu with the purpose to improve the mechanical properties of the Cu joints. The mechanical property of the FSW processed Cu joints is usually lower than that of the BM due to the annealing effect caused by very high heat input, even though the microstructure can be refined in the SZ due to dynamical recrystallization [12–14]. On the other hand, SiC reinforced Cu composites have been regarded as one of the best candidates for electronic packaging and thermal management application. However, the poor wettability and strong reaction between SiC and Cu melt are two critical problems standing in the way of fabricating SiC/Cu composite by using casting methods. It was observed that the contact angle was 137◦ for pure Cu droplet on sintered ␣-SiC and the interfacial reaction (Cu + SiC = Cu3 Si + C) can easily take place, which is detrimental to the thermo-physical properties due to the solid solution of Si and the formation of graphite [15–17]. From this point of view, FSW process might be a new approach to produce SiC reinforced Cu composites. After the FSW of Cu with SiC addition, the microstructure evolution and mechanical properties of the welds were investigated.
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tests were carried out using an Instron type testing machine at a crosshead speed of 1 mm/min. The tensile direction is perpendicular to the welding direction. 3. Results and discussion 3.1. Microstructure observation
Fig. 1. Sketch map showing the FSW process with the addition of SiC powder.
2. Experimental procedure Commercial purity Cu plates received in 1/2 Hard condition were used in this study as starting materials. For the FSW process, two Cu plates with 2 mm in thickness, 200 mm in length and 50 mm in width were clamped together on the back plate with a 1 mm wide gap between the two adjoining plates. Then the SiC particles with mean size of 5 m were inserted into the gap and pressed tightly. During the welding process, a WC-based alloy tool, which had a 12 mm-diameter shoulder, 4 mm-diameter probe and 2.0 mm probe length, was used with a constant applied load of 1000 kg and titled by 3◦ . A rotation speed of 500 rpm and welding speed of 100 mm/min were used during the welding process. Fig. 1 is the sketch map showing the main principle of FSW process with the addition of SiC powder. To make the SiC particles distribute uniformly in the Cu joints, the welding process were repeated at the same welding condition and regarded as the second pass welding process. After welding, optical microscopy (OM) was used to characterize the macrostructure of the joints. The samples for OM observation were cross-sectioned perpendicular to the welding direction, mechanically polished and then etched with a solution of iron chloride. The electron backscattered diffraction (EBSD) measurements were carried out by using a JEM-7001FA field emission scanning electron microscopy (FE-SEM) with a TSL orientation imaging system. For TEM sample preparation, thin plate were first cut at the target locations and then mechanically polished to a thickness of 100 m. The polished thin plates were finally twin-jet electropolished to make electron beam transparent thin film using a solution of HPO4 :CH4 O:H2 O = 1:1:2 at 5 V and 0 ◦ C. The thin films were observed with a Hitachi 800 TEM at 200 kV. The Vickers hardness profiles of the joint were measured along the centerline of the cross-section by using a Vickers indenter with a load of 0.98 N and a dwell time of 15 s. The tensile specimens were electrical discharge machined into a dog-bone shape with a gauge length of 100 mm, width of 10 mm and thickness of 2 mm, respectively. The tensile
The OM images of the Cu joints with SiC addition after one and two pass processing are shown in Fig. 2. The samples for OM observation were only mechanically polished without chemical etching. For both samples, an obvious area with dispersed particles can be observed and no obvious welding defects or superfluous flash can be found in the joints. After one pass welding process, the area with SiC distribution in the SZ is small and the particles aggregated seriously in the advancing side (AS). In addition, some particles-free regions can be observed and are surrounded by the SiC particles. While after two pass processing, the distribution of SiC particles became very homogenous and appeared in a larger area in the SZ. It was found that for both one pass and two pass processed samples, the SiC particles distribute more in the AS than in the retreating side (RS). According to A.C. Nunes’s physical model of the metal flow in the FSW joint, the material on the RS of the joint is directly transported from the front to the back of the rotation tool, with minimal residence time in the rotation field. This is referred to as the straight-through current flow of the metal. On the contrary, the materials on the AS of the joint resides long enough in the rotational flow around the tool and consists of rotation flow around the pin and downward flow to the top of the pin, which is referred to as the maelstrom current [18]. This model can well explain the asymmetrical distribution of the SiC particles in the Cu weld. The distribution of the SiC can reach the bottom of the AS and a swirl-like pattern formed in the center of the AS due to the maelstrom current of the metal during the FSW process. The gathering of the introduced particles in the AS was also found in the Al joint containing fullerene particles, which was explained to be related with the materials flow caused by the probe and the shoulder [19]. Nevertheless, the particles can be dispersed more homogeneously within an enlarged area by repeating the welding process. Fig. 3(a) and (b) shows the high magnification SEM images of the SZ after one pass and two pass processing, respectively. After one pass, the distribution of SiC particles is not homogeneous, indicating that the materials flow is not enough and the SiC particles cannot be scattered completely into the Cu matrix by the stirring of the rotation tools. The severe aggregation of the SiC particles leads to the pore formation around the particles due to the large difference of physical properties between Cu and SiC. This phenomenon was also found in the friction stir processed Al alloys with SiC particles, in which the distribution of SiC particles is not homogeneous,
Fig. 2. Cross-sectional macrostructure of (a) one pass and (b) two pass FSW processed Cu with SiC particles.
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Fig. 3. SEM micrographs showing the distribution of SiC particles in the SZ of FSW processed Cu after (a) one pass and (b) two pass.
defects at the bottom of the SZ generally forms due to the insufficient stirring during the processing [20]. However, after the second pass, the SiC particles can be uniformly dispersed into the Cu matrix. The SiC particles shows strong bonding with the Cu matrix and no pores or defects formed between Cu and SiC. From the SEM image at higher magnification inserted in the left corner of Fig. 3(b), some large SiC particles with crack inside can be observed, implying that some SiC particles can be broken into small pieces by the strong interactions with the rotation tools. The small SiC particles distributed in the Cu matrix are supposed to be fractured from the large particles. Fig. 4 shows the microstructural transition of the SiC contained Cu joint after two pass processing. Fig. 4(a) shows the microstructure of a large area of the joint, in which several typical zones, namely, BM, heat affected zone (HAZ) and SZ, can be easily distinguished and separated with dashed lines. However, the thermal-mechanical affected zone (TMAZ) cannot be easily distinguished. Fig. 4(b) shows the microstructure of BM, which exhibits a coarse equiaxial grain structure. Fig. 4(c) shows the microstructure
of HAZ, in which no SiC particles can be observed since no plastic deformation took place in this area. The main microstructural difference between the BM and HAZ is that the grain boundary in HAZ is not as obvious as that in the BM. Because the materials in HAZ experienced no plastic deformation but an annealing process during welding, it seems that the grain boundary started to immigrate into each other and have a tendency of grain growth. Fig. 4(d) shows the microstructure of the SZ, in which a banded structure can be observed. The formation of banded microstructure consisting of alternative particle-rich region and particle-free region has been reported by some investigators and it was said that the band spacing was directly related with the welding tool advance per revolution, which is also the main reason for the onion rings formation in the SZ [21]. However, it should be noted that the bands spacing in Fig. 4(d) is much smaller than the tool advance according to the revolution pitch of 0.2 mm/rev in the present study, because two pass FSW process were performed. In addition, it can be found that the grain size in the particle-free region is larger than that in the particle-rich region. However, the grain structure of the Cu matrix
Fig. 4. OM micrographs showing the microstructure transition in the SiC particles dispersed FSW processed Cu after two pass: (a) Large area OM image, (b) BM, (c) HAZ, and (d) SZ.
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Fig. 5. EBSP grain boundary maps showing the microstructure of (a) BM and (b) SZ after two pass processing. (Grey area indicates SiC particles.)
Fig. 6. TEM images showing the microstructure of the SZ after two pass processing: (a) SiC particle at the grain boundary and (b) SiC particle inside Cu grain.
in the particle-rich region cannot be clearly discerned due to the densely distributed SiC particles. In order to understand the effect of particles on the microstructure in the SZ, EBSD measurements were carried out in the SZ and the BM as well. Fig. 5(a) and (b) shows the grain boundary maps of the BM and the SZ of the two pass processed joint, respectively, in which the grain size of Cu matrix in different locations can be easily understood. The BM has a coarse grain structure with average grain size of about 16.2 m. As for SiC dispersed SZ shown in Fig. 5(b), the SiC particles can be easily distinguished due to its poor image quality and was indicated with grey color. It reveals that the banded structure consists of much refined Cu grains in the particle-rich region and relatively coarse Cu grains in the particlefree region. The particles-rich region and particles-free region have an average grain size of about 2 m and 8 m, respectively, which confirms that the particles can refine the microstructure of the FSW processed Cu. Fig. 6 shows the results of the TEM observation in the SZ of the Cu welds after two pass processing. Generally the SiC particles distributes in the Cu matrix through two kinds of ways. One is that the SiC particles distributed at the grain boundaries of Cu grains as shown in Fig. 6(a). Around the SiC particle, high density of dislocations can be observed inside the surrounding Cu grains. Another is that the SiC particle is embedded inside the Cu grain as shown in Fig. 6(b). Dislocations can also be found, however, with very low density. In addition, for all the SiC particles distributed through these two ways, a clear interface is
Fig. 7. Hardness profile along the centerline of the cross-section of the FSW processed Cu with and without SiC additions.
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Fig. 8. Tensile strength and the appearance of the fractured samples of BM and the SiC particle contained FSW Cu joint after one and two pass.
formed between the particles and the Cu matrix. No reaction or transitional layer between SiC particle and Cu matrix can be detected. It was said that during hot deformation of metallic materials containing particles, the particles generally exert three important impacts on the dynamic recrystallization of the matrix materials. One is that the particles can increase stored energy and hence the driving force for recrystallization. Another one is that the deformation heterogeneities at large particles may be sites at which recrystallization originates, which is called particle stimulated nucleation (PSN). The third one is that the particles can also pin the movement of the grain boundary and retard the grain growth after dynamic recrystallization [22]. The PSN during static recrystallization is commonly found in alloys containing hard particles such as carbides and oxides [23–26]. PSN has also been found during dynamic recrystallization in Cu containing Al2 O3 , SiO2 particles and the PSN was greatly depended on the particles size [27]. FSW is a well known severe plastic deformation process. The dynamic recrystallization usually occurs in the stir zone and thus the microstructure can be refined. In the present FSW of pure Cu, the smaller grain size in the particle-rich region is supposed to be caused by either PSN or the pinning effect. The pinning effect corresponds to the situation as shown in Fig. 6(a), the stress on the Cu grains by the particles will certainly cause the generation of high density dislocations. The PSN corresponds to the situation as shown in Fig. 6(b), the SiC particle acts as the nucleation site and the recrystallized Cu grain has a very low dislocation density.
3.2. Mechanical properties Fig. 7 shows the hardness distribution along the center line of the cross-section of the Cu joints after one and two pass processing. The hardness distribution of the Cu joint without SiC particles were also plotted for comparison. For both Cu joints dispersed with SiC particles, the hardness in the SZ is higher than that of the pure Cu counterpart. However, the hardness in the SZ of the one pass processed joint is not uniformly distributed. The low hardness value can be observed in some local areas, which corresponds to the particles-free region in the SZ. While for the joints after two pass processing, the hardness distribution in the SZ is rather homogenous. Although the banded structure consisting of particle-rich and particle-free region formed after two pass processing, the bands spacing is much smaller than the size of the indentation after the hardness tests. As a result, the hardness profile is rather smooth within the entire range of SZ. The hardness increase of the FSW processed Cu with SiC addition is supposed to be contributed by both the decrease of the average grain size and the presence of the hard SiC particles. Although it is hard to say which one of these two factors is more important to the increased hardness, the refinement of the microstructure cannot be achieved if no SiC particles were introduced into the materials during the welding process. In addition, the areas with the lowest hardness value are located in HAZ for all the joints, no matter whether the SiC particles are contained or not. Fig. 8 shows the tensile strength of the Cu joint with and without SiC additions. The one pass processed Cu joint shows a much lower
Fig. 9. Fractography of the tensile specimen after (a) one pass and (b) two pass FSW processing.
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strength of only 125 MPa and the two pass processed Cu joint shows a strength of about 217 MPa, which is very close to the 228 MPa of that without SiC additions. From the appearance of the fractured tensile specimen shown in Fig. 8, the one pass processed joints fractured in the area where the SiC greatly aggregated. However, the two pass processed joints and the joints without SiC particles fractured in the HAZ, corresponding to the lowest hardness value in the SZ. Fig. 9(a) and (b) shows the fractural morphology of the tensile specimens of SiC contained Cu joints after one and two pass processing, respectively. For the specimen after one pass processing, very large pores can be observed on the fracture plane, which is caused by the severe aggregation of SiC particles in the SZ. Although the hardness is generally high in the SZ due to the SiC reinforcement, the formation of pores might tear the continuity of the matrix and becomes the initialization site of the failure during tensile tests. As for the specimen after two pass processing, the fracture took place in the HAZ of the joint. However, the fracture should be close to the SZ, because sparsely distributed SiC particles can also be found on the fractured plane as shown in Fig. 9(b). In addition, large dimples formed due to the relatively large grain size of Cu matrix in HAZ. 4. Conclusions From the above descriptions, the following conclusions can be drawn: a) The micro-sized SiC particles can be introduced into the Cu welds by simply putting the SiC particles between the adjoining plates during the FSW process and the SiC particles can be uniformly distributed in the SZ after two pass processing. b) In the two pass processed joints, banded structure consisting of alternative particle-rich region and particle-free region formed in the SZ. The particle-rich region shows much refined grain structure with average grain size less than 2 m due to the SiC particles stimulated nucleation in the dynamic recrystallization of Cu during the FSW process, while the average grain size in the particle-free region is about 8 m. c) Comparing with the hardness of about 70 HV in the SZ of pure Cu joint without SiC addition, the hardness in the Cu welds can be remarkably increased to 110 HV by the introduction of SiC particles with a slight loss of tensile strength of 11 MPa.
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Acknowledgements The authors wish to acknowledge the financial support of a Grant-in-Aid for Science Research and the Collaborative Research Based on Industrial Demand from the Japan Society for the Promotion of Science and Technology of Japan, the Global COE Program (Project: Center of Excellence for Advanced Structural and Functional Materials Design) from the Ministry of Education, Sports, Culture, Science and Technology of Japan. References [1] R. Nandan, T. Debroy, H.K.D.H. Bhadeshia, Prog. Mater. Sci. 53 (2008) 980–1023. [2] R.S. Mishra, Z.Y. Ma, Mater. Sci. Eng. R 50 (2005) 1–78. [3] P.L. Threadgill, A.J. Leonard, H.R. Shercliff, P.J. Withers, Int. Mater. Rev. 54 (2009) 49–93. [4] S.H.C. Park, Y.S. Sato, H. Kokawa, J. Mater. Sci. 38 (2003) 4379–4383. [5] M. Ghosh, K. Kumar, R.S. Mishra, Scripta Mater. 63 (2010) 851–854. [6] S.M. Mousavizade, F.M. Ghaini, M.J. Torkamany, J. Sabbaghzadeh, A. Abdollahzadeh, Scripta Mater. 60 (2009) 244–247. [7] Y. Morisada, H. Fujii, T. Nagaoka, M. Fukusumi, Mater. Sci. Eng. A 433 (2006) 50–54. [8] Y. Morisada, H. Fujii, T. Nagaoka, M. Fukusumi, Mater. Sci. Eng. A 419 (2006) 344–348. [9] E.R.I. Mahmoud, M. Takahashi, T. Shibayanagi, K. Ikeuchi, Mater. Trans. 50 (2009) 1824–1831. [10] E.R.I. Mahmoud, M. Takahashi, T. Shibayanagi, Wear 268 (2010) 1111– 1121. [11] K. Inada, H. Fujii, Y.S. Ji, Y.F. Sun, Y. Morisada, Sci. Tech. Weld. Join 14 (2009) 41–46. [12] W.B. Lee, S.B. Jung, Mater. Lett. 58 (2004) 1041–1046. [13] Y.F. Sun, H. Fujii, Mater. Sci. Eng. A 527 (2010) 6879–6886. [14] G.M. Xie, Z.Y. Ma, L. Geng, Scripta Mater. 57 (2007) 73–76. [15] L. Zhang, X.H. Qu, X.B. He, B.H. Duan, S.B. Ren, M.L. Qin, Mater. Sci. Eng. A 489 (2008) 285–293. [16] V. Martinez, S. Ordonez, F. Castro, J. Mater. Sci. 38 (2003) 4047–4054. [17] C. Rado, B. Drevet, N. Eustathopoulou, Acta Mater. 48 (2000) 4483–4491. [18] J.A. Schneider, A.C. Nunes, Metall. Mater. Trans. A 35 (2004) 777–783. [19] Y. Morisada, H. Fujii, T. Nagaoka, K. Nogi, M. Fukusumi, Composites: Part A 38 (2007) 2097–2101. [20] E.R.I. Mahmoud, M. Takahashi, T. Shibayanagi, K. Ikeuchi, Sci. Tech. Weld. Join 14 (2009) 413–425. [21] M.A. Sutton, B. Yang, A.P. Reynolds, R. Taylor, Mater. Sci. Eng. A 323 (2002) 160–166. [22] W.L. Zhang, J.X. Wang, F. Yang, Z.Q. Sun, M.Y. Gu, J. Comp. Mater. 40 (2006) 1117–1131. [23] F.J. Humphreys, Met. Sci. 13 (1979) 136–145. [24] F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, 2nd ed., Pergamon Press, Oxford, 2002. [25] F.J. Humphreys, Scripta Mater. 43 (2000) 591–596. [26] R.L. Goetz, Scripta Mater. 52 (2005) 851–856. [27] F.J. Humphreys, M.G. Ardakani, Acta Mater. 44 (1996) 2717–2727.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
The effect of SiC particles on the microstructure and mechanical properties of friction stir welded pure copper joints Y.F. Sun, H. Fujii ∗ Joining and Welding Research Institute, Osaka University, 11-1 Mihogaoka, Ibaraki 567-0047, Japan
a r t i c l e
i n f o
Article history: Received 5 December 2010 Received in revised form 10 March 2011 Accepted 19 March 2011 Available online 25 March 2011 Keywords: Friction stir welding Copper SiC Dynamic recrystallization Nucleation site
a b s t r a c t The SiC particles with mean size of 5 m were introduced into the pure Cu joints during the friction stir welding process. After one pass processing, the distribution of SiC particle is not uniform and some pores formed around the aggregation of SiC particles. However, the particle distribution became rather uniform after two passes processing and the previous pores can be refilled. In addition, banded structure consisting of particle-rich and particle-free region can be observed in the stir zone. The particle-rich region has an average grain size of less than 2 m, much smaller than that of about 8 m in the particlefree region. Microstructural observation confirmed that the SiC particles can act as the heterogeneous nucleation site in the dynamic recrystallization of Cu grains. The SiC dispersed Cu joints exhibit a Vickers hardness of 110 HV, much higher than 70 HV in the stir zone without SiC particles. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Friction stir welding (FSW) is a kind of solid-state joining technique, which was invented at The Welding Institute (TWI) of UK in 1991 and was originally applied to Al alloys. With the rapid development of this technique and the application of high strength and durable rotation tools, the use of FSW has been expanded to many other materials including Mg, Cu, Ti, steels and Ni alloys, etc. [1–6]. Recently, metal matrix composites (MMCs) containing various kinds of particles have been successfully fabricated by FSW technique. During the FSW process, the particles are first put into the groove in the metallic materials. With the traveling of the high rotation tools, the particles can be completely scattered and merged into the plastic flow of the metallic materials. Since the entire fabrication process is performed under solid state, the problems that often occur in casting process can be avoided, such as the chemical reaction and poor wettability between the ceramic particles and the metal melts, and the aggregation of the particles in the metal melts. For example, Morisada et al. fabricated SiC reinforced AZ31 and carbon nanotube reinforced AZ31 alloy composites by FSW technique. It was found that the joints properties can be remarkably improved due to the microstructural refinements and the presence of the particles can greatly retard the grain growth of the Mg alloy matrix at higher temperature [7,8]. Similarly, Mahmoud et al. produced SiC contained Al alloys, which show higher
∗ Corresponding author. Tel.: +81 6 68798663; fax: +81 6 68798663. E-mail address: [email protected] (H. Fujii). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.03.077
hardness than the base metal (BM) and much improved wearresistant properties [9,10]. Recently, it was found that the particles can also be used to refill the gap or pores in the workpiece by a novel technique called friction stir powder process (FSPP), which was developed based on the principle of FSW. The relationship between the width of the pre-formed gap and the mechanical properties of the sample refilled by FSPP technique was also investigated [11]. In this study, the micro-scaled SiC particles were introduced into the FSW processed Cu with the purpose to improve the mechanical properties of the Cu joints. The mechanical property of the FSW processed Cu joints is usually lower than that of the BM due to the annealing effect caused by very high heat input, even though the microstructure can be refined in the SZ due to dynamical recrystallization [12–14]. On the other hand, SiC reinforced Cu composites have been regarded as one of the best candidates for electronic packaging and thermal management application. However, the poor wettability and strong reaction between SiC and Cu melt are two critical problems standing in the way of fabricating SiC/Cu composite by using casting methods. It was observed that the contact angle was 137◦ for pure Cu droplet on sintered ␣-SiC and the interfacial reaction (Cu + SiC = Cu3 Si + C) can easily take place, which is detrimental to the thermo-physical properties due to the solid solution of Si and the formation of graphite [15–17]. From this point of view, FSW process might be a new approach to produce SiC reinforced Cu composites. After the FSW of Cu with SiC addition, the microstructure evolution and mechanical properties of the welds were investigated.
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tests were carried out using an Instron type testing machine at a crosshead speed of 1 mm/min. The tensile direction is perpendicular to the welding direction. 3. Results and discussion 3.1. Microstructure observation
Fig. 1. Sketch map showing the FSW process with the addition of SiC powder.
2. Experimental procedure Commercial purity Cu plates received in 1/2 Hard condition were used in this study as starting materials. For the FSW process, two Cu plates with 2 mm in thickness, 200 mm in length and 50 mm in width were clamped together on the back plate with a 1 mm wide gap between the two adjoining plates. Then the SiC particles with mean size of 5 m were inserted into the gap and pressed tightly. During the welding process, a WC-based alloy tool, which had a 12 mm-diameter shoulder, 4 mm-diameter probe and 2.0 mm probe length, was used with a constant applied load of 1000 kg and titled by 3◦ . A rotation speed of 500 rpm and welding speed of 100 mm/min were used during the welding process. Fig. 1 is the sketch map showing the main principle of FSW process with the addition of SiC powder. To make the SiC particles distribute uniformly in the Cu joints, the welding process were repeated at the same welding condition and regarded as the second pass welding process. After welding, optical microscopy (OM) was used to characterize the macrostructure of the joints. The samples for OM observation were cross-sectioned perpendicular to the welding direction, mechanically polished and then etched with a solution of iron chloride. The electron backscattered diffraction (EBSD) measurements were carried out by using a JEM-7001FA field emission scanning electron microscopy (FE-SEM) with a TSL orientation imaging system. For TEM sample preparation, thin plate were first cut at the target locations and then mechanically polished to a thickness of 100 m. The polished thin plates were finally twin-jet electropolished to make electron beam transparent thin film using a solution of HPO4 :CH4 O:H2 O = 1:1:2 at 5 V and 0 ◦ C. The thin films were observed with a Hitachi 800 TEM at 200 kV. The Vickers hardness profiles of the joint were measured along the centerline of the cross-section by using a Vickers indenter with a load of 0.98 N and a dwell time of 15 s. The tensile specimens were electrical discharge machined into a dog-bone shape with a gauge length of 100 mm, width of 10 mm and thickness of 2 mm, respectively. The tensile
The OM images of the Cu joints with SiC addition after one and two pass processing are shown in Fig. 2. The samples for OM observation were only mechanically polished without chemical etching. For both samples, an obvious area with dispersed particles can be observed and no obvious welding defects or superfluous flash can be found in the joints. After one pass welding process, the area with SiC distribution in the SZ is small and the particles aggregated seriously in the advancing side (AS). In addition, some particles-free regions can be observed and are surrounded by the SiC particles. While after two pass processing, the distribution of SiC particles became very homogenous and appeared in a larger area in the SZ. It was found that for both one pass and two pass processed samples, the SiC particles distribute more in the AS than in the retreating side (RS). According to A.C. Nunes’s physical model of the metal flow in the FSW joint, the material on the RS of the joint is directly transported from the front to the back of the rotation tool, with minimal residence time in the rotation field. This is referred to as the straight-through current flow of the metal. On the contrary, the materials on the AS of the joint resides long enough in the rotational flow around the tool and consists of rotation flow around the pin and downward flow to the top of the pin, which is referred to as the maelstrom current [18]. This model can well explain the asymmetrical distribution of the SiC particles in the Cu weld. The distribution of the SiC can reach the bottom of the AS and a swirl-like pattern formed in the center of the AS due to the maelstrom current of the metal during the FSW process. The gathering of the introduced particles in the AS was also found in the Al joint containing fullerene particles, which was explained to be related with the materials flow caused by the probe and the shoulder [19]. Nevertheless, the particles can be dispersed more homogeneously within an enlarged area by repeating the welding process. Fig. 3(a) and (b) shows the high magnification SEM images of the SZ after one pass and two pass processing, respectively. After one pass, the distribution of SiC particles is not homogeneous, indicating that the materials flow is not enough and the SiC particles cannot be scattered completely into the Cu matrix by the stirring of the rotation tools. The severe aggregation of the SiC particles leads to the pore formation around the particles due to the large difference of physical properties between Cu and SiC. This phenomenon was also found in the friction stir processed Al alloys with SiC particles, in which the distribution of SiC particles is not homogeneous,
Fig. 2. Cross-sectional macrostructure of (a) one pass and (b) two pass FSW processed Cu with SiC particles.
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Fig. 3. SEM micrographs showing the distribution of SiC particles in the SZ of FSW processed Cu after (a) one pass and (b) two pass.
defects at the bottom of the SZ generally forms due to the insufficient stirring during the processing [20]. However, after the second pass, the SiC particles can be uniformly dispersed into the Cu matrix. The SiC particles shows strong bonding with the Cu matrix and no pores or defects formed between Cu and SiC. From the SEM image at higher magnification inserted in the left corner of Fig. 3(b), some large SiC particles with crack inside can be observed, implying that some SiC particles can be broken into small pieces by the strong interactions with the rotation tools. The small SiC particles distributed in the Cu matrix are supposed to be fractured from the large particles. Fig. 4 shows the microstructural transition of the SiC contained Cu joint after two pass processing. Fig. 4(a) shows the microstructure of a large area of the joint, in which several typical zones, namely, BM, heat affected zone (HAZ) and SZ, can be easily distinguished and separated with dashed lines. However, the thermal-mechanical affected zone (TMAZ) cannot be easily distinguished. Fig. 4(b) shows the microstructure of BM, which exhibits a coarse equiaxial grain structure. Fig. 4(c) shows the microstructure
of HAZ, in which no SiC particles can be observed since no plastic deformation took place in this area. The main microstructural difference between the BM and HAZ is that the grain boundary in HAZ is not as obvious as that in the BM. Because the materials in HAZ experienced no plastic deformation but an annealing process during welding, it seems that the grain boundary started to immigrate into each other and have a tendency of grain growth. Fig. 4(d) shows the microstructure of the SZ, in which a banded structure can be observed. The formation of banded microstructure consisting of alternative particle-rich region and particle-free region has been reported by some investigators and it was said that the band spacing was directly related with the welding tool advance per revolution, which is also the main reason for the onion rings formation in the SZ [21]. However, it should be noted that the bands spacing in Fig. 4(d) is much smaller than the tool advance according to the revolution pitch of 0.2 mm/rev in the present study, because two pass FSW process were performed. In addition, it can be found that the grain size in the particle-free region is larger than that in the particle-rich region. However, the grain structure of the Cu matrix
Fig. 4. OM micrographs showing the microstructure transition in the SiC particles dispersed FSW processed Cu after two pass: (a) Large area OM image, (b) BM, (c) HAZ, and (d) SZ.
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Fig. 5. EBSP grain boundary maps showing the microstructure of (a) BM and (b) SZ after two pass processing. (Grey area indicates SiC particles.)
Fig. 6. TEM images showing the microstructure of the SZ after two pass processing: (a) SiC particle at the grain boundary and (b) SiC particle inside Cu grain.
in the particle-rich region cannot be clearly discerned due to the densely distributed SiC particles. In order to understand the effect of particles on the microstructure in the SZ, EBSD measurements were carried out in the SZ and the BM as well. Fig. 5(a) and (b) shows the grain boundary maps of the BM and the SZ of the two pass processed joint, respectively, in which the grain size of Cu matrix in different locations can be easily understood. The BM has a coarse grain structure with average grain size of about 16.2 m. As for SiC dispersed SZ shown in Fig. 5(b), the SiC particles can be easily distinguished due to its poor image quality and was indicated with grey color. It reveals that the banded structure consists of much refined Cu grains in the particle-rich region and relatively coarse Cu grains in the particlefree region. The particles-rich region and particles-free region have an average grain size of about 2 m and 8 m, respectively, which confirms that the particles can refine the microstructure of the FSW processed Cu. Fig. 6 shows the results of the TEM observation in the SZ of the Cu welds after two pass processing. Generally the SiC particles distributes in the Cu matrix through two kinds of ways. One is that the SiC particles distributed at the grain boundaries of Cu grains as shown in Fig. 6(a). Around the SiC particle, high density of dislocations can be observed inside the surrounding Cu grains. Another is that the SiC particle is embedded inside the Cu grain as shown in Fig. 6(b). Dislocations can also be found, however, with very low density. In addition, for all the SiC particles distributed through these two ways, a clear interface is
Fig. 7. Hardness profile along the centerline of the cross-section of the FSW processed Cu with and without SiC additions.
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Fig. 8. Tensile strength and the appearance of the fractured samples of BM and the SiC particle contained FSW Cu joint after one and two pass.
formed between the particles and the Cu matrix. No reaction or transitional layer between SiC particle and Cu matrix can be detected. It was said that during hot deformation of metallic materials containing particles, the particles generally exert three important impacts on the dynamic recrystallization of the matrix materials. One is that the particles can increase stored energy and hence the driving force for recrystallization. Another one is that the deformation heterogeneities at large particles may be sites at which recrystallization originates, which is called particle stimulated nucleation (PSN). The third one is that the particles can also pin the movement of the grain boundary and retard the grain growth after dynamic recrystallization [22]. The PSN during static recrystallization is commonly found in alloys containing hard particles such as carbides and oxides [23–26]. PSN has also been found during dynamic recrystallization in Cu containing Al2 O3 , SiO2 particles and the PSN was greatly depended on the particles size [27]. FSW is a well known severe plastic deformation process. The dynamic recrystallization usually occurs in the stir zone and thus the microstructure can be refined. In the present FSW of pure Cu, the smaller grain size in the particle-rich region is supposed to be caused by either PSN or the pinning effect. The pinning effect corresponds to the situation as shown in Fig. 6(a), the stress on the Cu grains by the particles will certainly cause the generation of high density dislocations. The PSN corresponds to the situation as shown in Fig. 6(b), the SiC particle acts as the nucleation site and the recrystallized Cu grain has a very low dislocation density.
3.2. Mechanical properties Fig. 7 shows the hardness distribution along the center line of the cross-section of the Cu joints after one and two pass processing. The hardness distribution of the Cu joint without SiC particles were also plotted for comparison. For both Cu joints dispersed with SiC particles, the hardness in the SZ is higher than that of the pure Cu counterpart. However, the hardness in the SZ of the one pass processed joint is not uniformly distributed. The low hardness value can be observed in some local areas, which corresponds to the particles-free region in the SZ. While for the joints after two pass processing, the hardness distribution in the SZ is rather homogenous. Although the banded structure consisting of particle-rich and particle-free region formed after two pass processing, the bands spacing is much smaller than the size of the indentation after the hardness tests. As a result, the hardness profile is rather smooth within the entire range of SZ. The hardness increase of the FSW processed Cu with SiC addition is supposed to be contributed by both the decrease of the average grain size and the presence of the hard SiC particles. Although it is hard to say which one of these two factors is more important to the increased hardness, the refinement of the microstructure cannot be achieved if no SiC particles were introduced into the materials during the welding process. In addition, the areas with the lowest hardness value are located in HAZ for all the joints, no matter whether the SiC particles are contained or not. Fig. 8 shows the tensile strength of the Cu joint with and without SiC additions. The one pass processed Cu joint shows a much lower
Fig. 9. Fractography of the tensile specimen after (a) one pass and (b) two pass FSW processing.
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strength of only 125 MPa and the two pass processed Cu joint shows a strength of about 217 MPa, which is very close to the 228 MPa of that without SiC additions. From the appearance of the fractured tensile specimen shown in Fig. 8, the one pass processed joints fractured in the area where the SiC greatly aggregated. However, the two pass processed joints and the joints without SiC particles fractured in the HAZ, corresponding to the lowest hardness value in the SZ. Fig. 9(a) and (b) shows the fractural morphology of the tensile specimens of SiC contained Cu joints after one and two pass processing, respectively. For the specimen after one pass processing, very large pores can be observed on the fracture plane, which is caused by the severe aggregation of SiC particles in the SZ. Although the hardness is generally high in the SZ due to the SiC reinforcement, the formation of pores might tear the continuity of the matrix and becomes the initialization site of the failure during tensile tests. As for the specimen after two pass processing, the fracture took place in the HAZ of the joint. However, the fracture should be close to the SZ, because sparsely distributed SiC particles can also be found on the fractured plane as shown in Fig. 9(b). In addition, large dimples formed due to the relatively large grain size of Cu matrix in HAZ. 4. Conclusions From the above descriptions, the following conclusions can be drawn: a) The micro-sized SiC particles can be introduced into the Cu welds by simply putting the SiC particles between the adjoining plates during the FSW process and the SiC particles can be uniformly distributed in the SZ after two pass processing. b) In the two pass processed joints, banded structure consisting of alternative particle-rich region and particle-free region formed in the SZ. The particle-rich region shows much refined grain structure with average grain size less than 2 m due to the SiC particles stimulated nucleation in the dynamic recrystallization of Cu during the FSW process, while the average grain size in the particle-free region is about 8 m. c) Comparing with the hardness of about 70 HV in the SZ of pure Cu joint without SiC addition, the hardness in the Cu welds can be remarkably increased to 110 HV by the introduction of SiC particles with a slight loss of tensile strength of 11 MPa.
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Acknowledgements The authors wish to acknowledge the financial support of a Grant-in-Aid for Science Research and the Collaborative Research Based on Industrial Demand from the Japan Society for the Promotion of Science and Technology of Japan, the Global COE Program (Project: Center of Excellence for Advanced Structural and Functional Materials Design) from the Ministry of Education, Sports, Culture, Science and Technology of Japan. References [1] R. Nandan, T. Debroy, H.K.D.H. Bhadeshia, Prog. Mater. Sci. 53 (2008) 980–1023. [2] R.S. Mishra, Z.Y. Ma, Mater. Sci. Eng. R 50 (2005) 1–78. [3] P.L. Threadgill, A.J. Leonard, H.R. Shercliff, P.J. Withers, Int. Mater. Rev. 54 (2009) 49–93. [4] S.H.C. Park, Y.S. Sato, H. Kokawa, J. Mater. Sci. 38 (2003) 4379–4383. [5] M. Ghosh, K. Kumar, R.S. Mishra, Scripta Mater. 63 (2010) 851–854. [6] S.M. Mousavizade, F.M. Ghaini, M.J. Torkamany, J. Sabbaghzadeh, A. Abdollahzadeh, Scripta Mater. 60 (2009) 244–247. [7] Y. Morisada, H. Fujii, T. Nagaoka, M. Fukusumi, Mater. Sci. Eng. A 433 (2006) 50–54. [8] Y. Morisada, H. Fujii, T. Nagaoka, M. Fukusumi, Mater. Sci. Eng. A 419 (2006) 344–348. [9] E.R.I. Mahmoud, M. Takahashi, T. Shibayanagi, K. Ikeuchi, Mater. Trans. 50 (2009) 1824–1831. [10] E.R.I. Mahmoud, M. Takahashi, T. Shibayanagi, Wear 268 (2010) 1111– 1121. [11] K. Inada, H. Fujii, Y.S. Ji, Y.F. Sun, Y. Morisada, Sci. Tech. Weld. Join 14 (2009) 41–46. [12] W.B. Lee, S.B. Jung, Mater. Lett. 58 (2004) 1041–1046. [13] Y.F. Sun, H. Fujii, Mater. Sci. Eng. A 527 (2010) 6879–6886. [14] G.M. Xie, Z.Y. Ma, L. Geng, Scripta Mater. 57 (2007) 73–76. [15] L. Zhang, X.H. Qu, X.B. He, B.H. Duan, S.B. Ren, M.L. Qin, Mater. Sci. Eng. A 489 (2008) 285–293. [16] V. Martinez, S. Ordonez, F. Castro, J. Mater. Sci. 38 (2003) 4047–4054. [17] C. Rado, B. Drevet, N. Eustathopoulou, Acta Mater. 48 (2000) 4483–4491. [18] J.A. Schneider, A.C. Nunes, Metall. Mater. Trans. A 35 (2004) 777–783. [19] Y. Morisada, H. Fujii, T. Nagaoka, K. Nogi, M. Fukusumi, Composites: Part A 38 (2007) 2097–2101. [20] E.R.I. Mahmoud, M. Takahashi, T. Shibayanagi, K. Ikeuchi, Sci. Tech. Weld. Join 14 (2009) 413–425. [21] M.A. Sutton, B. Yang, A.P. Reynolds, R. Taylor, Mater. Sci. Eng. A 323 (2002) 160–166. [22] W.L. Zhang, J.X. Wang, F. Yang, Z.Q. Sun, M.Y. Gu, J. Comp. Mater. 40 (2006) 1117–1131. [23] F.J. Humphreys, Met. Sci. 13 (1979) 136–145. [24] F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, 2nd ed., Pergamon Press, Oxford, 2002. [25] F.J. Humphreys, Scripta Mater. 43 (2000) 591–596. [26] R.L. Goetz, Scripta Mater. 52 (2005) 851–856. [27] F.J. Humphreys, M.G. Ardakani, Acta Mater. 44 (1996) 2717–2727.