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Ultrasonic assisted fabrication of particle reinforced bonds joining aluminum metal matrix composites
Materials and Design 32 (2011) 343–347
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Materials and Design journal homepage: www.elsevier.com/locate/matdes
Short Communication
Ultrasonic assisted fabrication of particle reinforced bonds joining aluminum metal matrix composites Jiuchun Yan, Zhiwu Xu *, Lei Shi, Xing Ma, Shiqin Yang State Key Laboratory of Advanced Welding Production Technology, Harbin Institute of Technology, No. 92, Xidazhi Street, Harbin 150001, PR China
a r t i c l e
i n f o
Article history: Received 25 April 2010 Accepted 21 June 2010 Available online 23 June 2010
a b s t r a c t This paper presents a method of producing uniform particle strengthened bonds between pieces of aluminum metal matrix composite (Al-MMCs), of strength equal to that of the substrate material. SiC particle reinforced Zn-based filler metals were fabricated by mechanical stir casting and ultrasonic treatment, and then used to join pieces of SiCp/A356 composite with the aid of ultrasonic vibration. The filler metals made by mechanical stirring were porous and contained many particle clusters. Ultrasonic vibration was used to disperse the agglomerates and prevent further coagulation of SiC particles during joining, but the method failed to eliminate the porosity, resulting in a highly porous bond. The filler metal treated by ultrasonic vibration was free of defects and produced a non-porous bond strengthened with uniform particles between pieces of SiCp/A356 composite. The presence of surface oxide films at the bonding interface significantly degraded the performance of SiC particle reinforced bond. Removal of this oxide film by at least 4 s of ultrasonic vibration significantly increased the bond strength, reaching a value equal to that of the substrate metal. Crown Copyright Ó 2010 Published by Elsevier Ltd. All rights reserved.
1. Introduction The incorporation of silicon carbide particulates (SiCp) into an aluminum alloy matrix produces materials offering several advantages compared to monolithic Al alloys. These materials are highpotential candidates for structural applications in aerospace, automotive and sports goods where high specific stiffness and strength are particularly required. In recent decades, vigorous research has led to suitable, scalable and affordable processing methods to incorporate SiC particles into an aluminum matrix, most of which involve a molten aluminum phase. The manufacturing of SiCp reinforced Al-MMCs has become well established and these materials are available commercially. However, further application is restricted because joining these materials is an unresolved issue and therefore complex components cannot yet be easily and reliably fabricated [1–4]. One of the most important problems related to joining SiCp reinforced Al-MMC is that conventional aluminum alloy welding processes give joints free of SiCp reinforcement, with mechanical properties far lower than those of the substrate material [2,3]. To improve the mechanical properties of the welded structure, great efforts have been made to incorporate particulate reinforcement into the weld so as to form a composite joint [5–7]. The fusion welding process (e.g. gas tungsten arc welding (GTAW), gas metal arc welding (GMAW) and laser welding) gener* Corresponding author. Tel.: +86 451 86418695; fax: +86 451 86416186. E-mail address: [email protected] (Z. Xu).
ally causes partial melting of the substrate metal and the migration of a certain number of SiC particles into the weld pool. However, SiC particles are thermally unstable in the liquid Al at the processing temperatures, forming an undesirable compound of Al4C3. The weld zone is therefore depleted of the strengthening agent [3]. To compensate for the loss of SiC particles in the weld pool during welding, it is suggested that a filler metal containing SiC particles (i.e. composite filler metal) be used [5]. Unfortunately, the application of the composite filler metal produces a weld pool that has poor fluidity due to the presence of a high fraction of ceramic particulate reinforcement, and this results in high porosity in the solidified structure. Examination of the properties of these kinds of Al-MMC joints shows that the failure is usually located in the weld metal: the weld has a tensile strength less than 50% of that of the substrate material. In transient liquid phase (TLP) bonding of Al-MMC, reinforcing particles are incorporated into the bond region either by using a particle reinforced insert layer [7] or by the melt-back of the substrate metal occurring as a result of the eutectic reaction between the interlayer and the aluminum alloy [8]. A long period of isothermal solidification is needed in this process and particles contained in the liquid formed at the bonding temperature segregate to the joint centerline, resulting in a significant decrease of the joint strength. Based on numerous investigations, Zhou et al. [9,10] demonstrated that only by extremely precise control of the processing conditions can particle segregation be avoided. Moreover, typical limitations of this procedure as regards equipment cost
0261-3069/$ - see front matter Crown Copyright Ó 2010 Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2010.06.036
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J. Yan et al. / Materials and Design 32 (2011) 343–347
and adaptability to component shape and size make it necessary to consider the possibility of the application of more practical and efficient welding processes [11]. Recently, soldering or brazing has been attracting considerable attention for joining Al-MMC, since the application of this technique avoids the problems associated with fusion, and may offer improved flexibility in component design and production compared to solid state methods [11–13]. Wielage [13] developed SiC particle reinforced Sn-based solders and used them in ultrasonic assisted soldering of Al2O3p/6061 MMC. It was found that homogeneous SiC particle reinforced joints were obtained, and the strength of this joint improved approximately 100% compared to the unreinforced joint at elevated testing temperatures. The use of Snbased solder, however, formed a joint possessing strength far lower than that of the substrate material (only about 1/10 of the substrate material strength), inevitably limiting the application of this method. So, the present study aims to develop a SiC particle reinforced Zn-based filler, which has relatively higher strength, and assess the possibility of using ultrasound to fabricate a uniform particle strengthened bond in a SiCp/A356 composite. 2. Experimental The substrate material used was an Al–Si–Cu alloy (A356) reinforced with 20 vol.% SiC particles having an average diameter of 12.6 lm. The nominal composition of the matrix material is given in Table 1. Composite samples were manufactured by stir casting and supplied in the as-cast condition, showing a tensile strength of 240–250 MPa. A Zn–Al alloy was used as the matrix of the SiC particulate reinforced filler material (see Table 1 for detailed composition). The nominal strength of this alloy is lower than that of the substrate material, which is between 180 and 200 MPa. The addition of ceramic particles is expected to improve its properties. SiC particles with the same dimensions and volume fraction as those of the substrate material were used for this purpose, and were incorporated into the Zn–Al alloy by mechanical stirring. A portion of the SiCp/Zn–Al filler was re-melted and treated by ultrasonic treatment in a special experimental setup [14]. Both these filler metals were used to make ultrasonically assisted joint in the SiCP/A356 composite. Prior to the joining experiments, the fusion parameters of the composite filler (solidus and liquidus temperatures) were determined by differential scanning calorimeter (DSC). The ultrasonically assisted joining process of the SiCp/A356 composite samples is shown in Fig. 1. The 45 mm 10 mm 4 mm composite samples were placed in a single 15 mm overlap configuration. The filler materials were cut into 15 mm 10 mm 0.4 mm slices and located between the faying surfaces. The samples were subjected to ultrasonic vibration with an amplitude of 10 lm and frequency of 20 kHz at 420 °C. The ultrasonic vibration time ranged from 0.5 to 6 s, and the sonotrode was removed from the sample after ultrasonic vibration. The microstructures of SiCp/Zn–Al filler metals and joints were examined by optical microscope (Olympus–GX71) and scanning electron microscope (SEM, S-3400N) equipped with energy dispersive spectroscopy (EDS). Shear strength of the joint was evaluated using a specially designed fixture in a tensile testing machine (Instron5569) [11].
Fig. 1. Schematic of the ultrasonic-aided joining process used on SiCp/A356 composite samples.
3. Results and discussion Fig. 2 shows the microstructure of the synthesized SiCp/Zn–Al filler metals. It appears in the mechanically stirred cast filler metal, the SiC particles are mixed well from a macro-scale point of view (see Fig. 2a). However, there are a lot of voids in the filler and the surrounding matrix is characterized by serious particle segregation. When ultrasonic treatment was used, the voids and the local particle segregation in the SiCp/Zn–Al filler metal were eliminated and the distributions of SiC particles are seen to be homogeneous at both the macro and micro scales (Fig. 2b). Mechanical stirring has become an established method of incorporating larger solid ceramic particles into metal melts because it is simple, flexible and applicable to large-quantity production. A mechanical stirrer is used to vigorously stir the liquid metal to form a vortex at the surface of the melt, helping to transfer ceramic particles into the melt and maintain the particles in a state of suspension. However, it is extremely difficult for this method to distribute and disperse fine particles (considered to be <20 lm) uniformly in metal melts due to their large surface-to-volume ratio and low wettability, which induces agglomeration and clustering [15]. Additionally, a vigorously stirred melt is liable to entrap gas and draw it into the melt, resulting in high porosity in the cast product. In comparison, the injection of ultrasonic energy into molten alloys causes some important nonlinear effects such as cavitation, and acoustic streaming. Acoustic streaming is a liquid flow due to an acoustic pressure gradient, leading to highly effective stirring [16,17]. Acoustic cavitation involves the formation, growth, pulsation, and collapsing of micro-bubbles in liquids under cyclic high intensity ultrasonic waves. Some of these cavities grow rapidly under the influence of the alternating pressure and the unidirectional diffusion of dissolved gas products from the melt to the cavities. These large bubbles coagulate and float to the surface of the liquid pool due to the acoustically induced flows in the melt. As a result, the dissolved gases in the liquid melt are substantially removed. However, some of the cavitation bubbles do not have time to fill with gas dissolved in the melt and hence collapse under compression by the ultrasonic wave. In this case, local pressure pulses (up to 1000 MPa), cumulative liquid jets (with a speed being up to 100 m/s) and micro ‘‘hot spots” (with a temperature of about 5000 °C) are formed [15]. Transient cavitations produce an implosive impact strong enough to break up the clustered fine particles and disperse them more uniformly in the liquid phase. This strong
Table 1 Chemical compositions (wt.%) of A356 matrix and Zn–Al filler metal. Alloy
A356 Zn–Al
Element Mg
Cu
Si
Fe
Mn
Ag
Ti
Ni
Al
Zn
0.3–0.45 0.42
<0.2 2.5–3.5
6.5–7.5 0.80
<0.15 –
<0.1 0.20
– 0.40
<0.2 –
– 0.10
Bal. 4.0–5.20
<0.1 Bal.
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(a)
(b)
200 µm
200 µm
Fig. 2. Microstructures of reinforced filler metals: (a) made by stir casting and (b) ultrasonically assisted stir casting.
impact, when coupled with high local temperatures over a brief interval (in the order of 100 ms) can also enhance the wettability between metal melts and particles. These factors make the preparation of micro-particles (e.g. 10–50 lm) reinforced composites possible [18]. Fig. 3 shows the fusion parameters of the SiCp/Zn–Al filler. It can be seen that the melting point of SiCp/Zn–Al filler is in the range of 368–380 °C. Accordingly, joining temperatures below 450 °C can be used and the influence of the thermal cycle on the substrate material will be minimized [3,13]. The microstructure of a SiC particle reinforced bond fabricated between SiCp/A356 composite samples with mechanically-stir-cast SiCp/Zn–Al filler metal is shown in Fig. 4. Since the sizes of the cavitation bubbles are so small and their number is so large, when the cavitation phenomenon is created in a melt, degassing in a small volume of melt is extremely rapid. It only takes a few seconds to degas the filler melt used in each of the present joining cycle [16]. This is why the porous SiCp/Zn–Al filler metals made by mechanical stirring are used in the current study. It can be seen from Fig. 4 that particle clusters involved in the original filler metal disappeared; unfortunately a lot of voids are still detected in the bond produced with this kind of filler metal. This demonstrates that ultrasonic cavitation voids have been formed in the filler melt and while succeeding in dispersing the agglomerates and preventing further coagulation of SiC particles, not all the gas products are removed from the bond. One basic process of ultrasonic degassing is the appearance of coarse coalesced bubbles on the surface of the liquid bath. This process occurs due to the combined action of Stokes forces and acoustic flow [15]. In the present study, the filler metal was
Fig. 3. DSC experimental data for the SiCp/Zn–Al filler metal.
200 µm Fig. 4. Microstructure of a SiC particle reinforced bond fabricated with mechanically-stir-cast SiCp/Zn–Al filler metal.
inserted between the base metals to be joined. This meant that the top surface of the filler metal was covered by the solid substrate metal while it melted, and hence aggregated bubbles could only escape from the side surfaces of the filler melt, which are only about 0.4 lm in width. Thus, the effects of Stokes forces and acoustic flow on the removal of gaseous products from the filler melt were significantly depressed and most of the coarse and fine bubbles were observed to have remained in the solidified bond (Fig. 4). With the use of ultrasonically treated SiCp/Zn–Al filler metal, the microstructure of the SiC particle reinforced bond fabricated with SiCp/A356 composites is shown in Fig. 5. In contrast to Fig. 4, there is no voidage in the bonds shown in Fig. 5. The particle distribution within the bond is reasonably uniform, and is similar to the original filler metal (Fig. 2b). The use of a binary alloy system (i.e. Zn–Al, and Al–Si alloys) as a filler metal has an important feature, which is a significant advantage over Transient Liquid Phase bonding, in that slow isothermal solidification of the liquid zone is unnecessary. Thus, particle pushing by the solidifying front, which could lead to severe particle redistribution, is reduced. Fig. 5 also shows that the SiC particles in the bond retain their original morphology and size, and no signs of partial consumption or interfacial reaction are observed. This result is closely associated with the low temperatures used in the current study avoiding the reaction between molten aluminum and SiC particles to form Al4C3, which occurs at about 725 °C [11]. The formation of a surface oxide film is one of the greatest obstacles to proper metallic bonding between the filler and substrate metals. In the present joining process, the removal of the surface oxide film was accomplished via ultrasonic vibration: ultrasonic induced cavitation in the liquid metal can produce an implosive impact strong enough to damage the surface oxide film. It was found that substantial elimination of the surface oxide film
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(a)
(b) Surface oxide film
200 µm
200 µm
Fig. 5. Microstructure of a SiC particle reinforced bond fabricated with ultrasonically treated SiCp/Zn–Al filler for: (a) 0.5 s and (b) 2 s.
is only achieved when the ultrasonic vibration time exceeds 2 s (Fig. 5). When the ultrasonic vibration time is less than 2 s, continuous oxide films are still present at the bond interface. Further investigation of the bonding interface using back-scattered electron images shows that filler metal had penetrated into the substrate metal before the oxide film was completely destroyed, forming a zinc enriched diffusion layer under the oxide film (see Fig. 6a). The concentration of Zn in the diffusion layer reaches as high as 40–60 wt.%. The zinc enrichment of the substrate metal might lead to a decrease of the local melting temperature and produce partial melting of these penetrated zones. Once partial melting of substrate metal occurs, its surface oxide film becomes suspended. Logically, this oxide film might be more easily destroyed by ultrasonic cavitation, compared with that adhering to the solid substrate metal. In short, the penetration of filler metal into the substrate metal contributes to the removal of surface oxide films by ultrasonic treatment. As the surface oxide film is completely removed, the SiC particle reinforced filler metal is able to form a sound bond with the substrate metal, with aluminum dendrites growing from the substrate metal to the inside of the bond (see Fig. 6b). Fig. 7 shows the variation with ultrasonic vibration time of the strength of the SiC particle reinforced bond. A dramatic increase of the bond strength occurs as the ultrasonic vibration time increases from 0.5 to 2 s. When the ultrasonic vibration time exceeds 2 s, the value of bond strength increases slowly with time reaching that of the substrate metal as the ultrasonic vibration time exceeds 4 s. The microstructure of typical fracture surfaces of SiC particle reinforced bonds are shown in Fig. 8. The fracture surface of the bond appears smooth and contains many signs of being oxidized when ultrasonically treated for 0.5 s (see Fig. 8a), indicating that the fracture occurs at the interface of substrate metal/filler metal and that the trapped surface oxide films are responsible for the low strength of the bond. When the ultrasonic vibration time exceeds 2 s, the fracture surface is characterized by the presence of many shear ductile dimples and SiC reinforcing particles, as
(a)
AlK
ZnK
Fig. 7. Variation of bond shear strength with ultrasonic vibration time.
Fig. 5b shows. The fractures in these cases either occur from the bonding interface of substrate metal/filler metal and advances across the substrate metal subsequently, or occur solely in the substrate metal. The above results clearly indicate that the addition of SiC particles into the filler metal does improve the bond strength. In addition, Energy dispersive spectroscopy (EDS) microanalysis of the fracture surface (Table 2) demonstrates that the aluminum content in the bond matrix increases slightly with ultrasonic vibration time, accompanied by a slight drop of the zinc content. It is likely that the increase of Al/Zn ratio in the bond results from the increase in the amount of dissolved substrate material as the ultrasonic vibration time increases. Probably, this accounts for the slight increase of bond strength when the ultrasonic vibration time is increased from 2 to 6 s, since no obvious change in the distribution of SiC particles in the bond is observed.
(b)
Initial Interface
Diffusion layer
Fig. 6. Bond interfaces of substrate metal/filler metal for different ultrasonic vibration times: (a) 0.5 s and (b) 2 s.
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(b)
(a)
SiC particle
50 µm
50 µm
Fig. 8. Fracture surfaces of SiC particle reinforced bonds for different ultrasonic vibration times: (a) 0.5 s and (b) 2 s.
Team in the Harbin Institute of Technology, and the Development Program for Outstanding Young Teachers in Harbin Institute of Technology (HITQNJS 2008.019).
Table 2 Al, Zn contents (wt.%) in different bond matrixes. Element
Al Zn
Ultrasonic vibration time/s 0.5
1
2
4
6
03.21 89.30
04.35 86.69
06.15 84.50
09.39 81.35
12.83 77.34
4. Conclusions In this paper, SiC particulate reinforced Zn-based filler metals were fabricated and used to join the SiCp/A356 composites with the aid of ultrasonic vibration. The SiC particle reinforced filler metal fabricated by mechanically stirred casting showed serious particle accumulation and voidage due to the poor wetting of particles by the filler metal. During the joining of the SiCp/A356 composite samples with this filler metal, ultrasonic vibration dispersed the agglomerates and prevented further coagulation of SiC particles, but failed to remove the voids from the bond. The SiC particle reinforced filler metal treated by ultrasonic vibration was characterized by an absence of voidage and homogeneous distribution of particles. With this filler metal, non-porous bonds strengthened with uniform particles were produced in the SiCp/A356 composite by ultrasonic vibration. The removal of the surface oxide film was accomplished via 2 s of ultrasonic treatment during joining. Considerable elemental zinc from the filler metal penetrated the oxide film covering the substrate metal, causing partial melting. This was conducive to the removal of surface oxide film by ultrasonic cavitation. The presence of surface oxide films at the bonding interface significantly degraded the performance of the SiC particle reinforced bonds. When the ultrasonic vibration treatment time was greater than 4 s, the SiC particle reinforced bond produced was free of oxide film, and its strength increased to match that of the substrate metal. This investigation provides a method of producing high performance bonds of particle reinforced Al-MMCs at a relatively low temperature. Acknowledgments This research is supported by the National Natural Science Foundation of China (No. 50375039), the Program of Excellent
References [1] Tsunekawa Y, Suzuki H, Genma Y. Application of ultrasonic vibration to in situ MMC process by electromagnetic melt stirring. Mater Des 2001;22:467–72. [2] Ellis MBD. Joining of aluminium based metal matrix composites. Int Mater Rev 1996;42:41–58. [3] Ureña A, Escalera MD, Gil L. Influence of interface reactions on fracture mechanisms in TIG arc-welded aluminium matrix composites. Compos Sci Technol 2000;60:613–22. [4] Arik H, Aydin M, Kurt A, Turker M. Weldability of Al4C3–Al composites via diffusion welding technique. Mater Des 2005;26:555–60. [5] Burke MA, Lessmann GG. Method for laser beam welding metal matrix composite components. US Pat. No. 4803334. [6] Uzun H. Friction stir welding of SiC particulate reinforced AA2124 aluminium alloy matrix composite. Mater Des 2007;28:1440–6. [7] Huang JH, Dong YL, Wan Y, Zhao XK, Zhang H. Investigation on reactive diffusion bonding of SiCp/6063 MMC by using mixed powders as interlayers. J Mater Process Technol 2007;190:312–6. [8] Shirzadi AA, Wallach ER. New approaches for transient liquid phase diffusion bonding of aluminum based metal matrix composites. Mater Sci Technol 1997;13:135–42. [9] Zhou Y, North TH. Counteracting particulate segregation during transient liquid-phase bonding of MMC–MMC and Al2O3–MMC joints. J Mater Sci 1997;32:5571–5. [10] Chen RF, Zuo DW, Wang M. Improvement of joint strength of SiCp/Al metal matrix composite in transient liquid phase bonding using Cu/Ni/Cu film interlayer. J Mater Sci Technol 2006;22:291–4. [11] Ureña A, Gil L, Escriche E, Gómez de Salazar JM, Escalera MD. High temperature soldering of SiC particulate aluminum matrix composites (series 2000) using Zn–Al filler alloys. Sci Technol Weld J 2001;6:1–11. [12] Xu ZW, Yan JC, Wu GH, Yang SQ. Interface structure and strength of ultrasonic vibration liquid phase bonded joints of Al2 O3p/6061Al composites. Scripta Mater 2005;53(7):835–9. [13] Wielage B, Hoyer I, Weis S. Soldering aluminum matrix composites. Weld J 2007;86:67–70. [14] Xu ZW, Yan JC, Chen W, Yang SQ. Effect of ultrasonic vibration on the grain refinement and SiC particle distribution in Zn-based composite filler metal. Mater Lett 2008;62:2615–8. [15] Eskin GI. Cavitation mechanism of ultrasonic melt degassing. Ultrason Sonochem 1995;2:137–41. [16] Xu HB, Jian XG, Meek TT, Han QY. Degassing of molten aluminum A356 alloy using ultrasonic vibration. Mater Lett 2004;58:3669–73. [17] Li JW, Tadashi M, Yoshinori T, Fu Y. Application of ultrasonic treating to degassing of metal ingots. Mater Lett 2008;62:4152–4. [18] Cao G, Konishi H, Li X. Mechanical properties and microstructure of SiCreinforced Mg–(2,4)Al–1Si nanocomposites fabricated by ultrasonic cavitation based solidification processing. Mater Sci Eng A 2008;486:357–62.
Contents lists available at ScienceDirect
Materials and Design journal homepage: www.elsevier.com/locate/matdes
Short Communication
Ultrasonic assisted fabrication of particle reinforced bonds joining aluminum metal matrix composites Jiuchun Yan, Zhiwu Xu *, Lei Shi, Xing Ma, Shiqin Yang State Key Laboratory of Advanced Welding Production Technology, Harbin Institute of Technology, No. 92, Xidazhi Street, Harbin 150001, PR China
a r t i c l e
i n f o
Article history: Received 25 April 2010 Accepted 21 June 2010 Available online 23 June 2010
a b s t r a c t This paper presents a method of producing uniform particle strengthened bonds between pieces of aluminum metal matrix composite (Al-MMCs), of strength equal to that of the substrate material. SiC particle reinforced Zn-based filler metals were fabricated by mechanical stir casting and ultrasonic treatment, and then used to join pieces of SiCp/A356 composite with the aid of ultrasonic vibration. The filler metals made by mechanical stirring were porous and contained many particle clusters. Ultrasonic vibration was used to disperse the agglomerates and prevent further coagulation of SiC particles during joining, but the method failed to eliminate the porosity, resulting in a highly porous bond. The filler metal treated by ultrasonic vibration was free of defects and produced a non-porous bond strengthened with uniform particles between pieces of SiCp/A356 composite. The presence of surface oxide films at the bonding interface significantly degraded the performance of SiC particle reinforced bond. Removal of this oxide film by at least 4 s of ultrasonic vibration significantly increased the bond strength, reaching a value equal to that of the substrate metal. Crown Copyright Ó 2010 Published by Elsevier Ltd. All rights reserved.
1. Introduction The incorporation of silicon carbide particulates (SiCp) into an aluminum alloy matrix produces materials offering several advantages compared to monolithic Al alloys. These materials are highpotential candidates for structural applications in aerospace, automotive and sports goods where high specific stiffness and strength are particularly required. In recent decades, vigorous research has led to suitable, scalable and affordable processing methods to incorporate SiC particles into an aluminum matrix, most of which involve a molten aluminum phase. The manufacturing of SiCp reinforced Al-MMCs has become well established and these materials are available commercially. However, further application is restricted because joining these materials is an unresolved issue and therefore complex components cannot yet be easily and reliably fabricated [1–4]. One of the most important problems related to joining SiCp reinforced Al-MMC is that conventional aluminum alloy welding processes give joints free of SiCp reinforcement, with mechanical properties far lower than those of the substrate material [2,3]. To improve the mechanical properties of the welded structure, great efforts have been made to incorporate particulate reinforcement into the weld so as to form a composite joint [5–7]. The fusion welding process (e.g. gas tungsten arc welding (GTAW), gas metal arc welding (GMAW) and laser welding) gener* Corresponding author. Tel.: +86 451 86418695; fax: +86 451 86416186. E-mail address: [email protected] (Z. Xu).
ally causes partial melting of the substrate metal and the migration of a certain number of SiC particles into the weld pool. However, SiC particles are thermally unstable in the liquid Al at the processing temperatures, forming an undesirable compound of Al4C3. The weld zone is therefore depleted of the strengthening agent [3]. To compensate for the loss of SiC particles in the weld pool during welding, it is suggested that a filler metal containing SiC particles (i.e. composite filler metal) be used [5]. Unfortunately, the application of the composite filler metal produces a weld pool that has poor fluidity due to the presence of a high fraction of ceramic particulate reinforcement, and this results in high porosity in the solidified structure. Examination of the properties of these kinds of Al-MMC joints shows that the failure is usually located in the weld metal: the weld has a tensile strength less than 50% of that of the substrate material. In transient liquid phase (TLP) bonding of Al-MMC, reinforcing particles are incorporated into the bond region either by using a particle reinforced insert layer [7] or by the melt-back of the substrate metal occurring as a result of the eutectic reaction between the interlayer and the aluminum alloy [8]. A long period of isothermal solidification is needed in this process and particles contained in the liquid formed at the bonding temperature segregate to the joint centerline, resulting in a significant decrease of the joint strength. Based on numerous investigations, Zhou et al. [9,10] demonstrated that only by extremely precise control of the processing conditions can particle segregation be avoided. Moreover, typical limitations of this procedure as regards equipment cost
0261-3069/$ - see front matter Crown Copyright Ó 2010 Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2010.06.036
344
J. Yan et al. / Materials and Design 32 (2011) 343–347
and adaptability to component shape and size make it necessary to consider the possibility of the application of more practical and efficient welding processes [11]. Recently, soldering or brazing has been attracting considerable attention for joining Al-MMC, since the application of this technique avoids the problems associated with fusion, and may offer improved flexibility in component design and production compared to solid state methods [11–13]. Wielage [13] developed SiC particle reinforced Sn-based solders and used them in ultrasonic assisted soldering of Al2O3p/6061 MMC. It was found that homogeneous SiC particle reinforced joints were obtained, and the strength of this joint improved approximately 100% compared to the unreinforced joint at elevated testing temperatures. The use of Snbased solder, however, formed a joint possessing strength far lower than that of the substrate material (only about 1/10 of the substrate material strength), inevitably limiting the application of this method. So, the present study aims to develop a SiC particle reinforced Zn-based filler, which has relatively higher strength, and assess the possibility of using ultrasound to fabricate a uniform particle strengthened bond in a SiCp/A356 composite. 2. Experimental The substrate material used was an Al–Si–Cu alloy (A356) reinforced with 20 vol.% SiC particles having an average diameter of 12.6 lm. The nominal composition of the matrix material is given in Table 1. Composite samples were manufactured by stir casting and supplied in the as-cast condition, showing a tensile strength of 240–250 MPa. A Zn–Al alloy was used as the matrix of the SiC particulate reinforced filler material (see Table 1 for detailed composition). The nominal strength of this alloy is lower than that of the substrate material, which is between 180 and 200 MPa. The addition of ceramic particles is expected to improve its properties. SiC particles with the same dimensions and volume fraction as those of the substrate material were used for this purpose, and were incorporated into the Zn–Al alloy by mechanical stirring. A portion of the SiCp/Zn–Al filler was re-melted and treated by ultrasonic treatment in a special experimental setup [14]. Both these filler metals were used to make ultrasonically assisted joint in the SiCP/A356 composite. Prior to the joining experiments, the fusion parameters of the composite filler (solidus and liquidus temperatures) were determined by differential scanning calorimeter (DSC). The ultrasonically assisted joining process of the SiCp/A356 composite samples is shown in Fig. 1. The 45 mm 10 mm 4 mm composite samples were placed in a single 15 mm overlap configuration. The filler materials were cut into 15 mm 10 mm 0.4 mm slices and located between the faying surfaces. The samples were subjected to ultrasonic vibration with an amplitude of 10 lm and frequency of 20 kHz at 420 °C. The ultrasonic vibration time ranged from 0.5 to 6 s, and the sonotrode was removed from the sample after ultrasonic vibration. The microstructures of SiCp/Zn–Al filler metals and joints were examined by optical microscope (Olympus–GX71) and scanning electron microscope (SEM, S-3400N) equipped with energy dispersive spectroscopy (EDS). Shear strength of the joint was evaluated using a specially designed fixture in a tensile testing machine (Instron5569) [11].
Fig. 1. Schematic of the ultrasonic-aided joining process used on SiCp/A356 composite samples.
3. Results and discussion Fig. 2 shows the microstructure of the synthesized SiCp/Zn–Al filler metals. It appears in the mechanically stirred cast filler metal, the SiC particles are mixed well from a macro-scale point of view (see Fig. 2a). However, there are a lot of voids in the filler and the surrounding matrix is characterized by serious particle segregation. When ultrasonic treatment was used, the voids and the local particle segregation in the SiCp/Zn–Al filler metal were eliminated and the distributions of SiC particles are seen to be homogeneous at both the macro and micro scales (Fig. 2b). Mechanical stirring has become an established method of incorporating larger solid ceramic particles into metal melts because it is simple, flexible and applicable to large-quantity production. A mechanical stirrer is used to vigorously stir the liquid metal to form a vortex at the surface of the melt, helping to transfer ceramic particles into the melt and maintain the particles in a state of suspension. However, it is extremely difficult for this method to distribute and disperse fine particles (considered to be <20 lm) uniformly in metal melts due to their large surface-to-volume ratio and low wettability, which induces agglomeration and clustering [15]. Additionally, a vigorously stirred melt is liable to entrap gas and draw it into the melt, resulting in high porosity in the cast product. In comparison, the injection of ultrasonic energy into molten alloys causes some important nonlinear effects such as cavitation, and acoustic streaming. Acoustic streaming is a liquid flow due to an acoustic pressure gradient, leading to highly effective stirring [16,17]. Acoustic cavitation involves the formation, growth, pulsation, and collapsing of micro-bubbles in liquids under cyclic high intensity ultrasonic waves. Some of these cavities grow rapidly under the influence of the alternating pressure and the unidirectional diffusion of dissolved gas products from the melt to the cavities. These large bubbles coagulate and float to the surface of the liquid pool due to the acoustically induced flows in the melt. As a result, the dissolved gases in the liquid melt are substantially removed. However, some of the cavitation bubbles do not have time to fill with gas dissolved in the melt and hence collapse under compression by the ultrasonic wave. In this case, local pressure pulses (up to 1000 MPa), cumulative liquid jets (with a speed being up to 100 m/s) and micro ‘‘hot spots” (with a temperature of about 5000 °C) are formed [15]. Transient cavitations produce an implosive impact strong enough to break up the clustered fine particles and disperse them more uniformly in the liquid phase. This strong
Table 1 Chemical compositions (wt.%) of A356 matrix and Zn–Al filler metal. Alloy
A356 Zn–Al
Element Mg
Cu
Si
Fe
Mn
Ag
Ti
Ni
Al
Zn
0.3–0.45 0.42
<0.2 2.5–3.5
6.5–7.5 0.80
<0.15 –
<0.1 0.20
– 0.40
<0.2 –
– 0.10
Bal. 4.0–5.20
<0.1 Bal.
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(a)
(b)
200 µm
200 µm
Fig. 2. Microstructures of reinforced filler metals: (a) made by stir casting and (b) ultrasonically assisted stir casting.
impact, when coupled with high local temperatures over a brief interval (in the order of 100 ms) can also enhance the wettability between metal melts and particles. These factors make the preparation of micro-particles (e.g. 10–50 lm) reinforced composites possible [18]. Fig. 3 shows the fusion parameters of the SiCp/Zn–Al filler. It can be seen that the melting point of SiCp/Zn–Al filler is in the range of 368–380 °C. Accordingly, joining temperatures below 450 °C can be used and the influence of the thermal cycle on the substrate material will be minimized [3,13]. The microstructure of a SiC particle reinforced bond fabricated between SiCp/A356 composite samples with mechanically-stir-cast SiCp/Zn–Al filler metal is shown in Fig. 4. Since the sizes of the cavitation bubbles are so small and their number is so large, when the cavitation phenomenon is created in a melt, degassing in a small volume of melt is extremely rapid. It only takes a few seconds to degas the filler melt used in each of the present joining cycle [16]. This is why the porous SiCp/Zn–Al filler metals made by mechanical stirring are used in the current study. It can be seen from Fig. 4 that particle clusters involved in the original filler metal disappeared; unfortunately a lot of voids are still detected in the bond produced with this kind of filler metal. This demonstrates that ultrasonic cavitation voids have been formed in the filler melt and while succeeding in dispersing the agglomerates and preventing further coagulation of SiC particles, not all the gas products are removed from the bond. One basic process of ultrasonic degassing is the appearance of coarse coalesced bubbles on the surface of the liquid bath. This process occurs due to the combined action of Stokes forces and acoustic flow [15]. In the present study, the filler metal was
Fig. 3. DSC experimental data for the SiCp/Zn–Al filler metal.
200 µm Fig. 4. Microstructure of a SiC particle reinforced bond fabricated with mechanically-stir-cast SiCp/Zn–Al filler metal.
inserted between the base metals to be joined. This meant that the top surface of the filler metal was covered by the solid substrate metal while it melted, and hence aggregated bubbles could only escape from the side surfaces of the filler melt, which are only about 0.4 lm in width. Thus, the effects of Stokes forces and acoustic flow on the removal of gaseous products from the filler melt were significantly depressed and most of the coarse and fine bubbles were observed to have remained in the solidified bond (Fig. 4). With the use of ultrasonically treated SiCp/Zn–Al filler metal, the microstructure of the SiC particle reinforced bond fabricated with SiCp/A356 composites is shown in Fig. 5. In contrast to Fig. 4, there is no voidage in the bonds shown in Fig. 5. The particle distribution within the bond is reasonably uniform, and is similar to the original filler metal (Fig. 2b). The use of a binary alloy system (i.e. Zn–Al, and Al–Si alloys) as a filler metal has an important feature, which is a significant advantage over Transient Liquid Phase bonding, in that slow isothermal solidification of the liquid zone is unnecessary. Thus, particle pushing by the solidifying front, which could lead to severe particle redistribution, is reduced. Fig. 5 also shows that the SiC particles in the bond retain their original morphology and size, and no signs of partial consumption or interfacial reaction are observed. This result is closely associated with the low temperatures used in the current study avoiding the reaction between molten aluminum and SiC particles to form Al4C3, which occurs at about 725 °C [11]. The formation of a surface oxide film is one of the greatest obstacles to proper metallic bonding between the filler and substrate metals. In the present joining process, the removal of the surface oxide film was accomplished via ultrasonic vibration: ultrasonic induced cavitation in the liquid metal can produce an implosive impact strong enough to damage the surface oxide film. It was found that substantial elimination of the surface oxide film
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(a)
(b) Surface oxide film
200 µm
200 µm
Fig. 5. Microstructure of a SiC particle reinforced bond fabricated with ultrasonically treated SiCp/Zn–Al filler for: (a) 0.5 s and (b) 2 s.
is only achieved when the ultrasonic vibration time exceeds 2 s (Fig. 5). When the ultrasonic vibration time is less than 2 s, continuous oxide films are still present at the bond interface. Further investigation of the bonding interface using back-scattered electron images shows that filler metal had penetrated into the substrate metal before the oxide film was completely destroyed, forming a zinc enriched diffusion layer under the oxide film (see Fig. 6a). The concentration of Zn in the diffusion layer reaches as high as 40–60 wt.%. The zinc enrichment of the substrate metal might lead to a decrease of the local melting temperature and produce partial melting of these penetrated zones. Once partial melting of substrate metal occurs, its surface oxide film becomes suspended. Logically, this oxide film might be more easily destroyed by ultrasonic cavitation, compared with that adhering to the solid substrate metal. In short, the penetration of filler metal into the substrate metal contributes to the removal of surface oxide films by ultrasonic treatment. As the surface oxide film is completely removed, the SiC particle reinforced filler metal is able to form a sound bond with the substrate metal, with aluminum dendrites growing from the substrate metal to the inside of the bond (see Fig. 6b). Fig. 7 shows the variation with ultrasonic vibration time of the strength of the SiC particle reinforced bond. A dramatic increase of the bond strength occurs as the ultrasonic vibration time increases from 0.5 to 2 s. When the ultrasonic vibration time exceeds 2 s, the value of bond strength increases slowly with time reaching that of the substrate metal as the ultrasonic vibration time exceeds 4 s. The microstructure of typical fracture surfaces of SiC particle reinforced bonds are shown in Fig. 8. The fracture surface of the bond appears smooth and contains many signs of being oxidized when ultrasonically treated for 0.5 s (see Fig. 8a), indicating that the fracture occurs at the interface of substrate metal/filler metal and that the trapped surface oxide films are responsible for the low strength of the bond. When the ultrasonic vibration time exceeds 2 s, the fracture surface is characterized by the presence of many shear ductile dimples and SiC reinforcing particles, as
(a)
AlK
ZnK
Fig. 7. Variation of bond shear strength with ultrasonic vibration time.
Fig. 5b shows. The fractures in these cases either occur from the bonding interface of substrate metal/filler metal and advances across the substrate metal subsequently, or occur solely in the substrate metal. The above results clearly indicate that the addition of SiC particles into the filler metal does improve the bond strength. In addition, Energy dispersive spectroscopy (EDS) microanalysis of the fracture surface (Table 2) demonstrates that the aluminum content in the bond matrix increases slightly with ultrasonic vibration time, accompanied by a slight drop of the zinc content. It is likely that the increase of Al/Zn ratio in the bond results from the increase in the amount of dissolved substrate material as the ultrasonic vibration time increases. Probably, this accounts for the slight increase of bond strength when the ultrasonic vibration time is increased from 2 to 6 s, since no obvious change in the distribution of SiC particles in the bond is observed.
(b)
Initial Interface
Diffusion layer
Fig. 6. Bond interfaces of substrate metal/filler metal for different ultrasonic vibration times: (a) 0.5 s and (b) 2 s.
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(b)
(a)
SiC particle
50 µm
50 µm
Fig. 8. Fracture surfaces of SiC particle reinforced bonds for different ultrasonic vibration times: (a) 0.5 s and (b) 2 s.
Team in the Harbin Institute of Technology, and the Development Program for Outstanding Young Teachers in Harbin Institute of Technology (HITQNJS 2008.019).
Table 2 Al, Zn contents (wt.%) in different bond matrixes. Element
Al Zn
Ultrasonic vibration time/s 0.5
1
2
4
6
03.21 89.30
04.35 86.69
06.15 84.50
09.39 81.35
12.83 77.34
4. Conclusions In this paper, SiC particulate reinforced Zn-based filler metals were fabricated and used to join the SiCp/A356 composites with the aid of ultrasonic vibration. The SiC particle reinforced filler metal fabricated by mechanically stirred casting showed serious particle accumulation and voidage due to the poor wetting of particles by the filler metal. During the joining of the SiCp/A356 composite samples with this filler metal, ultrasonic vibration dispersed the agglomerates and prevented further coagulation of SiC particles, but failed to remove the voids from the bond. The SiC particle reinforced filler metal treated by ultrasonic vibration was characterized by an absence of voidage and homogeneous distribution of particles. With this filler metal, non-porous bonds strengthened with uniform particles were produced in the SiCp/A356 composite by ultrasonic vibration. The removal of the surface oxide film was accomplished via 2 s of ultrasonic treatment during joining. Considerable elemental zinc from the filler metal penetrated the oxide film covering the substrate metal, causing partial melting. This was conducive to the removal of surface oxide film by ultrasonic cavitation. The presence of surface oxide films at the bonding interface significantly degraded the performance of the SiC particle reinforced bonds. When the ultrasonic vibration treatment time was greater than 4 s, the SiC particle reinforced bond produced was free of oxide film, and its strength increased to match that of the substrate metal. This investigation provides a method of producing high performance bonds of particle reinforced Al-MMCs at a relatively low temperature. Acknowledgments This research is supported by the National Natural Science Foundation of China (No. 50375039), the Program of Excellent
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