Research Summary
Lead-Free Solder
Controlling Ag3Sn Plate Formation in Near-Ternary-Eutectic Sn-Ag-Cu Solder by Minor Zn Alloying Sung K. Kang, Da-Yuan Shih, Donovan Leonard, Donald W. Henderson, Timothy Gosselin, Sung-il Cho, Jin Yu, and Won K. Choi
As a result of extensive studies, nearternary-eutectic Sn-Ag-Cu (SAC) alloys have been identified as the leading lead-free solder candidates to replace lead-bearing solders for ball-grid array module assembly. However, recent studies revealed several potential reliability risk factors associated with the alloy system. The formation of large Ag3Sn plates in solder joints, especially when solidified at a relatively slow cooling rate, poses a reliability concern. In this study, the effect of adding a minor amount of zinc in SAC alloy was investigated. The minor zinc addition was shown to reduce the amount of undercooling during solidification and thereby suppress the formation of large Ag3Sn plates. In addition, the zinc was found to cause changes in both the microstructure and interfacial reaction of the solder joint. The interaction of zinc with other alloying elements in the solder was also investigated for a better understanding of the role of zinc during solidification of the nearternary-eutectic alloys.
plate-like Ag3Sn structures can grow rapidly within the liquid phase during cooling, before the final solidification of solder joints, as reported previously.6–8 In addition, when large Ag3Sn plates were present, adverse effects on the plastic-deformation properties of the solder7 were observed, as well as plastic strain localization at the boundary between the Ag3Sn plates and bounding β-Sn phase.7 In a study involving the thermo-mechanical fatigue testing of ceramic BGA solder joints, 8 strain localization at the boundary between the Ag3Sn plates and the β-Sn phase as well as preferred crack growth along the β-Sn/Ag3Sn interface were noted. The growth of large Ag3Sn plates in SAC alloys has been extensively studied in terms of cooling rate,8–10 alloy content,8,9 and minor alloying elements.11,12 The cooling rate during
the reflow of SAC alloys was found to be a critical factor in controlling the formation of large Ag3Sn plates in SAC joints.8–10 At a high cooling rate, such as 1.5°C/s or greater, the formation of large Ag3Sn plates can be kinetically suppressed during a reflow process. However, providing a high cooling rate is not always practical, especially in the case of large-thermal-mass chip carriers. To thermodynamically suppress the formation of large Ag3Sn plates in SAC alloys, the silver content was reduced with a fixed copper content.9,10 It was found that large Ag3Sn plate formation was substantially reduced in alloys with a silver content less than 3 wt.%, even under extremely slow cooling conditions, such as 0.02°C/s.9 In addition, the copper content in SAC alloys was found to have a less significant effect than silver content on the formation of
INTRODUCTION The near-ternary-eutectic Sn-Ag-Cu (SAC) alloys with a melting temperature around 217°C are becoming leading candidates1–3 for lead-free solder technology. Among the potential applications are surface-mount technology card assembly including ball-grid array (BGA) solder joints. In a study of solidification behavior of the neareutectic SAC alloys,4 it was reported that Ag3Sn plate nucleation and ensuing growth may occur with minimal undercooling. In contrast, the β-Sn phase required significantly greater undercooling in order to induce nucleation and bring about final solidification.4,5 As a consequence of this disparity, large, 34
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Figure 1. Optical micrographs of Sn-3.8Ag-0.7Cu solder balls solidified at two different cooling rates: (a) and (b) as-received, fast cooled, (c) and (d) slow cooled at 0.02°C/s after reflow at 250°C.
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large Ag3Sn plates. However, lowering the copper content was also found to be beneficial in reducing the so-called pasty range (the temperature range between liquidus and solidus) of the molten solder, resulting in fewer joint defects.9 Another method reported to control the growth of large Ag3Sn plates is to reduce the amount of undercooling required for the solidification of β–Sn. 11,12 This can be achieved by adding minor alloying elements, such as Zn, Al, Sb, and others to pure Sn13 or near-eutectic SAC alloys. 11,12 This article describes the effects of minor zinc alloying additions in controlling the formation of large Ag3Sn plates. See sidebar for experimental details.
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RESULTS Melting Temperature and Undercooling of Sn-Ag-Cu Alloys
Figure 2. Optical micrographs of Sn-3.8Ag-0.7Cu + 0.1Zn solder balls solidified at two different cooling rates: (a) and (b) as-received, fast cooled; (c) and (d) slow cooled at 0.02°C/s after reflow at 250°C.
The melting temperature of Sn3.8Ag-0.7Cu was measured to be around 217°C, which is in good agreement with the reported value.3,4 The addition of a small amount of zinc alloying element up to 0.7 wt.% did not cause much of a change in melting point as listed in Table A. The amount of undercooling, defined as the temperature difference (∆T = T 1–T2) between the melting point during heating (T1) and the onset temperature of β-Sn solidification during cooling (T2), is estimated for each alloy and listed in Table A. The undercooling required for the β-Sn solidification is around 30°C for Sn-3.8Ag-0.7Cu; for the zinc-added SAC, undercooling is drastically reduced to 4.4°C for the 0.1Zn alloy and 3.6°C for the 0.7Zn. Microstructure of BGA Solder Balls The microstructure of Sn-3.8Ag0.7Cu solder balls was studied in terms of zinc content and cooling rate. A comparison of a typical microstructure of Sn-3.8Ag-0.7Cu solder balls solidified at two different cooling rates is shown in Figure 1. The fast-cooling rate (>> 10°C/s) is associated with the as-received solder balls and the very slow cooling rate (0.02°C/s) is provided by a furnace cool. The solidification microstructure associated with the fast cooling rate 2004 June • JOM
Figure 3. Optical micrographs of Sn-3.8Ag-0.7Cu + 0.7Zn solder balls solidified at two different cooling rates: (a) and (b) as-received, fast cooled; (c) and (d) slow cooled at 0.02°C/s after reflow at 250°C.
exhibits a much finer dendritic structure than the slow-cooled microstructure. A salient feature observed in the slowcooled microstructure is the presence of large Ag3Sn plates as shown in Figure 1c and d. The large Ag3Sn plates are expected to form due to the large amount of undercooling found with SAC alloys in the present study as well as in those reported previously.8–10 Figures 2 and 3 show a typical microstructure of zinc-
added SAC alloys solidified at two different cooling rates: fast cool (>> 10°C/s) vs. slow cool (0.02°C/s). No large Ag3Sn plates were observed in zinc-added SAC alloys, even in the very slow-cooled condition. The population of large Ag3Sn plates in the slow-cooled solder balls was estimated by counting the number of solder balls containing Ag3Sn plates longer than 100 µm on their cross sections of 100 solder balls. 35
EXPERIMENTAL PROCEDURE The Sn-Ag-Cu (SAC) solder alloys investigated in this study are in the form of solder balls, about 890 µm in diameter, commercially produced for ball-grid array (BGA) module assembly. Solder compositions investigated include Sn-3.8Ag-0.7Cu (SAC), Sn-3.8Ag-0.7Cu + 0.1Zn (SAC + 0.1Zn), and Sn-3.8Ag-0.7Cu + 0.7Zn (SAC + 0.7Zn) (all in weight percent with a nominal composition variation of ± 0.2 wt.%). The zinc was added to the SAC alloys using a commercial process for producing BGA solder balls. Table A summarizes the alloy compositions with their melting points and the amount of undercooling measured by using differential scanning calorimetry. During thermal analysis, a heating rate of 1°C/min. and a cooling rate of 6°C/min. were normally employed. The as-received BGA solder balls were solidified at a fast cooling rate, possibly much faster than 10°C/s. To examine the cooling-rate effects on the formation of large Ag3Sn plates, an additional reflow process was applied to the as-received solder balls. After the reflow process at 250°C for 10 min. dwell time, a very slow cooling rate (0.02°C/s) was applied for all BGA solder balls. The microstructures of both as-received and slow-cooled (0.02°C/s) balls were examined to find the population of Ag3Sn plates and any changes in the general microstructure due to the different cooling rates. To reveal the solder microstructure more clearly, the β-Sn matrix was lightly etched with a diluted etchant of 5% HNO3, 3% HCl, and 92% CH3OH for several seconds. Microhardness tests were performed using 10 g load on cross sections of multiple solder balls to detect the mechanical property changes due to cooling rate and alloy composition. Each hardness value reported is an average value of ten indentations or more. An electron-microprobe analysis was performed to investigate the interaction of zinc with other alloying elements in the solder. X-ray elemental mapping for zinc, copper, silver, and tin was obtained in a back-scatter mode. Due to the small amount of alloying elements, a typical collection time of x-ray signals to produce a decent elemental map often took a few hours. For example, 5 hours were needed for zinc. The interfacial reactions between SAC + Zn solders and copper or nickel metallization were investigated by attaching BGA solder balls to a plastic module having copper or Au/Ni pads in a forced convection oven under an N2 atmosphere with the peak temperature ranging from 235°C to 245°C. The cooling rate after the reflow was 0.2°C/s. The cross-sectional microstructure was examined to find the presence of large Ag3Sn plates in the solder joints. Table A. Composition, Melting Temperature, and Undercooling of Sn-3.8Ag-0.7Cu + Zn BGA Solder Balls Solder Composition Sn-3.8Ag-0.7Cu SAC + 0.1 Zn SAC + 0.7 Zn
Melting Temp. during Heating (T1) 217.0 217.7 217.2
The results are shown in Table I. For the regular SAC alloy, 74 out of the 100 balls contained large Ag3Sn plates, while no large Ag3Sn plates were observed in any of 100 balls for both zinc-added alloys. By adding zinc to SAC, the tin dendrite structure in SAC becomes coarse with a decrease in its volume fraction, while the eutectic phase(s) increases in volume fraction. The trends appear to be true for both cooling rates. This suggests that the addition of a small amount of zinc into SAC alloys not only reduces the amount of undercooling, but also modifies the bulk microstructure of SAC alloys. To understand the effects 36
Peak Temp. during Cooling (T2) 187.2 213.3 213.6
Undercooling ∆T (T1–T2) 29.8 4.4 3.6
of zinc addition on the microstructure modification, an electron microprobe analysis was performed to detect a possible segregation of zinc atoms. Figure 4 is an example of elemental mapping of Zn, Cu, Ag, and Sn atoms, respectively. By comparing the series of elemental maps, it was found that zinc atoms were preferentially segregated to copper atoms, possibly forming Cu-Zn intermetallic phases.
increases slightly as the zinc content increases for both the as-received and the slow-cooled solders. The most drastic change in hardness is due to the cooling rate. The fast-cooled solder balls of the fine dendrite microstructure have a much higher microhardness than the slow-cooled solders of large dendrites. The microhardness value of the slow-cooled solder balls is close to that of the SAC annealed at 150°C for 48 h,14 indicating the slow-cooling treatment caused a similar effect on the microhardness of SAC as the annealing treatment previously studied. It is also noted that the zinc-added SAC alloys were less prone to the hardness reduction at the slow cooling rate. Interfacial Reactions in Solder Joints Figure 6 shows the interfacial reactions of SAC + Zn solder balls on copper pads reflowed in a forced convection furnace with a peak temperature of 235–245°C and then solidified at a cooling rate of 0.2°C/s. For SAC + 0.1Zn, large Ag3Sn plates formed at the copper pads as well as inside the solder ball (Figure 6a and b). This is not consistent with the case of the freestanding solder balls, such as those shown in Figure 2c and d. This suggests that 0.1% zinc in SAC alloy was not enough to suppress the formation of large Ag 3Sn plates when joined to copper metallization, because zinc atoms preferentially segregated to the copper pad. However, for SAC + 0.7Zn, the formation of large Ag3Sn plates appears to be well under control, as shown in Figure 6c and d, where only a few small Ag3Sn plates were observed at the copper pad. The microstructure of solder balls attached to copper pads in a module was also influenced by the addition of zinc atoms: tin dendrites become coarser and the volume fraction Table I. Population of Large Ag3Sn Plates in Sn-3.8Ag-0.7Cu + Zn Solder Balls Solidified at a Very Slow Cooling Rate (0.02°C/s)
Microhardness of BGA Solder Balls
Solder Composition (wt.%)
Figure 5 plots microhardness data as a function of zinc content and cooling rate. The microhardness of SAC alloys
Sn-3.8Ag-0.7Cu SAC + 0.1 Zn SAC + 0.7 Zn
# of Solder Balls with Large Ag3Sn Plates 74/100 0/100 0/100
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Figure 4. An electron microprobe analysis of Sn-2.8Ag-0.7Cu + 0.7Zn solder balls, slow cooled at 0.02°C/s after reflow at 250°C.
of the eutectic phase increases as the zinc content increases. DISCUSSION Effects of Zinc Addition to SAC A small addition of zinc into SAC alloys was effective in reducing the amount of undercooling required for β-Sn solidification and thereby controlling the formation of large Ag3Sn plates in SAC solder balls. The thermal analysis data obtained by differential scanning calorimetry with the zinc-added SAC alloys support the microstructural observation of the solder balls. This was true for the wide range of cooling rates, from the very slow (0.02°C/s) to the very fast (>> 10°C/s), applied to the solder balls during solidification. Undercooling for SAC alloys may also be reduced by adding other alloying elements such as aluminum, antimony, and others as reported in the literature.11,13 Another effect of zinc addition was observed in the as-solidified microstructures. As the zinc content increases, tin dendrites become coarser and the volume fraction of the eutectic 2004 June • JOM
phase found between tin dendrites increases at the equivalent cooling rate. As zinc atoms preferentially react with copper atoms, the role of zinc is regarded to influence the nucleation of the eutectic phase(s) and facilitate the formation of the eutectic phase(s) even under non-equilibrium solidification conditions. This trend was observed with both the cooling rates. The microstructure of as-solidified
SAC alloys was significantly changed by the cooling rate applied. The high cooling rate produced a fine dendrite structure with a larger volume fraction of dendrites in comparison with the solder balls solidified at a slow cooling rate. The microhardness measurements have well confirmed the microstructural difference produced by the different cooling rates. The microhardness of the rapidly cooled solder balls is much Figure 5. The microhardness of Sn-3.8Ag-0.7Cu + Zn solder balls: fast-cooled (as-received) vs. slow cooled (0.02°C/s) at 250°C.
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comparison to the size of Ag3Sn plates, such as in flip-chip solder bumps,6,19 large Ag3Sn plates become critical in influencing the mechanical properties of the solder joints and thereby increase their reliability risk factors. Accordingly, it is recommended to control the formation of large Ag3Sn plates in SAC solder joints by either adding a small amount of zinc (or others elements) and/or by reducing silver content.17,18 References
Figure 6. The interfacial reactions of zinc-added Sn-3.8Ag-0.7Cu solder balls on copper pads in a plastic module reflowed in a forced convection oven with the peak temperature of 235–245°C and solidified at 0.2°C/s: (a) and (b) for 0.1Zn, and (c) and (d) for 0.7Zn.
higher than the very slow cooled solder balls. The zinc addition appears to be less significant in causing the hardness changes than the cooling rate effect. Interfacial Reactions of SAC+Zn Solders In the presence of copper metallization, as shown in Figure 6, the addition of 0.1% zinc into SAC was insufficient to control the formation of large Ag3Sn plates, while higher zinc content (0.7%) appears to be adequate to suppress the growth of large Ag3Sn plates. This was attributed to the preferential reaction of zinc atoms with copper atoms as demonstrated by the electron microprobe analysis such as that shown in Figure 4. During reflow at around 240°C, copper atoms from a copper pad can dissolve into the molten tin up to its solubility limit (e.g., 1.2 wt.%).15 Most dissolved copper atoms precipitate out as Cu-Sn intermetallics during solidification to reach the solubility of copper in tin at room temperature (about 0.16 wt.%).4 However, for the zinc-added SAC solders, some of the copper atoms would react with zinc atoms to form Cu-Zn intermetallics instead. In addition, the common Cu-Sn intermetallics formed on a copper pad are expected to be replaced by the Cu-Zn intermetallics,
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as reported in the case of Sn-Zn-Bi alloys.16 The thin intermetallic (IMC) layers on the copper pads in Figure 6 are therefore expected to be Cu-Zn IMCs instead of Cu-Sn IMCs, although this has not yet been confirmed. Based on the interfacial reactions, it is recommended that a zinc content of more than 0.1 wt.%, and possibly around 0.7 wt.%, be used to effectively control the formation of large Ag3Sn plates in SAC alloys. Control of Large Ag3Sn Plates in Solder Joints The role of large Ag3Sn plates during fatigue-crack propagation has been extensively studied through the failure analysis of SAC solder joints with reduced silver contents.17,18 From the interaction of Ag3Sn plates with the crack propagation path it was not clear whether they facilitated or impeded the crack propagation during thermomechanical cycling. Since examples of both were observed in the SAC solder joints that had undergone similar fatigue testing, it is believed that large Ag3Sn plates are detrimental only when the plates are aligned in the direction of maximum shear stress, which facilitates crack propagation. However, when the volume of solder joints becomes much smaller in
1. 2002 Lead-Free Roadmap (Tokyo, Japan: Japan Electronics and Information Technology Industries Association, 2002). 2. J. Bath, C. Handwerker, and E. Bradley: “Research Update: Lead Free Solder Alternatives,” Circuits Assembly, 11, (2000), pp. 45–52. 3. I.E. Anderson et al., J. Electronic Materials, 30 (2001), pp. 1050–1059. 4. K.W. Moon et al., J. of Electronic Materials, 29 (2000), pp. 1122–1136. 5. I. Ohnuma et al., J. of Electronic Materials, 29 (2000), pp. 1137–1144. 6. D.R. Frear et al., JOM, 53 (6) (2001), pp. 28–32. 7. K.S. Kim, S.H. Huh, and K. Suganuma, Materials Science and Engineering, A333 (2002), pp. 106–114. 8. D.W. Henderson et al., J. of Materials Research, 17 (11) (2002), pp. 2775-2778. 9. S.K. Kang et al., Proc. 53rd ECTC (Piscataway, NJ: IEEE, 2003), pp. 64–70. 10. S.K. Kang et al., JOM, 55 (6) (2003), pp. 61–65. 11. K L. Buckmaster et al. (Paper presented at the 2003 TMS Fall Meeting, Chicago, November 2003). 12. S.K. Kang et al. (Paper presented at 2004 TMS Annual Meeting, Charlotte, NC, March 2004). 13. A. Ohno and T. Motegi, J. of Japan Inst Metals, 37 (1973), pp. 777–780. 14. P. Lauro et al., J. Electronic Materials, 32 (12) (2003), pp. 1432–1440. 15. S.K. Kang et al., Proc. 52nd ECTC (Piscataway, NJ: IEEE, 2002), pp. 147–153. 16. P. Harris, Surface Mount Tech. (U.K.), 11 (3) (1999), pp. 46–52. 17. S.K. Kang et al., to be published in Materials Transactions (The Japan Inst. Metals in 2004). 18. S.K. Kang et al., to be published in Proc. Electronic Comp. Tech. Conf. (Piscataway, NJ: IEEE, June 2004). 19. K. Zeng and K.N. Tu, Materials Sci. & Eng., R 38 (2002), pp. 55–105.
Sung K. Kang, Da-Yuan Shih, and Donovan Leonard are with IBM T.J. Watson Research Center in Yorktown Heights, NY. Donald W. Henderson and Timothy Gosselin are with IBM Microelectronics in Endicott, NY. Sung-il Cho and Jin Yu are with the Department of Materials Science and Engineering at KAIST in Daejon, Korea. Won K. Choi is with Samsung Advanced Institute of Technology in Suwon, Korea. For more information, contact Sung K. Kang, IBM T.J. Watson Research Center, 1101 Kitchawan Road, Route 134, P.O. Box 218, Yorktown Heights, NY 10598; (914) 945-3932; fax (914) 945-2141; e-mail
[email protected].
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