Figure 2 shows the EBSD orientation image map (OIM) of each Cu electroplated film in the normal direction (ND). The color for each grain orientation in the Cu films is represented by an inverse pole figure (IPF) of Cu in Fig. 2c. For example, green in the IPF represents a grain with a high {110} texture. Figure 2a shows the OIM and the IPF with the distribution intensity for each orientation (IPF-DI) in the Cu A film. The Cu A film was clearly {110}-preferred, and that was supported by the high intensity of {110} in its IPF-DI. The ratio of the {110} grains in the OIM was 78.1% by tolerance of 15° between the crystal directions and the axis of ND. The Cu B and Cu C films also possessed {110}-preferred orientations in ND, as shown in their OIMs and IPFs-DI (Figs. 2b and 2c). The ratios of {110} grains in the Cu B and Cu C films were 73.2% and 81.7%, respectively. The ratio of {110} in each Cu film was summarized in Table 1. Hence, the grain orientations of all Cu films in ND can all be regarded as highly < 110>-oriented, as shown in their OIMs and IPFs-DI.
Table 1
Ratio of the {110}-preferred grain orientation in each Cu film under a tolerance of 15°.
Cu electroplated film | A | B | C |
{110} ratio | 78.1 % | 73.2 % | 87.1 % |
The grain orientation at a metallic surface plays a significant effect on the surface diffusion. For example, the surface diffusivity is the fastest on the {111} planes in a face-centered cubic crystal because they are the close-packed planes [9, 10]. Based on the theoretical calculation, the diffusivities of Cu on the {111}, {100}, and {110} planes are 9.42⋅10− 6, 1.19⋅10− 6, and 5.98⋅10− 6 cm2/s, respectively [11]. On the other hand, “Kirkendall effect” is a theory to explain the vacancy diffusion through an interface between two dissimilar solid materials, and “Kirkendall voids” are formed at one side material with a higher diffusivity. This is the reason that void formation frequently appeared at the interface between Sn-based solder and Cu because the diffusion of Cu into Sn is much faster than that of Sn into Cu [3, 6, 12]. Consequently, the control of the Cu surface orientation is very important to investigate the formation of Kirkendall voids at the SAC305/Cu interface. Because the three Cu electroplated films were all {110}-preferred on their surface, the effect of surface diffusivity on the void formation can be ignored in this study. In other words, the effects of microstructure and impurity concentration on the void formation can be sure compared.
Figure 3 shows the cross-sectional FIB images of the Cu A, B, and C films. The grains in Cu A were embedded with a few twin boundaries (Fig. 3a). In Cu B and C, their grains were columnar-shaped as shown in Figs. 3b and c. Therefore, the microstructure of Cu B was similar to that of Cu C, but they were dissimilar to that of Cu A without the columnar grains. Some articles emphasized that the microstructures of the Cu films significantly induced or suppressed the Kirkendall effect [7, 8, 12]. Among them, the smaller grain size in a Cu film provided numerous grain boundaries for Cu fast diffusion into the Sn-rich solder [7, 13], and the grain boundaries between columnar grains were the paths for vacancy aggregation on the Cu surface [8]. They enhanced the Kirkendall effect at the Sn-rich solder/Cu interface unless Cu nanotwin existed in the Cu films [8, 12]. According to the above-mentioned inference, although the Cu A film possessed twin boundary, but the number of twins was much lower than that previously reported capable of suppressing the void formation. The Cu B and C films were with silimar columnar microstructure that enhanced the formation of Kirkendall voids. Therefore, from the microstructural point-of-view the three Cu films all faced high risks of the void formation.
Figure 4 depicts the interfacial reactions of the SAC305 solder with the Cu A, B, and C films during a thermal aging process at 200 °C for 1000 h. Basically, three intermetallic compounds (IMCs) were observed at the SAC305/Cu interface. The layer-type IMCs were identified as the Cu6Sn5 and Cu3Sn phase using EDX which were formed as a result of the interfacial reactions between Sn-rich solder and Cu. The Ag3Sn particles embedded at the SAC305/Cu6Sn5 interface were precipitates from the SAC305 solder matrix. At the SAC305/Cu A interface (Fig. 4a), a void-free interlayer was observed after the thermal aging, which was similar to that at the SAC305/Cu B interface (Fig. 4b), although the microstructural difference between Cu A and B was significant (Fig. 3a versus 3b). The Cu B film with columnar grains provided several channels for the diffusion of vacancies from the Cu interior toward the film surface. However, voids were not observed at the SAC305/Cu B interface. Conversely, at the SAC305/Cu C interface where the Cu C film possessed a similar microstructure with Cu B, plenty of voids were formed after thermal aging for 1000 h. Kirkendall effect was clearly observed in the SAC305/Cu C system. Moreover, these voids even propagated across the entire SAC305/Cu interface and connected in series to form many crevices (or cracks) parallel to the interface, forming a IMC/crevice alternating structure. Based on the above examinations, it was concluded that the Cu microstructure is not the only essential factor to induce the formation of Kirkendall voids at the interface between SAC305 and Cu.
On the other hand, some impure species originating from the organic substances in the plating solution usually co-deposit in a Cu electroplated film and segregate in the grain boundaries. An increase in the atomic disarrangement due to the impurity incorporation gives rise to plenty of vacancies in the Cu film, and the void formation is induced by the segregation of vacancies during thermal aging. Therefore, the impurities are considered as another possible source of producing the voids in the Cu films [13]. Herein, the incorporation of impurities such as oxygen (O), sulfur (S), and carbon (C) in the Cu A, B, and C films was examined using SIMS as shown in Fig. 5. Obviously, the incorporation levels of O, S, and C were in the same order of Cu C > Cu B ≈ Cu A. In Fig. 4, the SAC305/Cu A and SAC305/Cu B solder joints were void-free. Although the microstructure of Cu B resembled that of Cu C, serious void propagation occurred in the interlayer between SAC305 and Cu C during thermal aging. The results demonstrated that, rather than the microstructures of the Cu films, the impurity level in the Cu films were highly related to the void formation in the SAC305/Cu solder joints thermally aged at 200 °C. The impurities incorporated in the Cu films played the roles to produce plenty of vacancies nucleating at the SAC305/Cu C interface. Segregation of the impure species or their derivatives in the grain boundaries might also annihilate the vacancy sinking sites which accelerated the accumulation of vacancies to form voids [14–16]. Severe void propagation gave rise to the formation of continuous crevices or cracks which was a kind of volumetric defects in crystal caused by void nucleation during annealing [17, 18].
To further investigate the impurity effect, a new Cu electroplated film (termed as Cu C*) was prepared by reducing the additive concentration. The Cu C* film possessed a microstructure similar to that of Cu C but a lower impurity concentration as shown in the SIMS patterns in Fig. 5. Figure 6 shows the cross-sectional SEM image of the SAC305/Cu C* interface after thermal aging at 200 °C for 1000 h. Discrete voids were found but these voids had not propagated and aggregated together to form continuous crevices as those in the SAC305/Cu C system shown in Fig. 4(c) and 4(d). Such different void evolution can be attributed to the level of impurity incorporated in the Cu films. As shown in Fig. 5, the impurity level in the Cu C* film was lower than that of Cu C, so the formation rate of voids caused by the oversaturation of vacancies was lower in the Cu C* film than in Cu C. Therefore, voids were formed discretely instead of continuous crevices in the SAC305/Cu C* system. In other words, the impurity concentration in the Cu electroplated films plays a crucial role to dominate the void evolution in the form of discrete voids or continuous crevices caused by severe void propagation.