Solid-State Reactions between Cu(Ni) Alloys and Sn

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1 ournal of ELECTRONIC MATERIALS, ol. 3, No. 1, 27 DOI: 1.17/s Ó 27 TMS Regular Issue Paper Solid-State Reactions between (Ni) Alloys and ESA UORINEN, 1,3 TOMI LAURILA, 1 TONI MATTILA, 1 ERKKI HEIKINHEIMO, 2 and ORMA K. KIILAHTI 1 1. Laboratory of Electronics Production Technology, Helsinki University of Technology, P.O. Box 3, Espoo, 215 HUT, Finland. 2. Laboratory of Metallurgy, Helsinki University of Technology, P.O. Box 3, Espoo, 215 HUT, Finland esa.uorinen@tkk.fi Solid-state interfacial reactions between and (Ni) alloys have been investigated at the temperature of 125 C. The following results were obtained. Firstly, the addition of.1 at.% Ni to decreased the total thickness of the intermetallic compound (IMC) layer to about half of that observed in the binary / diffusion couple; the Ni addition decreased especially the thickness of 3. Secondly, the addition of 1 to 2.5 at.% Ni to further decreased the thickness of 3, increased that of 5 (compared to that in the binary / couple) and produced significant amount of voids at the / 3 interface. Thirdly, the addition of 5 at.% Ni to increased the total thickness of the IMC layer to about two times that observed in the binary / diffusion couple and made the 3 disappear. Fourthly, in contrast to the previous case, the addition of 1 at.% Ni to decreased the total IMC ( 5 ) thickness again close to that of the / couple. With this Ni content no voids were detected. The results are rationalized with the help of the thermodynamics of the --Ni system as well as with kinetic considerations. Key words: Intermetallic reactions, --Ni system, Kirkendall voids, thermodynamics, kinetics INTRODUCTION Copper is the most common conductor metal. Because of its good electrical conductivity, solderability and processbility it is generally used in printed wiring boards (PWB) and component metallizations, where it is in contact with solder alloys. The dissolution rate of to liquid solder is, however, relatively fast. 1,2 This, especially in fine-pitch applications, can lead to a situation where too large a fraction of the solder joint is transformed into intermetallic compounds (IMCs) or the entire metallization layer has dissolved into liquid solder. To circumvent this possibility, electrochemical or electroless Ni metallizations with Au surface finishes have been used as diffusion barrier layers owing to NiÕs slower dissolution into -based solders compared to. 1,2 However, the usage of electroless or electrochemical Ni gives rise to other problems. For example, when Ni metallizations are (Received February 2, 27; accepted une 28, 27; published online September 11, 27) used with -based solders that do not contain, redeposition of (Au,Ni) 4 on top of the Ni 3 4 can take place. 3 5 On the other hand, when Ni/Au is used with -containing solders, (Ni,) 3 4, (,Ni) 5 or both IMC layers can form on top of the Ni metallization depending on the content of the solder. 9 Finally, when electroless Ni(P) is used, the interfacial reaction product layer can become structurally complex and brittle If the good wetting properties of could be feasibly combined with the slower dissolution and reaction kinetics of Ni, (Ni) alloys would be used as metallization material. In this regard, Ohriner 13 has carried out an extensive investigation of the reactions between various -based alloys and -containing solders. The substrates were exposed to liquid solder for 5 s at a temperature of 3 C above the liquidus temperature of the solder alloy and subsequently annealed in the solid state at different temperatures ranging from 15 C to 225 C. 13 He observed interesting behaviour when Ni-containing alloys of were soldered with the 95-5Ag, 95-5Sb and -Pb solders. The rate of intermetallic formation was found to be dependent 1355

2 135 uorinen, Laurila, Mattila, Heikinheimo, and Kivilahti on solder composition, temperature, and the Ni content of the (Ni) alloys. The maximum growth rate was always between % and 9 at.% of Ni in. 13 The change in the reaction layer thickness was suggested to be related to the changes in the composition of the intermetallic compound and that the addition of nickel to the 5 can change the concentration of structural vacancies in the IMC. 13 Another important observation made by Ohriner was that the only IMC to grow between (Ni) alloys and 95-5Ag, 95-5Sb, and -Pb solders was (,Ni) 5 while 3 was not detected within the resolution limits of scanning electron microscopy (SEM). 13 When considering these results it is to be noted that is known to be the faster diffusing element in 5, whereas is the main diffusing species in 3 at temperatures above 1 C 14,15 and that Yu et al. have demonstrated that the driving force for the diffusion of through 5 is increased, when the Ni content of the phase is increased. 1 These factors are addressed in detail later in this paper. Additionally, Paul has made interesting observations when studying solid-state reactions (T = 225 C) between pure and alloyed with 5 to 15 at.% Ni. 14 He also detected significant increase in the growth rate of (,Ni) 5 within this concentration range, as well as the absence of 3. Paul also observed that the grain size of (,Ni) 5 was more than one order of magnitude smaller than the grain size in the case of. 14 We have also observed the high growth rates of (,Ni) 5 ) within the same concentration range, not only in the solid state but also as in solid/liquid reaction couples. 17 Recently, some authors have observed the formation of voids in 3 or at the / 3 interface in the case of Pb and Ag solders. 7,18 Related to the formation of Kirkendall voids, Paul 14 has shown that Kirkendall planes (of constant compoisition) in a given reaction couple can be multiple, stable, unstable or virtual. He used an approach based on the so-called Kirkendall velocity plot 19,2 to explain and predict the Kirkendall plane formation. This velocity plot can be constructed from the available experimental data (intrinsic diffusion coefficients, marker displacement, etc.). 14,19,2 Based on the calculations by Paul et al. 14 in the binary / system no stable Kirkendall plane should form inside the 3 phase. The purity of is important, since the results from the velocity construction are strongly dependent, among other factors, on the compositions of the endmembers of the diffusion couple. In fact, Paul 14 detected extensive porosity inside 3 when the layer in the /1%Ni diffusion couple. Accordingly, since the interfacial intermetallic structure can drastically influence the solder joint reliability especially under mechanical shock loading conditions 21 we will study systematically the effect of Ni (.1 at.% to 1 at.%) in, firstly, on the growth of (,Ni) 5, secondly, on the possible absence of 3, and, thirdly, on the formation of voids inside the reaction zone at a temperature of 125 C. In order to rationalize the experimental results obtained, ternary --Ni phase diagram data together with available diffusion kinetic information are utilized. MATERIALS AND METHODS The (Ni) alloys were made from commercially pure metal powders (99.9%+, entron corporation) by melting the premixed powders in a Al 2 O 3 crucible for 4 h at 14 C under a 9%Ar1%H 2 atmosphere. The (Ni) alloys as well as high-purity (99.99%+, Outokumpu) copper were cut into 1-mmthick plates (25 mm 2 ), which were placed on a PWB. Prefluxed (rosin mildly activated) 25-lm-thick pure (99.95%+, Goodfellow Ltd.) foils were set on top of the (Ni) plates. The soldering was carried out in a conventional forced convection reflow oven (EPM/EWOS 5.1 N 2 ) under air atmosphere. The reflow profiles measured from three different locations on the board are presented in Fig. 1. After the soldering the samples were annealed in a forced convection oven (Heraeus Instruments UT) for different periods of time up to 35 h at 125 C. The cross-sections of the samples were prepared with the standard metallographic methods, and they were analyzed with a field-emission scanning electron microscope (EOL 335F FE-SEM) equipped with an energy-dispersive spectrometer (EDS, Oxford ISIS). The reported IMC thicknesses are the average values of 2 measurements. The amount of voids is defined as the estimated cross sectional area of voids relative to the area of 3. In order to reveal the intermetallic compounds more clearly some of the samples were etched with a dilute HCL/ methanol solution. RESULTS The following results were obtained from the experiments carried out with several (Ni)/ diffusion couples at the temperature of 125 C. When high-purity was soldered with pure and the diffusion couple was subsequently annealed for Fig. 1. The measured reflow profile from three different locations on the PWB.

3 Solid-State Reactions between (Ni) Alloys and h the resulting reaction layers were as shown in Fig. 2a. Both 5 and 3 are clearly visible. If one compares the thickness ratio of 5 to 3 to that typically observed after reflow (Fig. 2b), it is evident that the relative thickness of the 3 layer has increased markedly during the solid-state ageing. The total thickness of the IMC layers after 35 h of annealing is about 7 lm, i.e., about three times of the thickness after reflow. These results are in agreement with the previous studies on the kinetics of / reactions available in the literature. 22,23 When.1 at.% of Ni is added to and the diffusion couple is annealed after soldering for 35 h the resulting interfacial reaction layers are as shown in Fig. 3. Both 5 and 3 are present in the diffusion couples. What is different, as compared to the binary case, is that the relative thickness of 3 is much smaller. Thus, it is evident that the addition of.1 at.% of Ni has a marked effect on the growth kinetics of the IMC layers. If one compares the total thickness of the reaction layers in 1%Ni/ to those in binary / couple, it can be observed (Figs. 2 and 3) that the total IMC thickness has also decreased, due to the reduction in the thickness of 3. The average thickness of 5 is, however, about the same as in the / couple. By increasing the Ni content in up to 1 at.%, the interfacial microstructure after the solid-state annealing continues to change (see Fig. 4). In addition to the reduction in the thickness of 3, there is a marked density of voids, mostly located at the / 3 interface or inside the 3 layer. Some authors have claimed these to be Kirkendall voids, but as has already been discussed elsewhere 7 the root cause is likely more complex. The addition of 2.5 at.% of Ni to does not provide any major microstructural changes in comparison to the 1%Ni case, as can be seen from Fig. 5. Both compounds are still present, but the 3 is very thin and its porosity is even higher. The 5 shows a wide variation in thickness, but the average thickness is somewhat larger than in the 1%Ni case. However, when the Ni concentration in the (Ni) alloy is increased to 5 at.% (for 35 h), 3 is only barely detectable. The thickness of the (,Ni) 5 (containing about to 8 at.% Ni) is much larger than in the previous cases (Fig. ). The amount of voids has reduced significantly. Increasing the Ni content to 1 at.% does not change the microstructure markedly. (,Ni) 5 is the only phase present but its thickness has decreased and there are hardly any voids (Fig. 7). Accordingly, Fig. 8a and b sums up the changes in the reaction kinetics in diffusion, when Ni is added to the following phenomena take place in the (Ni)/ diffusion couples: (1) the addition of Fig. 2. Backscattered SEM images taken from the sample where high-purity is soldered with pure. (a) After reflow the diffusion couple has been subsequently annealed at 125 C for 35 h and (b) interfacial reaction zone after reflow. Fig. 3. Backscattered SEM images taken from the Ni/ sample annealed at 125 C for 35 h. Fig. 4. Backscattered SEM image taken from the 99.1.Ni/ sample annealed at 125 C for 35 h showing the formation of voids at the / 3 interface..1 at.% of Ni to decreased the total thickness of the IMC layer to about half of that observed in the binary / diffusion couple; the Ni addition decreased especially the thickness of 3 ; (2) the

4 1358 uorinen, Laurila, Mattila, Heikinheimo, and Kivilahti 1 9 a 8 voids of Precentage Fig. 5. Backscattered SEM image taken from the Ni/ sample annealed at 125 C for 35 h Ni +1.Ni +2.5Ni +5.Ni +1Ni b I MC thickness [ µ m ] Ni +1.Ni +2.5Ni +5.Ni +1Ni Fig.. Backscattered SEM image taken from the 95.5.Ni/ sample annealed at 125 C for 35 h. I MC thickness [ µ m ] c 2,1 1 2,5 5 1 Ni content of the (Ni) alloys [at.%] Fig. 8. (a) The percentage of voids at the / 3 interface and (b) IMC thickness versus the Ni content of the (Ni) alloy substrate after annealing for 35 h at 125 C. Fig. 7. Backscattered SEM image taken from the 91Ni/ sample annealed at 125 C for 35 h. addition of 1 to 2.5 at.% of Ni to further decreased 3, increased the thickness of 5 (in comparison to that in the binary / couple) and produced significant amount of voids at the / 3 interface; (3) the addition of 5 at.% Ni to increased the total thickness of the IMC layer to about two times of that observed in the binary / diffusion couple and made 3 to disappear, and (4) the addition of 1 at.% of Ni to decreased the total IMC ( 5 ) thickness again close to that of the / couple. With this Ni content no voids were detected.

5 Solid-State Reactions between (Ni) Alloys and 1359 DISCUSSION The reason for the absence of 3 in (Ni) alloys with more than 5 at.% Ni can be thermodynamic, kinetic, or a combined effect, as suggested also by Ohriner. 13 In order to study the thermodynamics behind this effect, the relative stabilities of the phases (i.e., b-, (,Ni) 3, (,Ni) 5, (Ni,) 3 4, (Ni,) 3 2, (Ni,) 3, and Ni 5 2 ) in the --Ni system at 125 C have been assessed by making use of the experimental information and the available thermodynamic data. 7,17,24 28 The results of the assessment are presented as the metastable isothermal section at 125 C in Fig. 9. This metastable diagram includes neither the solid miscibility gap in the -Ni system nor the ternary phase Ni (s). 24 This is because the formation of the s-phase equires very long annealing times (up to 1, h) and, furthermore, this compound has not been reported elsewhere. The other ternary compound, Ni 5 2 (C 1 ) 28 was taken into account in the assessment of the system for the reasons discussed later. It should be noted that up to 5% metastable solubility of Ni in the sublattice of 5 has been reported at 22 C. 15 At the moment all the phase equilibria in the ternary --Ni system are not unambiguously determined. One reason for the uncertainty concerning the --Ni phase diagram is that in the low- region of the system the kinetics is very slow at the temperatures relevant to soldering. Thus, even after extensive solid-state annealing, it is not clear whether the equilibria have been established. Therefore, the phase equilibria are presented with dashed lines in the low- Ni Mole Fraction of Ni Γ Ni 3.2 Ni Mole Fraction of Ni 5 3 Fig. 9. Isothermal section at 125 C from the assessed --Ni ternary phase diagram with superimposed diffusion path. region. On the basis of the diagram shown in Fig. 9, the addition of more than 5 at.% of Ni to will lead to the local equilibrium either between (Ni) and (Ni,) 3 or between (Ni) and Ni 5 2, contrary to the diagram by Lin et al. 25 So, the diffusion path presented with the dotted line in Fig. 9 can go ether through the (Ni,) 3 or ternary phase region to (Ni,) 3 2 where from it goes via (,Ni) 5 to b-. Hence, in these reaction sequences the 3 is not stable and so cannot form. In addition, because the growth rates of Ni intermetallics are known to be very low they can be present but as very thin layers that they are outside the resolution limits of the FE-SEM. Explaining the effect of dissolved Ni on the diffusion kinetics in 3 (and 5 ), Garner et al. used another approach based on phase diagram information on the --Ni ternary system. 29 Their approach was based on a rule of thumb originally presented by Birchenall. 3 According to this rule, the interdiffusion coefficient D ~ in a solution phase is decreased, if the addition of component i raises its melting point (i.e., liquidus temperature), and vice versa. On the basis of their interpretation of the 3 -Ni 3 isopleth, 31 Garner et al. concluded that the addition of Ni slows the 3 growth kinetics. This conclusion is not easy to accept. Firstly, the rule is based on the experimental results obtained with binary solution phases and not with ternary ordered compounds (i.e., 3 ). Secondly, in the case of the 3 -Ni 3 isopleth, 31 Ni does not increase the melting point of 3, as the authors claim, 29 but of another high-temperature phase ( 4, namely the c-phase). In fact, addition of Ni to 3 decreases the solvus curve between the e and c phases. As to the rationalization of the results related to the accelerated growth of (,Ni) 5, and the formation of voids at the / 3 interface we need to supplement the thermodynamics of the system with kinetic considerations. Figure 1 presents the Gibbs free energy diagram of the binary - system at 125 C, including also the effects of dissolved Ni. Gibbs energy curves have been calculated from the same assessed thermodynamic data as in the isothermal section (Fig. 9). A detailed description of the data and the calculation procedure are presented in earlier publication. 32 It is to be noted that in ternary systems the tie lines are not usually in the plane of the vertical sections. 33 Therefore the free energy values must be obtained along the diffusion paths (following the tie lines). As shown in Fig. 1, the driving force for the diffusion of atoms through 5 (from interface I to II) is about 31 /mol (DG, left-hand side of Fig. 1) and that of atoms through 3 (from interface III to II) is about 1, /mol (DG, right-hand side of Fig. 1) in the binary / diffusion couple. When Ni is incorporated into the binary IMCs, the driving forces are altered markedly. The EDS measurements show that the Ni content of the

6 13 uorinen, Laurila, Mattila, Heikinheimo, and Kivilahti Gibbs energy [/mol] G 5 G * -1 G ** Mole-fraction G g β g η (β) 5 (η) 3 (ε) (α) g ε, (,Ni) 5 is close to that of the original (Ni) alloy. Therefore, when 1 at.% of Ni is added to, the Ni content of (,Ni) 5 is about 1 at.% and its stability is increased to g g*. The 3 in equilibrium with it is practically free of Ni, as can be seen from the direction of the tie lines in Fig. 9, and so its stability will remain practically the same as that of pure 3 (g e* ). As a result, the driving forces for the diffusion over the IMCs have changed and are DG (about twice DG ) for and DG (about 9/1 DG ) for atoms. If the Ni content of (,Ni) 5 increases to 5 at.%, the stability of the g phase is increased even more (g g** ) and the (,Ni) 3 in equilibrium with it has also stabilized slightly (g e** ). Thus, the driving force for the diffusion of atoms though ð,niþ 5 ðg Þ has increased up to 4 to 5 times and the driving force for the diffusion of atoms though ð,niþ 3 ðg Þ has decreased to about two-thirds of that in the binary - system. Further, if 3 does not form, as the experimental evidence suggests, the driving force for the diffusion of through (,Ni) 5 will increase even further (dashed tangent line in Fig. 1 for g g* and g a*, which is the Gibbs free energy curve for the fcc(a)-phase containing 5 at.% Ni). The diffusion flux of component i ( i ) is directly proportional to the driving force (chemical potential gradient, dl/dx), according to the well-known Nernst Einstein relation: g η g η G g ε g α G * G G ** Fig. 1. Gibbs free energy diagram at 125 C showing the driving forces for diffusion over the schematically represented interfacial reaction zone. I II i ¼ C i M i i where C = concentration [mol/m 3 ], M (= D * /RT) is mobility and D * = tracer diffusion coefficient [m 2 /s], R = universal gas constant (8.314 /kmol), and T = absolute temperature. Therefore the material flux can be expected to grow in 5 and to decrease in 3 when Ni is brought into the - system. By extrapolating the data reported by Oh 15 (at temperatures between 18 C and 22 C) down to 125 C, the following ratios for the intrinsic diffusion coefficients are obtained: D 3 D 3 3 and D 5 D 3 3:5: ð2þ Paul 14 determined the ratio of diffusion fluxes at 22 C as 5! 5 ¼ D 5 5 D 5 ¼ 1: 5 ð3þ 3 and 3 ¼ :9: where D / i = intrinsic diffusion coefficient of component i in phase / and / i = molar volume of component i in phase /. On the basis of these, the following estimation is made for the relative diffusion fluxes in the binary - system at 125 C : 5 ¼ 1ð¼ reference valueþ, 5 ¼ :; 3 ¼ :3 and 3 ¼ :9. By assuming that the changes in the fluxes of the fast-moving species (i.e., in 5 and in 3 ) are proportional to the changes in the driving forces due to dissolved Ni and that the changes in the magnitudes of 5 and 3 can be ignored (because the fluxes are so small in comparison to 5 and 3 ), the following flux values are obtained: 5 ¼ 2; 5 ¼ :; 3 ¼ :3 and 3 ¼ :8, when the Ni content of (, Ni) 5 is 1 at.% and 5 ¼ 4:5; 5 ¼ :, 3 ¼ :3 and 3 ¼ :, when the Ni content of (,Ni) 5 is 5 at.%. Besides, due to Ni addition, the vacancy flux balance ð 5 v and 3 v Þ is therefore also significantly changed as can be seen from Fig. 11. At the moment it is not unambiguously clear what microstructural effects Ni induces in the (, Ni) 5 compound layer. However, what is evident is that the presence of Ni will accelerate the diffusion of through this layer. an Loo et al. 34 have analyzed the effect of intrinsic diffusion coefficients and initial and interfacial concentrations of each phase on the creation and annihilation of vacancies.

7 Solid-State Reactions between (Ni) Alloys and 131 Any interphase boundary must be able to serve as either a source or a sink for vacancies to accommodate the differences in vacancy fluxes. In our case, the 5 / 3 interface must be a source of vacancies for both IMC layers. This is because is the main diffusing species in 5 and is the main diffusing species in 3. In order to maintain the vacancy flux balance, the vacancies being created at the 5 / 3 interface must be annihilated at other interfaces or in the bulk phases. When the flux through 5 increases the 5 / 3 interface must produce more vacancies to keep up with the incoming atoms. Since the 3 / 5 interface provides vacancies for both IMC phases, more vacancies are pumped into the 3 layer as well. However, based on the presented analysis and experimental results the diffusion of is not markedly accelerated in 3 (so vacancies are not consumed) and therefore vacancies diffuse to the / 3 interface, where they will be annihilated or, if in excess (as in this case), precipitated as macroscopic pores (i.e., Kirkendall voids). Furthermore, if we assume that the ratio of the mobilities of and in (,Ni) 3 does not markedly change as a result of the addition of Ni, the Kirkendall effect should be enhanced, as the material flux through (,Ni) 3 is reduced. When more Ni is added to, the fluxes in (,Ni) 3 will be reduced even further, finally resulting in the disappearance of (,Ni) 3 as well as the voids. Thus, the changes in the diffusion fluxes resulting from the addition of Ni can result in both the observed changes in the thickness of the IMC layers and the formation and disappearance of voids (Fig. 11). Pure/ 1.Ni/ 5.Ni/ (α) I I 3 (ε) (Ni)(α) (,Ni) 3 I (Ni)(α) (,Ni) II II T=125 C II (,Ni) 5 (,Ni) 5 5 (η) (,Ni) 5 (,Ni) 5 (,Ni) 5 (,Ni) 5 (,Ni) 5 (,Ni) III III III (β) (β) (β) CONCLUSION The interfacial reactions between and (Ni) alloys have been investigated as a function of Ni content at the temperature of 125 C. The effects of Ni on the interfacial intermetallic growth kinetics as well as on the formation of voids has been studied based on the combined thermodynamic, kinetic, and microstructural approach. The results show that the addition of Ni to has significant effects on the total thickness of the IMC layers, their relative growth kinetics, and the formation of voids at the / 3 interface. The thickness of the 5 layer was observed to grow as a function of increasing Ni content up to 5. at.%, after which it decreased. On the other hand, the thickness of 3 decreased as a function of increased Ni content. The number of voids in the 3 layer increased as the layer became thinner. The disappearance of 3 with the higher Ni content could not be explained solely on the basis of phase stability information. However, based on the assessed thermodynamic data it was shown that the driving forces for the diffusion of the component atoms as well as the mobilities of the components are altered in such a manner that the growth of Fig. 11. The effect of Ni on the intrinsic fluxes through the IMCs. Note that the thicknesses of the IMC layers are not to scale, but the flux vectors are. 5 is favored over 3, when Ni has dissolved into the intermetallics. Furthermore, the changes in the intrinsic diffusion fluxes through the interfacial reaction zone are proposed to result in the void formation. ACKNOWLEDGEMENTS The authors greatly acknowledge Dr. Hao Yu for the assessment of the --Ni system. This work was supported financially by the Academy of Finland. REFERENCES 1. W. Bader, Weld.. 48(12), 551 (199). 2. W. Bader, Proc. Conf. Physical Metall., Metal oining, St. Louis, MO, Oct. 1 17, (198), TMS/AIME, Warredale, USA. 3. Z. Mei, A. Eslambochi, and P. ohnsson, Proc. 48th Electronic Components and Tech. Conf., IEEE, Piscataway, N, 952, (1998). 4. A. Minor and. Morris, Metall. Mater. Trans. 31, 798 (2). 5. T. Laurila,. uorinen, T. Mattila, and.k. Kivilahti,. Electron. Mater. 34(1), 13 (25).

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