Microstructure and Mechanical Properties of Sn-8.55Zn-1Ag-XAl Solder Alloys

Transcription

1 Materials Transactions, Vol. 46, No. 1 (2005) pp. 42 to 47 #2005 The Japan Institute of Metals Microstructure and Mechanical Properties of Sn-8.55Zn-1Ag-XAl Solder Alloys Shou-Chang Cheng 1; * and Kwang-Lung Lin 2 1 Department of Electronics Engineering and Computer Science, Tung-Fang Institute of Technology, Kaohsiung, Taiwan 829, R.O. China 2 Department of Materials Science and Engineering, National Cheng-Kung University, Tainan, Taiwan 701, R.O. China The microstructure and mechanical properties of as cast lead-free solders including eutectic Sn-9Zn, Sn-8.55Zn-1Ag and Sn-8.55Zn-1Ag- XAl (X ¼ 0:010:45 mass%) alloys were investigated. Microstructures of the Sn-Zn-Ag-XAl alloys consist of compound, phase, Al-rich segregation and hypoeutectic phase. The addition of Al dramatically improves the 0.2% offset yield stress, Vickers hardness, ultimate tensile stress (UTS) and total tensile strain of the Sn-8.55Zn-1Ag-XAl alloys. The increase in the Al content of the Sn-8.55Zn-1Ag-XAl alloy from 0.01 to 0.45 mass% increases the average yield stress from 49.9 to 54.0 MPa (47:6 MPa for eutectic Sn-9Zn) and the Vickers hardness from 17.0 to 18.3 Hv (16:8 Hv for eutectic Sn-9Zn). The average values of UTS of the Sn-8.55Zn-1Ag-XAl alloys with X ¼ 0:01, 0.1, 0.25 and 0.45 mass% Al are 55, 58, 55 and 60 MPa, respectively, while those of the tensile strain are 47%, 52%, 58% and 45%, respectively (52 MPa and 50% for the eutectic Sn-9Zn alloy). Fracture mechanisms of the Sn-8.55Zn-1Ag-XAl alloys are correlated with the phase and Al segregation. (Received May 14, 2004; Accepted November 4, 2004) Keywords: fracture, hardness, lead-free solder, mechanical property, microstructure, tension test 1. Introduction In terms of mechanical integrity, Sn-Ag system alloys have been recognized as one of the best choices as the lead-free solder. 1 3) However, the Sn-Ag alloys have higher melting temperatures of K than that of 456 K for the eutectic Sn-Pb alloy. 4,5) In contrast, the eutectic Sn-9Zn solder alloy has a melting temperature of 471 K, which is close to that of the eutectic Sn-Pb alloy. The mechanical strength of the eutectic Sn-Zn alloy is comparable to that of the eutectic Sn- Pb alloy. 6) The eutectic Sn-Zn alloy has been investigated for its fatigue, 5) creep, 7) soldering behavior, 8) shear strength, 9) wettability 5,10) and mechanical alloying. 11) The microstructure and soldering ability of Sn-Zn system alloys with Cu have also been examined ) It has been reported that Sn- Zn-Al solder alloys have good wettability on a Cu substrate ) The microstructure of a Sn-9Zn-0.45Al lead-free solder alloy has been investigated using scanning electron microscopy. 19) The corrosion behavior of Sn-Zn-Al 20,21) and Sn-Zn-Al-In 22) alloys in a 3.5% NaCl aqua has been investigated by potentiodynamic polarization. Small alloying addition of Ag dramatically improved the mechanical properties of the ternary Sn-8Zn-5In alloy. 22) The wettability of Sn-8.55Zn-1Ag-XAl solder alloys was improved by the increasing addition of Al. 23) Aluminum has a melting temperature of K and good electrical conductivity. According to the Al-Sn, 24) Al-Zn 25) and Al-Ag 26) binary phase diagrams, Al may form solid solutions with Sn, Zn and Ag. It is desired to have a lead-free solder with a melting temperature close to the eutectic temperature of the Sn-Pb alloy. A soldering temperature higher than that for the eutectic Sn-Pb alloy can cause undesirable thermal damage and/or warpage to printedcircuit boards, semiconductor chips and electronic packaging materials. 27) The present work tried to investigate the Fig. 1 mechanical properties and microstructures of Sn-8.55Zn- 1Ag-XAl solder alloys with various Al contents. 2. Experimental Procedures Schematic diagram of tensile specimen. Sn-9Zn, Sn-8.55Zn-1Ag and Sn-8.55Zn-1Ag-XAl solder alloys were prepared from pure elements with 99.9% purity. The Al content of the Sn-8.55Zn-1Ag-XAl alloys investigated ranges from 0.01 to 0.45 mass%. The molten alloys were homogenized at 773 K for 1 hour, and then casted in a stainless steel mold at a cooling rate of 95 K/min. The ascast alloys were machined with a wire cutting machine into tensile specimens as shown in Fig. 1. Tensile tests were performed at a crosshead speed of 1 mm/min according to ASTM-A370 standard. The Vickers hardness of the alloy was measured with a loading of 98 N for 15 sec. The specimens were prepared for hardness testing with ASTM-E3 standard. Phase identification of the alloy was carried out by X-ray diffractometry (XRD) at 30 kev using Cu-K radiation with diffraction angles (2) from 35 to 50 and a constant scanning speed of 1 /min. The microstructure of the alloy was investigated with a scanning electron microscope (SEM). The composition of each phase in the alloy was analyzed with an energy dispersive spectrometer (EDS). *Corresponding author, matkllin@mail.ncku.edu.tw

2 Microstructure and Mechanical Properties of Sn-8.55Zn-1Ag-XAl Solder Alloys 43,, (a) Sn-Zn eutectic (d), P-Sn, A B (b) (e),, (c) (f) Fig. 2 SEM micrographs of (a) Sn-9Zn, (b) Sn-8.55Zn-1Ag, (c) Sn-8.55Zn-1Ag-0.01Al, (d) Sn-8.55Zn-1Ag-0.1Al, (e) Sn-8.55Zn-1Ag- 0.25Al and (f) Sn-8.55Zn-1Ag-0.45Al alloys. 3. Results 3.1 Microstructure Figure 2 shows the microstructures of the Sn-9Zn, Sn-8.55Zn-1Ag and Sn-8.55Zn-1Ag-XAl (X ¼ 0:010:45 mass%) alloys. The microstructure of the Sn-9Zn alloy in Fig. 2(a) shows a typical eutectic structure with the phase distributing uniformly in the matrix. The addition of 1 mass% Ag to the Sn-Zn alloy results in a few needle-like precipitates of the phase, dendritic blocks of the compound (" phase) and hypoeutectic structure in the matrix as indicated in Fig. 2(b). The eutectic structure (points A and B) and the -Sn phase (P-Sn) are also formed upon the Ag addition. The increasing addition of Al from 0.01 to 0.25 mass%, to the Sn-Zn-Ag alloy tends to enlarge the needle-like precipitate of the phase and the dendritic block of the compound as seen in Figs. 2(c) (e). At an Al content of 0.45 mass%, Fig. 2(f), the microstructure of the alloy consists of near diamond shape Al-rich phase (Al seg.), rod-like phase and rounded compound. The phase stays together with the compound [Figs. 2(d) and (f)]. These phases became more abundant in the matrix in Fig. 2(f), but the phase shrinks in the matrix of the Sn-8.55Zn-1Ag-0.45Al alloy. When the Al content increases up to 0:10:45 mass%, the eutectic boundaries became prominent in the matrix [Figs. 2(d)(f)].

3 44 S.-C. Cheng and K.-L. Lin Fig. 4 The 0.2% offset yield stresses of Sn-9Zn, Sn-8.55Zn-1Ag and Sn- 8.55Zn-1Ag-XAl alloys. Fig. 3 XRD patterns of Sn-9Zn, Sn-8.55Zn-1Ag and Sn-8.55Zn-1Ag-XAl alloys. 3.2 XRD analysis Figure 3 shows typical results of XRD analysis of the Sn- Zn-Ag-Al alloys. The Bragg peaks corresponding to Sn(211), Sn(220), Sn(101), Zn(101) and Zn(100) are observed in the eutectic Sn-9Zn and Sn-8.55Zn-1Ag alloys. The latter, in addition, exhibits the Bragg peaks for (002) and (101). The XRD analysis of the Sn-9Zn alloy shows the phases of the eutectic structure. The diffraction patterns of the Al (0.01 to 0.45 mass%) containing alloys are almost the same as that of the Sn-8.55Zn-1Ag alloy. 3.3 The 0.2% offset yield stress Figure 4 shows the 0.2% offset yield stress of the alloys investigated. The average values of the yield stress of all the alloys are between 41.2 and 54.0 MPa. Those of the Sn- 8.55Zn-1Ag-XAl alloys with X ¼ 0:01, 0.1, 0.25 and 0.45 mass% Al are 52, 54, 50 and 54 MPa, respectively. The average value of the yield stress of the eutectic Sn-9Zn alloy is 47.6 MPa. The Sn-8.55Zn-1Ag-0.45Al alloy exhibits the highest average yield stress of 54.0 MPa, while the Sn- Zn-1Ag alloy shows the lowest one of 41.2 MPa. It is noticeable that the Al containing Sn-8.55Zn-1Ag-XAl alloys possess greater average values of the yield stress than that of the Sn-9Zn alloy. The addition of Ag lowers the yield stress of the Sn-9Zn alloy. The yield stress of the Sn-8.55Zn-1Ag alloy is improved by the addition of Al. 3.4 Ultimate tensile stress and total tensile strain The average values of the UTS (ultimate tensile stress) of all the alloys are within 45 and 60 MPa, while those of the total strain are within 45 and 58% as presented in Fig. 4. The average values of the UTS and total strain of the Sn-9Zn alloy are 52 MPa and 50%, respectively. The Sn-8.55Zn-1Ag alloy has the lowest average value of the UTS of 45 MPa and the average tensile strain of 50%. The average values of the UTS of Sn-8.55Zn-1Ag-XAl alloys with X ¼ 0:01, 0.1, 0.25 and 0.45 mass% Al are 55, 58, 55 and 60 MPa, respectively, while those of the average tensile strain are 47, 52, 58 and 45%, respectively. 3.5 Vickers hardness The average values of the hardness, Fig. 4, of the Sn-9Zn, Sn-8.55Zn-1Ag and Sn-8.55Zn-1Ag-XAl alloys with 0.01, 0.1, 0.25 and 0.45 mass% Al are 16.8, 13.7, 18.3, 17.0, 16.9 and 17.2 HV, respectively. The Sn-8.55Zn-1Ag alloy has the lowest average value of the hardness of 13.7 HV. The Al containing Sn-8.55Zn-1Ag-XAl alloys possess greater hardness values than that of the Sn-9Zn alloy. The Ag addition lowers the hardness of the Sn-9Zn alloy. The result also shows that the hardness of the Sn-8.55Zn-1Ag alloy is improved by the addition of Al. According to the results in Sections 3.2 to 3.5, it is concluded that the average values of the UTS, yield stress and hardness for all the Al-containing alloys are much greater than those of the Sn-9Zn and Sn-Zn- 1Ag alloys. The tensile strain was also greater for the 0.1Aland 0.25Al-containing alloys than for the Sn-9Zn and Sn-Zn- 1Ag alloys. 3.6 Fracture morphology Figures 5(a) to (f) show morphology of the fracture surfaces for all the alloys investigated. The Sn-9Zn alloy shows a typical dimple fracture surface as indicated in Fig. 5(a). The morphology of the fracture surfaces of all the other alloys also show dimples. These dimple structures along with a magnified fractograph in Fig. 6 of the Sn- 8.55Zn-1Ag-0.01Al alloy indicate the ductile fracture behavior. Figure 7 shows the cross-sectional microstructures of fracture of various alloys. For the Sn-9Zn alloy, after tensile

4 Microstructure and Mechanical Properties of Sn-8.55Zn-1Ag-XAl Solder Alloys 45 fracture surface fracture surface voids voids crack voids (a) (b) (a) (b) crack microvoids microvoids fracture surface (c) (d) (c) (d) (e) (f) (e) Fig. 5 Fractographs of (a) Sn-9Zn, (b) Sn-8.55Zn-1Ag, (c) Sn-8.55Zn- 1Ag-0.01Al, (d) Sn-8.55Zn-1Ag-0.1Al, (e) Sn-8.55Zn-1Ag-0.25Al and (f) Sn-8.55Zn-1Ag-0.45Al alloys. Fig. 7 Cross sectional SEM micrographs of fracture for (a) Sn-9Zn, (b) Sn- 8.55Zn-1Ag, (c) Sn-8.55Zn-1Ag-0.1Al, (d) and (e) Sn-8.55Zn-1Ag- 0.45Al alloys. ~ through the eutectic and the phase as shown in Figs. 7(d) and (e). 4. Discussion Fig. 6 Morphology inside dimple of Sn-8.55Zn-1Ag-0.01Al alloy. process, cracks propagated through voids, and then the fracture surface was formed, as observed in Fig. 7(a). The crack also propagated through the void in the matrix of the Sn-8.55Zn-1Ag alloy after tensile process as shown in Fig. 7(b). In the figure, there are particles of the compound at the fracture surface. The addition of Al to the Sn-Zn-Ag alloy resulted in the propagation of the crack through the eutectic as can be seen for the 0.1 mass% Al-containing alloy in Fig. 7(c). For the 0.45 mass% Al-containing alloy, the crack propagated 4.1 Effects of Ag and Al on microstructural evolution The addition of 1 mass% Ag to the Sn-Zn alloy gave rise to the needle-like precipitate of the phase, the dendritic block of the compound and the hypo-eutectic structure in the matrix as shown in Fig. 2(b). The formation of the phase in the Sn-Zn-1Ag alloy is induced by the precipitation of the compound. The phase and the compound are observed at the same time, because the addition of 1 mass% Ag is not enough to consume Zn in this alloy completely. On the other hand, no Ag-Sn compound is formed. 28) The formation of the compound lowers the Zn content of the matrix and thus gives rise to the hypo-eutectic structure with coarsened particles of the phase in the matrix. For the Sn- 8.55Zn-1Ag-0.01Al alloy, most of Al was dissolved in the eutectic and primary -Sn phase. Both Al and Zn possess very small solubilities in Sn, and Al has little solubility in Zn. The addition of 0.01 mass% Al to the Sn-8.55Zn-1Ag alloy seems to enforce aggregation of Zn. This is ascribed to the competition for the hypo-eutectic matrix by Al. This behavior results in coarsening of the phase as seen in Figs. 2(c) to (f). The Al-rich segregation starts to show up at 0.45 mass% Al. The Al segregation has also been detected for the Sn-Zn-Al system. 29) The appearance of the Al-rich

5 46 S.-C. Cheng and K.-L. Lin segregation at the addition of 0.45 mass% Al is due to saturation of Al in this alloy. Aluminum has less solid solubility in Sn (1 mass% at 505 K) 24) and Zn (1:5 mass% at 654 K) 25) than in Ag (6 mass% at 723 K). 26) Aluminum may form a compound with Ag, while Ag may also form a compound with Zn. 30) The SEM micrographs in Fig. 2 along with the XRD patterns in Fig. 3 indicate that the compound was formed while no Al-Ag, Al-Zn and Ag-Sn compounds were detected in the Sn-8.55Zn-1Ag-XAl alloys. The Gibbs free energy for the formation of the compound at 773 K G 773 K is 9:535 to 14:414 kj/mol 31) for, 0:359 to 1:782 kj/mol for -Ag 3 Al, 4:724 to 8:025 kj/mol for -Ag 2 Al, 32) 2:698 to 4:149 kj/mol for "-Ag 3 Sn and 0:324 to 2:024 kj/mol for Al-Zn. 33,34) These thermodynamic data explain the preferential formation of the compound rather than that of the Al-Ag, Al-Zn and "- AgSn 3 compounds in the alloys investigated. Consequently, the compound formation between Ag and Zn 23,24) reduces the Zn content of the matrix and enhances the formation of the hypoeutectic structure in the Sn-8.55Zn-1Ag-XAl alloys. 4.2 Variation in mechanical properties The result in Fig. 4 shows that the average values of the UTS, yield stress and hardness of the Sn-8.55Zn-1Ag alloy were the lowest among all the alloys. The tensile strain did not obviously change. The 1 mass% Ag addition to the Sn- 9Zn alloy destroys the uniform eutectic structure and forms the needle-like precipitate of the phase, the dendritic block of the compound, and the hypoeutectic matrix. The average values of the UTS, yield stress and hardness of the Sn-8.55Zn-1Ag Al alloys are improved by the addition of Al. It is believed that the strengthening effect of the Al addition is attributed to the dissolution of Al in Sn ) However, the strengthening effect of Al was not improved by increasing Al content, as the solubility of Al in Sn is rather low. The addition of 0.25 mass% Al to the Sn- 8.55Zn-1Ag alloy could be harmful to the average values of the UTS and yield stress, when the Al content is beyond the solubility limit in Sn. This contributes to the formation of the needle-like precipitate of the phase and the coarsening of the compound, while the eutectic became prominent in the matrix. In the addition of Al up to 0.45 mass%, the segregation of Al, indicating super-saturation, was formed. However, the microstructure with the Znrich phase, the compound and the Al segregation were so uniformly distributed that the average values of the UTS and yield stress in the 0.45Al-containing alloy increase. 38) The tensile strain was obviously decreased because stress concentration occurred at the interface between the compound or segregation and the matrix. 4.3 Fracture behavior The Sn-9Zn and Sn-8.55Zn-1Ag alloys show the crack propagated through the void during necking process as shown in Figs. 7(a) and (b). The compound exists adjacent to the fracture surface of the Sn-8.55Zn-1Ag alloy. The tensile test results in deformation and necking. The internal damage may exist in the form of void or cavity that accumulates while the cross-sectional area reduces. 39) The fracture in engineering alloys can occur through transgranular (through the grain) or intergranular (along the grain ) fracture path. 40) When the Al content increases up to 0.1 to 0.45 mass%, the eutectic became prominent and then the fracture path was formed in the matrix. The 0.1 Al-containing alloy showed clearly the propagation of the crack through the void and the eutectic as shown in Fig. 7(c). Additionally, when overload is a principal cause of fracture, ductile materials fail by a process known as micro-void nucleation at regions of localized strain discontinuity, such as that associated with second-phase particles, inclusions, grain boundaries and dislocation pile-ups. As the strain in the material increases, the micro-void grows, coalesces, and eventually forms a continuous fracture surface. 40) The increasing addition of Al did show association of the fracture surface with the phase and the eutectic, as indicated in Figs. 7(c) and (e). It is believed that the compound, the phase, the Al segregation and the eutectic can all provide the origin for the fracture behavior of the alloys. 5. Conclusions The addition of 1 mass% Ag to the Sn-Zn alloy destroys the eutectic structure and results in the formation of the needle-like precipitate of the phase, the dendritic block of compound, and the hypo-eutectic structure. The variation in the microstructure lowers the UTS (ultimate tensile stress), yield stress and Vickers hardness of the alloy. The increasing addition of Al from 0.01 to 0.25 mass% to the Sn-Zn-Ag alloy tends to enlarge the needle-like precipitate of the phase and the dendrite block of the compound. The UTS, yield stress and Vickers hardness of the 0:010:1 Al-containing alloys were enhanced by the Al addition. The improvement in the mechanical properties was ascribed to solid solution strengthening by dissolution of Al in the hypoeutectic matrix. The mechanical properties of the 0.25 Al-containing alloy were deteriorated by coarsening of the compound and the phase. An increase of the Al content to 0.45 mass% gave rise to precipitation of the Al-rich phase with a near diamond shape. The compound, the phase, the Al segregation and the eutectic provide the origin and path for the fracture of the Sn-8.55Zn-1Ag-XAl alloys. Acknowledgements The financial support of this work from the National Science Council of the Republic of China under the projects NSC E , NSC E , and NSC E is gratefully acknowledged. REFERENCES 1) M. McCormack and S. J. Jin: JOM 45 (1993) ) C. Melten: JOM 45 (1993) ) C. H. Miller, I. E. Anderson and J. F. Smith: J. Electron. Mater. 23 (1994) ) M. McCormack and S. J. Jin: J. Electron. Mater. 23 (1994) ) C. M. Chuang, T. S. Lui and L. H. Chen: J. Electron. Mater. 30 (2001) ) H. Mavoori, S. Vaynman, J. Chin, B. Moran, L. M. Keer and M. E.

6 Microstructure and Mechanical Properties of Sn-8.55Zn-1Ag-XAl Solder Alloys 47 Fine: Mater. Res. Soc. Symp. Proc. 390 (1995) ) H. Mavoori, J. Chin, S. Vaynman, B. Moran, L. M. Keer and M. E. Fine: J. Electron. Mater. 26 (1997) ) S. P. Yu, H. J. Lin, M. H. Hon and M. C. Wang: J. Mater. Sci., Mater. Electron. 11 (2000) ) J. Foley, A. Gickler, F. H. Leprevost and D. Brown: J. Electron. Mater. 29 (2000) ) M. E. Loomans, S. Vaynman, G. Ghosh and M. E. Fine: J. Electron. Mater. 23 (1994) ) M. L. Huang, C. M. L. Wu, J. K. L. Lai, L. Wang and F. G. Wang: J. Mater. Sci., Mater. Electron. 11 (2000) ) K. Suganuma, K. Niihara, T. Shoutoku and Y. Nakamura: J. Mater. Res. 13 (1998) ) K. Suganuma, T. Murata, H. Noguchi and Y. Toyoda: J. Mater. Res. 15 (2000) ) H. M. Lee, S. W. Yoon and B. J. Lee: J. Electron. Mater. 27 (1998) ) K. L. Lin and L. H. Wen: J. Mater. Sci., Mater. Eelctron. 9 (1998) ) K. L. Lin and Y. C. Wang: J. Electron. Mater. 27 (1998) ) S. P. Yu, H. C. Wang, M. H. Hon and M. C. Wang: JOM. 52(June) (2000) ) S. P. Yu, M. C. Wang and M. H. Hon: J. Mater. Res. 16 (2001) ) K. L. Lin, L. H. Wen and T. P. Liu: J. Electron Mater. 27 (1998) ) K. L. Lin and T. P. Liu: Mater. Chem. Phys. 56 (1998) ) K. L. Lin, F. C. Chung and T. P. Liu: Mater. Chem. Phys. 53 (1998) ) M. McCormack and S. J. Jin: J. Electron. Mater. 23 (1994) ) S. C. Cheng and K. L. Lin: J. Electron. Mater. 31 (2002) ) T. B. Massalski, J. L. Murray, L. H. Bennett, H. Baker and L. Kacprzak: 1987) pp ) T. B. Massalski, J. L. Murray, L. H. Bennett, H. Baker and L. Kacprzak: 1987) pp ) T. B. Massalski, J. L. Murray, L. H. Bennett, H. Baker and L. Kacprzak: 1987) pp ) S. K. Bhattacharya and D. F. Baldwin: Adv. Packaging. Sep ) K. L. Lin and C. L. Shih: J. Electron. Mater. 32 (2003) ) K. L. Lin, L. H. Wen and T. P. Liu: J. Electron. Mater. 27 (1998) ) T. B. Massalski, J. L. Murray, L. H. Bennett, H. Baker and L. Kacprzak: 1987) pp ) R. Hultgren, P. D. Desai, D. T. Hawkins, M. Gleiser and K. K. Kelley: (Metal Park, ASM International, OHIO, 1973) pp ) R. Hultgren, P. D. Desai, D. T. Hawkins, M. Gleiser and K. K. Kelley: (Metal Park, ASM International, OHIO, 1973) pp ) R. Hultgren, P. D. Desai, D. T. Hawkins, M. Gleiser and K. K. Kelley: (Metal Park, ASM International, OHIO, 1973) pp ) R. Hultgren, P. D. Desai, D. T. Hawkins, M. Gleiser and K. K. Kelley: (Metal Park, ASM International, OHIO, 1973) pp ) Y. Kariya and M. Otsuka: J. Electron. Mater. 27 (1998) ) Y. Kariya and M. Otsuka: J. Electron. Mater. 27 (1998) ) Y. Kariya, Y. Hirata and M. Otsuka: J. Electron. Mater. 28 (1999) ) M. McCormack, S. J. Jin, G. W. Kammlott and H. S. Chen: Appl. Phys. Lett. 63 (1993) ) G. W. Powell, S. E. Mahmoud and K. Mills: Metals Handbook, 9th edn, Vol. 8, (Metals Park, ASM International, OHIO, 1986) pp ) G. W. Powell, S. E. Mahmoud and K. Mills: Metals Handbook, 9th edn, Vol. 12, (Metals Park, ASM International, OHIO, 1986) pp