Fracture Behavior of Sn-3.5Ag-0.7Cu and Pure Sn Solders as a Function of Applied Strain Rate

Fracture Behavior of Sn-3.5Ag-0.7Cu and Pure Sn Soldersas a Function of Applied Strain Rate

K.E. YAZZIE,1 J.J. WILLIAMS,1 and N. CHAWLA1,2

1.—Materials Science and Engineering, School for the Engineering of Matter, Transport,and Energy, Arizona State University, Tempe, AZ 85287-6106, USA. 2.—e-mail: [email protected]

The demand for environmentally benign Pb-free solders is increasing, and thepush toward smaller portable electronics will make it more likely for solderinterconnects to encounter mechanical shock through dropping or mishan-dling. Thus, fundamental understanding of the relationship between soldermicrostructure and mechanical shock resistance is essential for developingreliable numerical models of mechanical shock behavior. In this paper wereport on the strain rate-dependent mechanical behavior of pure Sn andSn-3.5Ag-0.7Cu solders, measured from tensile tests conducted in the strainrate range from 10�3 s�1 to 30 s�1. Local strain and strain rate distributionswere measured by digital image correlation. Finally, the strain rate depen-dence of fracture mechanisms is discussed. For a given strain rate, water-quenched tin–silver–copper (SAC) had the greatest ultimate tensile strength(UTS), followed by furnace-cooled SAC, then pure Sn. Furnace-cooled SAC hadlower ductility than water-quenched SAC, due to large Ag3Sn needles thatnucleated elongated voids which easily coalesced.

Key words: Pb-free solder, mechanical shock, fracture

INTRODUCTION

The need to develop environmentally benignelectronic packages has generated great interest inPb-free alloys.1–6 As electronic packages are madesmaller for portable devices, there is an increasedprobability that solder joints may fail by accidentaldropping during manufacture, shipping, or use. Thestrain rates experienced by solders during drop, i.e.,mechanical shock, are in an intermediate rangebetween the quasistatic and dynamic regimes, i.e.,between 10�1 s�1 and 102 s�1.7–13 The microstruc-ture, creep, and thermal fatigue behavior of Pb-freesolders is well understood,14–17 and a good under-standing of the mechanical shock behavior of Pb-freesolders is being established.18–22 However, funda-mental understanding of the relationship betweensolder microstructure and mechanical shock resis-tance is needed. Understanding this relation-ship requires quantitative analysis of strainrate-dependent fracture mechanisms.

In this study the mechanical behavior of Sn-3.5Ag-0.7Cu solder was systematically studied, and com-pared with that of pure Sn, at strain rates rangingfrom 10�3 s�1 to 30 s�1. Local values of strain andstrain rate were measured at the onset of necking, inthe necking region of the tensile specimen, usingdigital image correlation (DIC). Fracture surfaces offailed tensile specimens were quantitatively ana-lyzed using scanning electron microscopy (SEM).The first part of the paper describes the micro-structural characterization and analysis of experi-mental tensile test data. The second part discussesstrain rate-dependent fracture mechanisms.

EXPERIMENTAL PROCEDURES

Sn (99.999% pure; Alfa Aesar, Ward Hill, MA,USA) and Sn-3.5Ag-0.7Cu (SAC; Indium Corpora-tion, Ithaca, NY, USA) ingots were used in thisstudy. The ingots were reflowed in a graphite-coatedaluminum mold with the following dimensions:10.5 cm long, 1 cm wide, and 0.8 cm high. A ther-mocouple placed at the bottom of the solder was

(Received March 3, 2012; accepted June 14, 2012;published online July 6, 2012)

Journal of ELECTRONIC MATERIALS, Vol. 41, No. 9, 2012

DOI: 10.1007/s11664-012-2180-9� 2012 TMS

2519

used to measure the cooling rate. The samples wereheated to 20�C above the melting point, held for20 s, and cooled. Fine and coarse microstructureswere obtained with a water-quench cooling rate of16.5�C/s and a furnace cooling rate of 1.0�C/s,respectively. A fine microstructure was produced inpure Sn by recrystallization after cold working 35%by rolling. Pure tin specimens were prepared by coldworking to obtain a more uniform grain size in theingots from which the specimens were machined.The samples were polished to a final finish using a0.05-lm colloidal silica solution. SEM (FEI-XL30)was used to characterize the final microstructuresand to perform fractography.

Tensile specimens were machined from a sectionnear the bottom of the reflowed blank, where thecooling rate was measured. Microstructural charac-terization of the tensile specimens prior to testingrevealed uniform microstructure throughout thespecimens. The tensile specimens had 10 mm gagelength. A detailed schematic is available else-where.18 Tensile specimens were coated with a whitespray paint base, followed by black spray paintspeckle, for DIC analysis. A high-speed camera(Phantom; Vision Research, Wayne, NJ) recordedthe deformation of the tensile specimen at framerates of up to 4200 frames per second. Tensile tests

were performed on a MTS 810 servohydraulic loadframe at nominal strain rates ranging from 10�3 s�1

to 30 s�1. Inertial effects were negligible for the bulksamples, such that the strain rate was linear duringthe entire test. Details about controlling strain rateare available elsewhere.18 The strain produced inthe specimen during the tensile test was analyzedusing commercially available DIC software (AR-AMIS; Trillion Quality System, Plymouth Meeting,PA, USA). Details about the DIC parametersare available elsewhere.19 Energy-dispersive x-rayspectroscopy (EDS) was used to identify precipitateson the fracture surfaces.

RESULTS AND DISCUSSION

Tensile Behavior

The microstructures of pure Sn, water-quenchedSAC, and furnace-cooled SAC are shown in Fig. 1.Pure Sn had final grain size of approximately400 lm, as shown in Fig. 1a. The water-quenchedSAC microstructure consisted of Sn-rich dendritesand a eutectic mixture of Sn, Cu6Sn5, and Ag3Sn, asshown in Fig. 1b. The furnace-cooled SAC micro-structure consisted of Ag3Sn needles and Cu6Sn5

intermetallic several micrometers in size, asshown in Fig. 1c. Figure 2a–c shows representative

Fig. 1. (a) Microstructure of recrystallized pure Sn. Average grain size was approximately 400 lm. (b) Microstructure of furnace-cooledSn-3.5Ag-0.7Cu. Large Ag3Sn needles and blocky Cu6Sn5 particles exist in a Sn matrix. (c) Microstructure of water-quenched Sn-3.5Ag-0.7Cu.Sn-rich dendrites are surrounded by a eutectic mixture of Sn, Ag3Sn, and Cu6Sn5.

Yazzie, Williams, and Chawla2520

engineering stress–strain curves for tensile tests ofpure Sn, water-quenched SAC, and furnace-cooledSAC, respectively. At the higher strain rates, oscil-latory waves caused by reflected elastic waves weresuperimposed on the stress–strain curves. In gen-eral, all specimens showed an increase in flow stresswith increasing strain rate. Pure Sn had the lowestflow stress for a given strain rate. Furnace-cooledSAC had higher ultimate tensile strength (UTS)values than pure Sn, due to strengthening fromlarge Ag3Sn needles and Cu6Sn5 particles. Water-quenched SAC had the largest UTS for a given strainrate, due to a finer microstructure, particularly inthe form of a fine dispersion of Ag3Sn and Cu6Sn5

particles, which increased the number of obstaclesfor dislocation motion.

Figure 3 shows the UTS, averaged from atleast three tests, plotted as a function of appliedstrain rate for pure Sn, and water-quenched and

furnace-cooled SAC. The slope of the Log–Log plotcorresponds to the strain rate sensitivity, m, whichwas found to be 0.15 for pure Sn, 0.10 for water-quenched SAC, and 0.08 for furnace-cooled SAC.These values are similar to those found in the lit-erature for Sn,23,24 and for Sn–Ag–Cu solders withsimilar composition tested at low25–28 and highstrain rates.20–22 Empirically, strain rate sensitivityindicates a material’s ability to resist tensile insta-bility (necking). Higher m also typically indicates adelayed onset of tensile failure, and hence greaterductility. Therefore, the strain rate sensitivitiesmeasured here indicated that Sn should have thegreatest ductility, followed by water-quenched SAC,and then furnace-cooled SAC. While classicalstrengthening theory states that materials with finemicrostructures should have less ductility thanmaterials with coarse microstructures, the initialanalysis of the strain rate sensitivities indicated

Fig. 2. Representative stress–strain curves for (a) pure Sn, (b) furnace-cooled SAC, and (c) water-quenched SAC. All stress–strain curves showincreasing flow stress with strain rate. It is interesting to note that, in general, furnace-cooled SAC is less ductile than water-quenched SAC.

Fracture Behavior of Sn-3.5Ag-0.7Cu and Pure Sn Solders as a Function of Applied Strain Rate 2521

that the reverse was true in the case of water-quenched and furnace-cooled SAC. Indeed, thisstudy will show that water-quenched SAC did havegreater strain to failure than furnace-cooled SAC.

Far-Field and Local Strain and Strain RateDistributions

Far-field strain was measured from the displace-ment of the MTS 810 linear variable differentialtransducer (LVDT). The far-field strain coincidingwith the onset of necking (UTS) was measured, sincethis can be an important constitutive input for reli-ability models. The far-field strain at UTS is plottedas a function of the applied strain rate in Fig. 4. The

far-field strain at UTS was greatest for pure Sn, whilethe strains for furnace-cooled and water-quenchedSAC were essentially equal. Interestingly, the far-field strain at UTS did not show strain rate depen-dence. However, when a local measurement of thestrain at necking was made using DIC, the strainrate-dependent deformation became apparent. Thelocal strain in the neck, measured at the onset ofnecking (UTS), is plotted in Fig. 5a as a function ofthe applied strain rate. These are averages of thelocal strains measured only in the necking region ofthe tensile specimen (identified from high-speedvideo). Again, pure Sn had the greatest strain, while

Fig. 3. Log UTS plotted as a function of Log applied strain rateto show the strain rate dependence of the flow stress.

Fig. 4. Far-field strain computed using displacement from LVDT.The far-field strain corresponding to necking is plotted as a functionof the applied strain rate.

Fig. 5. (a) Local strain measured in the necking region of the sample at the onset of necking computed using DIC and plotted as a function of theapplied strain rate. (b) Strain contours computed using DIC and corresponding to the onset of necking for pure Sn, furnace-cooled SAC, andwater-quenched SAC tested at 10�1 s�1. The trend observed in (a) is visualized using these contour plots.


the strains for furnace-cooled and water-quenchedSAC were essentially equal. Local strain measure-ments showed that there was an increase in the strainin the neck with the applied strain rate. This trend isvisualized in Fig. 5b, which shows images of the gagesection of the tensile specimen at the onset of necking.Color contours corresponding to the longitudinal (ex)strain distribution are overlaid on the gage sections.The fact that the far-field and local strains at neckingfor furnace-cooled and water-quenched SAC wereequal can be explained by considering that, while voidnucleation can occur at low plastic strains, significantvoid growth occurs rapidly after the onset of neck-ing.29–32 Therefore, it is not expected that the strainswould be appreciably different for furnace-cooled andwater-quenched SAC up to necking. However, micro-structure-dependent differences in void growthbehavior after necking should be reflected in thestrain-to-failure measurements.

The local strain rate in the neck, at the onset ofnecking, was measured from DIC. The local strainrate in the neck is plotted as a function of theapplied strain rate in Fig. 6a. The strain rate in theneck was very close to the applied strain rate.Again, this is not surprising since the measure-ments were taken at the onset of necking. It wouldbe expected that at later stages of necking the localstrain rates would increase significantly. Some dif-ficulty was encountered in carrying out DIC analy-sis beyond necking and close to failure of thespecimen. The large deformation in the gage sectiontended to spall the speckle pattern, making it diffi-cult to continue DIC. Figure 6b shows strain ratecolor contours overlaid on images of the specimens.It should be noted that the strain rate concentra-tions computed by DIC, shown in Fig. 6b, coincidedwith the actual strain localizations in Fig. 5b.

The strain to failure, measured from the LVDT, isplotted as a function of the applied strain rate inFig. 7. Pure Sn had the greatest ductility for a givenstrain rate, followed by water-quenched SAC, andthen furnace-cooled SAC. This is an interestingresult considering that classical strengthening the-ory states that materials with coarser microstruc-tures usually exhibit greater ductility. The reasonfor this apparent contradiction can be attributed tothe void growth mechanisms which are expoundedupon in the fractography section. There does not

Fig. 6. (a) Local strain rate in the neck computed using DIC and corresponding to the onset of necking as a function of the applied strain rate.The local strain rate is close to the applied strain rate. (b) Strain rate contours computed using DIC and corresponding to the onset of necking forpure Sn, furnace-cooled SAC, and water-quenched SAC.

Fig. 7. Strain to failure measured from displacement of LVDT as afunction of applied strain rate. Though strong strain rate dependencewas observed in the flow stress, there appears to be no strain ratedependence in the macroscopic strain to failure within the range ofstrain rates tested.


appear to be strain rate dependence of the strainto failure. Other researchers have found similarresults in the strain to failure for bulk Sn–Ag–Cualloys at low27 and high33,34 strain rates; however,no explanations were given for this behavior. In thecase of bulk Sn-3.5Ag-0.7Cu, it will be shown thatfracture is controlled by nucleation of voids atCu6Sn5 and Ag3Sn precipitates. This mechanism isnot applicable in the case of pure Sn though. How-ever, it may be reasonable to postulate that there isno significant strain rate-dependent ductility overthe range of strain rates tested, but that at higherstrain rates some effect may be observable.

Strain Rate Dependence of FractureMechanisms

Ductile deformation was initiated at Cu6Sn5 andAg3Sn precipitates, as identified by EDS spot scans.In particular, the fracture behavior was controlled

by Ag3Sn due to its higher volume fraction inSn-3.5Ag-0.7Cu solder. The evidence for this wasfine, submicrometer Ag3Sn precipitates found at thebottom of ductile dimples in water-quenched SAC,or in the case of furnace-cooled SAC, large Ag3Snneedles at the bottom of elongated ductile dimples.The ductile dimples in water-quenched SAC becamemore numerous with increasing strain rate. In fur-nace-cooled SAC, large Ag3Sn needles began tofracture along their length, with the average frac-tures generally increasing with strain rate. Theelongated nature of the Ag3Sn needles appears to beresponsible for nucleation of larger voids, and thus asmaller intervoid spacing, which could coalescemore easily, resulting in lower ductility. In bothcases, increasing strain rate caused an increase inthe stress triaxiality that developed around theprecipitates, thereby initiating more numerousductile voids in water-quenched SAC and fracturingAg3Sn needles in furnace-cooled SAC.

Fig. 8. (a) Average number of fractures per length of Ag3Sn needle, analyzed from fractography of furnace-cooled SAC and plotted as a functionof strain rate. (b) Example of intact Ag3Sn needle in ductile dimple from 10�3 s�1 test. (c) Example of fractured Ag3Sn needle in ductile dimplefrom 30 s�1 test.


Comprehensive fractographic analysis of all thefailed tensile specimens was performed to quantifythe strain rate dependence of the aforementionedfracture mechanisms. For each specimen, a mini-mum of 10 measurements were made in 10 randomlocations at magnification of 10009, correspondingto a minimum area of 0.12 mm2 analyzed. In fur-nace-cooled SAC the number of fractures in a Ag3Snneedle was counted, then divided by the total lengthof the needle, to compute the average fractures permicron of length. The average fractures per micronof length of Ag3Sn needle is plotted as a function ofthe applied strain rate in Fig. 8a. The averagefractures in Ag3Sn needles increased monotonicallywith strain rate. The relatively large standarddeviations may be due to the random orientations ofthe Ag3Sn needles in the as-cast ingots. Needlesthat are aligned with the loading axis will bear moreload than those that are not aligned.35 Figure 8band c show representative SEM micrographs of thefracture surface of furnace-cooled SAC specimens

tested at 10�3 s�1 and 30 s�1, respectively. In Fig. 8bthe Ag3Sn needles were mostly intact and sat inshallow ductile dimples. However, in Fig. 8c theAg3Sn needles were fractured into multiple piecesand sat in deep dimples, indicating the effect ofincreasing stress triaxiality with strain rate.

In water-quenched SAC the number of voids in agiven micrograph were counted, then divided by thetotal area of the micrograph to compute the averagenumber of voids per lm2. The average number ofvoids per lm2 is plotted as a function of the appliedstrain rate in Fig. 9a. The number of ductile voidsinitiated by spherical Ag3Sn precipitates increasedmonotonically with strain rate. The large standarddeviations at 10�1 s�1 and 1 s�1 are due to partic-ularly uneven fracture surfaces for these two spec-imens. Figure 9b and c show representative SEMmicrographs of the fracture surface of water-quen-ched SAC specimens tested at 10�3 s�1 and 30 s�1,respectively. Figure 9b shows ductile dimples thatare relatively large and shallow, some on the order

Fig. 9. (a) Average number of voids per area from fractography of water-quenched SAC, plotted as a function of strain rate. (b) Example of lowaverage voids per unit area (�0.004 lm�2) in 10�3 s�1 test. (c) Example of high average voids per unit area (�0.016 lm�2) from 30 s�1 test.


of 20 lm in diameter. Figure 9c shows a very roughfracture surface composed of ductile dimples thatare relatively small and deep. The smaller, morenumerous ductile dimples were due to developmentof higher stress triaxiality around the small,spherical Ag3Sn precipitates as the strain rate wasincreased.

CONCLUSIONS

Tensile tests of pure Sn and Sn-3.5Ag-0.7Cu spec-imens with coarse and fine microstructures wereconducted over the strain rate range from 10�3 s�1 to30 s�1 to understand the relationship betweenmicrostructure and strain rate-dependent mechani-cal behavior. Local strain and strain rate evolution atthe onset of necking were analyzed using DIC. Frac-ture surfaces were quantitatively analyzed to deter-mine the effect of microstructure and strain rate onfracture mechanisms. The following conclusionswere drawn from the experimental results:

1. Flow stress increased with increasing appliedstrain rate. For a given strain rate, water-quenched SAC had the greatest UTS, followedby furnace-cooled SAC, and then pure Sn. Strainrate sensitivity measured from plots of Log UTSversus Log strain rate indicated that water-quenched SAC should have greater ductilitythan furnace-cooled SAC.

2. While the far-field strain at the onset of neckingdid not show strain rate dependence, local strainin the neck increased with applied strain rate.The strain rate in the neck was close to theapplied strain rate.

3. Furnace-cooled SAC had lower ductility thanwater-quenched SAC, due to large Ag3Sn needlesthat nucleated elongated voids which easilycoalesced. At the lowest strain rate Ag3Sn nee-dles were found intact in the ductile dimples, butbegan to fracture along their length with increas-ing strain rate. This fracture behavior wasquantified by measuring the average fracturesper micron length, which increased monotoni-cally with the Log of the applied strain rate. Inwater-quenched SAC, small ductile voids werenucleated by submicrometer Ag3Sn precipitates.The ductile dimples became smaller, deeper, andmore numerous with increasing strain rate. Thisfracture behavior was quantified by measuringthe average number of voids per lm2 area, whichincreased monotonically with the Log of theapplied strain rate.

ACKNOWLEDGEMENTS

The authors are grateful for the financial supportfor this work from the National Science FoundationDivision of Materials Research—Metals Division(Drs. Alan Ardell, Bruce MacDonald, Harsh Chopra,

and Eric Taleff, Program Directors). The authorsgratefully acknowledge the use of facilities withinthe Center for Solid State Science at Arizona StateUniversity.

REFERENCES

1. N. Chawla, Int. Mater. Rev. 54, 368 (2009).2. J. Glazer, J. Electron. Mater. 23, 693 (1994).3. J. Glazer, Int. Mater. Rev. 40, 65 (1995).4. M. McCormack and S. Jin, JOM 45, 36 (1993).5. P.R. Vianco and D.R. Frear, JOM 45, 14 (1993).6. W.J. Plumbridge, J. Mater. Sci. 31, 2501 (1996).7. T.Y. Tee, H.S. Ng, C.T. Lim, E. Pek, and Z. Zhong, Micro-

electron. Reliab. 44, 1131 (2004).8. D. Reiff and E. Bradley, Electron. Compon. Technol. Conf.

Proc. 55, 1519 (2005).9. M. Date, T. Shoji, M. Fujiyoshi, K. Sato, and K.N. Tu,

Electron. Compon. Technol. Conf. Proc. 54, 668 (2004).10. K.T. Tsai, F.L. Liu, E.H. Wong, and R. Rajoo, Solder. Surf.

Mt. Technol. 18, 12 (2006).11. P. Pandher and M. Boureghda, Reliab. Phys. Symp. 45, 107

(2007).12. K. Newman, Electron. Compon. Technol. Conf. Proc. 55,

1194 (2005).13. J.Y.H. Chia, B. Cotterell, and T.C. Chai, Mater. Sci. Eng. A

417, 259 (2006).14. R.S. Sidhu and N. Chawla, Metall. Mater. Trans. A 39, 340

(2008).15. R.S. Sidhu, X. Deng, and N. Chawla, Metall. Mater. Trans.

A 39, 349 (2008).16. R.S. Sidhu and N. Chawla, Metall. Mater. Trans. A 39, 799

(2008).17. M. Kerr and N. Chawla, Acta Mater. 52, 4527 (2004).18. K. Yazzie, H. Fei, J. Williams, H. Jiang, and N. Chawla,

J. Electron. Mater. 38, 2746 (2009).19. K.E. Yazzie, J.J. Williams, D. Kingsbury, P. Peralta,

H. Jiang, and N. Chawla, JOM 62, 16 (2010).20. X. Long, I. Dutta, V. Sarihan, and D.R. Frear, J. Electron.

Mater. 37, 189 (2007).21. E.H. Wong, C.S. Selvanayagam, S.K.W. Seah, W.D. van

Driel, J.F.J.M. Caers, X.J. Zhao, N. Owens, L.C. Tan, D.R.Frear, M. Leoni, Y.-S. Lai, and C.-L. Yeh, Mater. Lett. 62,3031 (2008).

22. F. Qin, T. An, and N. Chen, J. Appl. Mech. 77, 011008(2010).

23. M.J. Mayo and W.D. Nix, Acta Metall. 36, 2183 (1988).24. R.E. Frenkel, O.D. Sherby, and J.E. Dorn, Acta Metall. 3,

470 (1955).25. N. Bai, X. Chen, and Z. Fang, J. Electron. Mater. 37, 1012

(2008).26. K. Kim, S. Huh, and K. Suganuma, Mater. Sci. Eng. A 333,

106 (2002).27. I. Shohji, T. Yoshida, T. Takahashi, and S. Hioki, Mater. Sci.

Eng. A 366, 50 (2004).28. M. Amagai, M. Watanabe, M. Omiya, K. Kishimoto, and T.

Shibuya, Microelectron. Reliab. 42, 951 (2002).29. V. Tvergaard and A. Needleman, Acta Metall. 32, 157

(1984).30. W.M. Garrison Jr. and N.R. Moody, J. Phys. Chem. 48, 1035

(1987).31. A.A. Benzerga, J. Besson, and A. Pineau, Acta Mater. 52,

4623 (2004).32. A. Weck, D.S. Wilkinson, E. Maire, and H. Toda, Acta Mater.

56, 2919 (2008).33. B.L. Boyce, L.N. Brewer, M.K. Neilsen, and M.J. Perricone,

J. Electron. Packag. 133, 031009-1 (2011).34. F. Qin, T. An, N. Chen, and J. Bai, J. Electron. Packag. 131,

031001-1 (2009).35. N. Chawla, F. Ochoa, V.V. Ganesh, X. Deng, M. Koopman,

K.K. Chawla, and S. Scarritt, J. Mater. Sci. Mater. Electron.15, 385 (2004).


Documents

Fracture Behavior of Sn-3.5Ag-0.7Cu and Pure Sn Solders as a Function of Applied Strain Rate