- Email: [email protected]
Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints
MR-11967; No of Pages 11 Microelectronics Reliability xxx (2016) xxx–xxx
Contents lists available at ScienceDirect
Microelectronics Reliability journal homepage: www.elsevier.com/locate/mr
Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints Chaoran Yang, Fuliang Le, S.W. Ricky Lee ⁎ Electronic Packaging Laboratory, Center for Advanced Microsystems Packaging, Hong Kong University of Science & Technology, Clear Water Bay, Kowloon, Hong Kong
a r t i c l e
i n f o
Article history: Received 8 May 2015 Received in revised form 14 March 2016 Accepted 14 March 2016 Available online xxxx Keywords: Cu–Sn IMCs Brittle failure Fractographic analysis Charpy impact test Cu3Sn-controlling thermal degradation mechanism
a b s t r a c t A detailed experimental study on the fracture mechanism of Cu–Sn intermetallic compounds (IMCs) in the Pb-free solder was presented in this paper. The growth behaviors of the Cu6Sn5 and Cu3Sn IMCs were inspected and the respective evolution pattern of their microstructures was investigated. Then, a detailed fractographic analysis on brittle fractured solder joints was conducted after the high speed ball pull test. The fracture locations in the Cu–Sn IMC layers during different periods of aging process were identified. The fracture modes of Cu6Sn5 and Cu3Sn were determined as well. Afterwards, the fracture energies of different Cu–Sn IMC materials were directly compared using the Charpy impact test with a specially designed specimen. It was found that the grain boundary of Cu3Sn is the weakest link in the Cu–Sn IMC system. Finally, based on these three parts of study, a mechanism to explain the thermal degradation of Cu–Sn IMCs was proposed. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction The intermetallic compound (IMC) layer plays a significant key role in the solder joint of microelectronics. The presence of a continuous IMC layer is a must to form a physical bonding between the solder and the substrate metal [1,2]. Yet on the other hand, the IMC material is brittle in nature. It tends to fracture when the solder joint is subjected to a high strain-rate stress condition, causing a brittle failure of the solder joint [3–7]. With the explosive popularity of portable electronics and the introduction of more rigid lead-free solder alloys since the beginning of this century, this issue has become an even greater concern. As a result, improving the mechanical strength of the IMC layer instead of losing it becomes a natural demand. However, when the solder joint is in an elevated temperature environment, the IMC strength will gradually be reduced, and the solder joint is more vulnerable to generate brittle failure [8,9]. This phenomenon is widely known as the thermal aging effect. Among numerous IMC systems in microelectronics, Cu–Sn IMCs, consisting of a layer of Cu6Sn5 (η-phase) and a layer of Cu3Sn (ε-phase), suffer most from this aging issue. Due to the reason that the Cu–Sn IMCs are directly contacted with the Cu substrate without any barrier layer in between, the growth rate of the Cu–Sn IMCs is higher than the other IMC systems, such as Ni–Sn IMCs, and the deterioration of the IMC strength during the aging process is also more severe.
⁎ Corresponding author. E-mail address: [email protected] (S.W.R. Lee).
Even though that the thermal degradation of Cu–Sn IMCs brings a great concern about the long-term reliability of the solder joint, the mechanism behind is still not clear. Peoples just simply attribute the strength reduction to the excessive growth of the IMC layer, for it is the most salient difference during the aging period [10,11]. Most of the research reports do support the statement that “if the IMC layer is thicker, then its strength would be lower” [8,9], but obviously this is only a phenomenological description without considering the root cause. Even this statement should also be carefully applied. Our previous study showed that when the thickness of the Cu–Sn IMC layers was affected by the element doping inside the solder matrix, this “thickness-strength” relation would not be valid any more [12]. Zeng observed [13] a voiding phenomenon inside the Cu3Sn layer when the solder joint was subjected to thermal aging. The void density increased with age. He proposed that this void formation was due to the different diffusion rates of Cu and Sn atom, also known as the Kirkendall effect, causing the deterioration of the solder joint. Xu and Pang [14] and Mei [15] have reported similar results. However, they provided no evidence that the fracture was either initiated at or propagated along those defects. Some recent studies have shown that voids are actually caused by impurities in the electroplated Cu layer instead of by the Kirkendall effect [16–18]. The void density strongly depends on the quality of the Cu foil, and it can drastically be reduced if the Cu foil has an optimized annealing process after electroplating. Therefore, the void formation should not be the intrinsic reason of the thermal degradation of Cu–Sn IMC layers. The present research work is aimed to unveil the thermal degradation mystery of the Cu–Sn IMCs, by starting from the very basic studies:
http://dx.doi.org/10.1016/j.microrel.2016.03.021 0026-2714/© 2016 Elsevier Ltd. All rights reserved.
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
2
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
the microstructure characteristics and the fracture morphology of the Cu–Sn IMC layers during different periods of the aging process. The growth behavior, and more importantly, the microstructure evolution of the Cu–Sn IMC layers during the aging process were inspected at first. High magnification observation technique, including scanning electron microscope (SEM) and transmission electron microscope (TEM), was employed. Then, the brittle failures of solder joints were generated using the high speed ball pull test and the fractographic analysis was conducted on both sides of the fracture surface to determine the fracture location. A quantitative evaluation was performed as well using a specially designed Charpy impact specimen so that the fracture energy of different Cu–Sn IMC materials can be directly compared. Based on these three parts of study, a mechanism to explain the thermal degradation of Cu–Sn IMC layers is proposed at the end of this paper. 2. Experimental procedures 2.1. SEM/TEM sample preparation The first step to understand the degradation mechanism of Cu–Sn IMCs is to investigate their microstructure evolution during the aging process. Due to the different grain sizes of Cu6Sn5 and Cu3Sn, a Scanning Electron Microscope (SEM) technique was used to investigate the growth behavior of the Cu6Sn5 layer, while the inspection of Cu3Sn grains was conducted with the help of a Transmission Electron Microscope (TEM). Accordingly, different sample preparation methods were adopted for these two microscopic observation techniques. The samples for SEM observation were prepared through the conventional reflow soldering process. In the current investigation, Sn3.0 wt%Ag0.5 wt%Cu (SAC305) lead-free solder ball was selected for the sample preparation. The solder balls were attached to a conventional PCB through reflow process soldering. The pad finish of the PCB was Organic Solderability Preservative (OSP) to make sure that the Cu–Sn IMCs were formed during the chemical reaction between the solder and base metal. An industrial-level eight-zone reflow oven was used, and the reflow temperature profile is plotted in Fig. 1. The peak temperature was around 243 °C. After the solder ball attachment, the samples, except for those reserved for the as-reflow condition inspection, were put into a high temperature oven for the isothermal aging test. The aging temperature was maintained at 125 °C for a total of 1000 h. The samples under the as-reflow condition and after aging for 600 h and 1000 h were used for IMC observation and characterization. Some of the samples were molded using transparent epoxy for the subsequent cross-section preparation. Conventional mechanical grinding and polishing techniques were employed. Afterwards, the cross-sectional area was slightly etched with a 2%HCl–98%C2H5OH solution for a few seconds to expose the IMC morphology. To fabricate the TEM specimen, a different approach was used. A polycrystalline copper plate was cut to a size of 10 × 10 × 1.5 mm3. The copper surfaces were ground and carefully polished with 0.5 μm
Fig. 1. Reflow temperature profile.
Fig. 2. Setup of the Charpy impact test.
diameter alumina soliquoid and then rinsed in acetone, methanol and distilled water in an ultrasonic bath. Afterwards, two copper plates were immediately covered by a commercial Sn0.7Cu solder paste and clamped together. Instead of the reflow oven, a hot plate was used to melt the solder and the surface temperature of the hot plate was maintained at around 243 °C. The sample was put on the hot plate, and after the solder started melting, the sample was kept on the hot plate for about 4 ~ 5 s before being removed and air cooled to room temperature. Thus, the TEM specimen was actually a Cu–Sn–Cu sandwich-like structure, and the same aging condition was applied to promote the growth of the IMC layer. Before doing the cross-section, the TEM specimen was first cut into a thin slice with a thickness of around 1 mm. This slice was temporarily attached to a TEM holder using an instant wax so that it could be thinned to less than 20 μm using mechanical grinding and polishing techniques. After that, a very thin Cu–Sn–Cu specimen slice
Fig. 3. Modified Charpy impact test sample.
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
Fig. 4. IMC growth in the modified Charpy impact test samples (a) After Aging for 250 h; (b) After Aging for 500 h.
Fig. 5. Effect of thermal aging on growth of Cu6Sn5 layer (a) 0 h aging; (b) 600 h aging; (c) 1000 h aging.
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
3
4
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
also investigated along with the mechanical test, to identify the facture location during different aging period. Therefore, the finished PCB coupons and the separated solder balls were carefully collected during the test, and they were put in a Scanning Electron Microscope (SEM) for fractographic observation. 2.3. Modified Charpy impact test
Fig. 6. Microstructure of the Cu3Sn layer at as-reflow condition (before thermal aging).
was transferred to an ion milling system (Gatan 600 CTMP) for the final thinning. The milling parameter was 5.0 keV, 5 μA and low milling angle at 5°. TEM JEM 2010 from JOEL were used in this study. In order to capture the microstructure evolution of the Cu3Sn grains at the early stage of thermal aging, samples after 100 h and 200 h of aging were used to get the TEM image.
Although the ball pull test result can provide an overall assessment of the mechanical integrity of the IMC layers, it cannot represent the fracture toughness of a single IMC material, due to their very complex microstructure and fracture behavior. In this study, a Charpy impact test was performed to quantitatively characterize the relevant toughness of the IMC materials separately. A typical setup of the Charpy impact test is shown in Fig. 2. Although the present study followed the ASTM E23 Standard in principle, there was a modification in the test samples. Instead of making a V-notch on a whole piece of Cu specimen, the present samples consisted of two identical half-pieces (with one corner trimmed) which were soldered together as shown in Fig. 3 (all dimensions followed ASTM A370). The solder joint between the two half-pieces was formed by electro-plated Sn with thermal compression at 400 °C. After 10 min of heating, the whole bond line became a layer full of Cu6Sn5 IMC. The layer thickness was about 5 μm. Some of the prepared samples were subject to subsequent thermal aging at 150 °C for up to 500 h. Fig. 4 shows further IMC growth during thermal aging. It was observed that the solder joint between the two half-pieces of Cu became a mixture of Cu3Sn and Cu6Sn5 first and then turned into a full layer of Cu3Sn after 500 h of thermal aging. 3. Results and discussion
2.2. High speed ball pull test 3.1. Growth behavior of Cu–Sn IMCs under thermal aging conditions Joint-level ball pull and ball shear tests are widely used in industry to evaluate the integrity of the solder joint. In this study, high speed solder ball pull tests were performed with a commercial bond tester (DAGE 4000HS), following JEDEC Standard JESD22-B115. The same solder ball composition and pad finish were used and the thermal aging condition (125 °C, 1000 h) was applied as well to accelerate the growth of Cu–Sn IMC layers. The nominal diameter of solder balls was 0.76 mm and the solder pads were solder mask defined (SMD). Based on the results of previous research [6,9], the ball pull speed was set at 500 mm/s, which can guarantee brittle fracture of solder joints at the IMC area and has a good correlation to the failure mode in mechanical drop tests. In addition to the fracture strength data, the fracture surface was
3.1.1. Growth behavior of Cu6Sn5 during thermal aging The morphologies of the Cu6Sn5 layer during different aging periods observed from the top and the cross-section are shown in Fig. 5. For the as-reflow condition, the Cu6Sn5 layer appeared to be a scallop-like morphology, leading to a very rough and non-uniform IMC layer. For the top view observation, it can be found that the grain size of the Cu6Sn5 is around 1 ~ 2 μm. With the increase of the aging hours, the thickness of the IMC layer also increased. The top view observation shows the coarsening of the Cu6Sn5 grains and that the initial scalloped intermetallic layer became planar. This planarization process is due to the shorter diffusion distance between the scallop valleys and the Cu substrate.
Fig. 7. Microstructure of the Cu3Sn layer at the early stage of aging process (a) After 100 h of thermal aging; (b) After 200 h of thermal aging.
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
5
Thus, the Cu will diffuse faster to the scallop valley than to the scallop peak, leading to a faster growth rate at the valley and subsequent planarization of the layer. This process, as explained by Choi and Lee [19] from a thermodynamic point of view, is that the larger surface area of a scallop-like morphology, versus that of a planar morphology with a similar thickness, results in a large driving force for the reduction in surface-free energy of the layer. 3.1.2. Growth behavior of Cu3Sn during thermal aging The TEM image of the Cu3Sn layer at the as-reflow condition is shown in Fig. 6. The Cu3Sn grain size is around 150 nm and demonstrated a columnar structure. One important feature is that only one stack of Cu3Sn grains appeared between the Cu6Sn5 and the Cu substrate. Therefore, compared with the Cu6Sn5 layer, the Cu3Sn layer was relatively uniform. Different from Cu6Sn5, two kinds of growth mechanisms can be observed in Cu3Sn, as shown in Fig. 7. After 100 h of thermal aging, some of the Cu3Sn grains became elongated towards the direction perpendicular to the Cu substrate, as indicated by the yellow dots in Fig. 7(a). In addition, at some of the triple-junction sites along the interface between the Cu3Sn grains and Cu, the new nucleated Cu3Sn grains can be detected, as marked with the white dots. Since these two processes happen simultaneously, after the sample was solid state aged for 200 h, the grain size of the Cu3Sn has become non-uniform. At some sites, the Cu3Sn layer was converted to a multi-stack structure. After 1000 aging hours, as shown in Fig. 8, the Cu3Sn layer becomes a complete multi-stack structure.
Fig. 9. Ball pull force vs. thermal aging time.
Figs. 9 and 10 present the results of high speed solder ball pull tests in terms of ball pull force versus thermal aging time and IMC thickness, respectively. Note that the IMC thickness was evaluated by an area averaging approach based on cross-section pictures similar to those typical ones given in Fig. 2. The trends in Figs. 9 and 10 indicate that the ball pull strength decreases with increasing the aging time and the IMC thickness. To understand the mechanism of these two figures, especially Fig. 10, the corresponding fracture surfaces on top of the copper pads and at the bottom of the solder balls were both inspected to determine the fracture location. The observation results of samples under three thermal aging conditions are presented here: without aging, after aging for 200 h and after aging for 1000 h, and the SEM observation results are demonstrated from Figs. 11–13.
To determine the fracture location, not only the fracture morphology inspection is needed, but also the elemental characterization. The fracture surface of the solder joints presented a very complicated morphology. However, through high magnification SEM observation, it can be found that for the sample without thermal aging, two types of morphology can be distinguished on the fracture surface at the pad side. One is relatively flat, demonstrating a plateau-like structure with the outline of grain boundaries. The other is like bunches of small particles randomly distributed among those plateaus. The EDX result showed that those plateaus were Cu6Sn5 and the small particles were Cu3Sn grains. The fracture surface at the bottom of solder ball also presented two types of morphology that respectively matched very well with the two morphologies at the pad side. One was also a plateau-like morphology with a flat surface and the other was just like bunches of dents embedded on the fracture surface. The element analysis showed that the alloy compositions at the bottom of the solder balls were all Cu6Sn5. It can be determined that for as-reflow solder joint samples, there are two dominant fracture locations: (1) inside the Cu6Sn5 grains, where the fracture surface shows flat morphology at both the pad side and the ball side; (2) along the interface between the Cu6Sn5 and Cu3Sn layer, where the fracture surface shows the exposed small Cu6Sn5 particles at the pad side and the dents embedded at the bottom of the Cu6Sn5 grains.
Fig. 8. Microstructure of the Cu3Sn layer after 1000 h of thermal aging.
Fig. 10. Ball pull Force vs. IMC thickness.
3.2. Thermal aging effect on fracture location after ball pull test
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
6
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
Fig. 11. Fracture Surface of the Solder Ball after Pull Test (without aging) (a) Overview of the Fracture Surface; (b) Close-up View of Fracture Surface of the Solder Ball shown in (a); (c) Close-up View of the Fracture Surface shown in (b).
For the samples after 200 h of thermal aging, it seems that mild thermal aging would not significantly alter the dominant fracture location inside the Cu–Sn IMC layers, but after the solder joint went through 1000 h of thermal aging, a remarkable change in the fracture surface can be observed. Another type of morphology appeared on the fracture surface at both the pad side and solder ball side, and covered most of the area. This morphology showed a relatively fine texture due to the small grain size. The EDX spectrum confirms that the composition of those fine textures was Cu3Sn. This result suggests that when solder joints endure a long period of thermal aging, the dominant fracture location shifts to the Cu3Sn layer. Yet the other two kinds of aforementioned morphology (the flat plateau and the small particles/dents) still can be observed at the fracture surface, meaning that the fracture location is either inside the Cu6Sn5 or along the Cu6Sn5/Cu3Sn interface at those specified areas. A close-up observation of the fractured Cu6Sn5 grains and Cu3Sn grains shows that they have completely different fracture mechanisms. It was found that the fracture of the Cu6Sn5 occurred by transgranular cleavage and the river line between different cleavage planes is indicated
in Fig. 14. The fracture of the Cu3Sn layer, as shown in Fig. 15, appeared an intergranular fracture mode as the crack path follows the Cu3Sn grain boundaries. Further microscopic inspections were performed on the fractured pad surfaces cross-sectioned by Focused Ion Beam (FIB) after the high speed ball pull tests. Two benefits can be obtained by utilizing FIB technology. Before etching (this is how FIB ‘cuts’ the sample), the FIB will deposit an extra thin layer on the surface of the target area at first. Usually this thin layer of deposition is tungsten or platinum, and it protects the material underneath it from etching damage. This ensures that no information about the fracture surface will be lost due to the crosssectioning. Moreover, with a dual beam system, the FIB is also equipped with a SEM imaging unit. Therefore, the cross-sectioning can be very site-specific. Cross-section observations of the fracture surfaces (pad side) during different aging periods are shown in Figs. 16–18. The top view images from the secondary electron field are provided as well to indicate the cross-sectioning location, which is marked by the red dash line. In order to be aligned with previous fractographic analysis, samples
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
7
Fig. 12. Fracture Surface of the Solder Ball after Pull Test (200 h aging) (a) Overview of the Fracture Surface; (b) Close-up View of Fracture Surface of the Solder Ball shown in (a); (c) Closeup View of the Fracture Surface shown in (b).
without thermal aging, after aging for 200 h and for 1000 h were selected. From the cross-section observation, the layered structure below the fracture surface can be displayed. Taking the as-reflow condition sample as the example, as shown in Fig. 16, the cross-sectioning area intentionally runs through from the flat area (left-hand side of the SE image) to the rough area (right-hand side of the SE image). In the secondary ion field, it was recognized that the flat area was the bottom part of a Cu6Sn5 grain, underneath which there was a thin layer of Cu3Sn grains. In the right part of the figure, this Cu3Sn layer became exposed at the top of fracture surface. This result further confirms that those small particles on top of the fracture surface are indeed Cu3Sn grains. The 200-h aging sample displayed a very similar fracture surface morphology to the sample without thermal aging, yet a much thicker Cu3Sn layer was inspected in the cross-section observation. While for the sample aged for 1000 h, the Cu6Sn5 disappear and only Cu3Sn remained on the pad side. These results provide a strong indication that a change in the IMC related failure mechanism exists when the SAC solder joint is subject to prolonged thermal aging. Most likely, such a variation also leads to a reduction in the solder joint strength. In other words, the decrease of IMC strength is not simply
a result of the increase in IMC thickness, but due to the shift of fracture location in different IMCs. 3.3. Fracture energy comparison of Cu–Sn IMCs The Charpy impact test result is presented in Fig. 19. Compared with the result at time zero, the samples after 250 h of thermal aging exhibited a slight decrease in the fracture energy. After 500 h of thermal aging, the fracture energy dropped substantially. Note that the failure modes of all samples in the Charpy impact tests were brittle fracture. However, there existed various possibilities of fracture locations. Therefore, more detailed fractographic inspections were required. After the Charpy impact test, similar to the previous ball pull test, the observation on both sides of the fractured impact test sample was carried out. During the inspection, multiple sites were observed to make sure the dominant fracture location for each aging condition, but only the typical fracture morphology is presented in this paper, shown from Figs. 20 to 22. It was found that at time zero, both sides of the fracture surface were Cu6Sn5, while after 500 h of thermal aging, both sides
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
8
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
Fig. 13. Fracture surface of the solder ball after pull test (1000 h aging) (a) Overview of the fracture surface; (b) Close-up view of fracture surface of the solder ball shown in (a); (c) Closeup view of the fracture surface shown in (b).
Fig. 14. Transgranular fracture of Cu6Sn5 grains.
Fig. 15. Intergranular fracture of Cu3Sn grains.
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
9
Fig. 16. Cross-section observation of the fracture surface (without thermal aging) (a) FIB cutting location in secondary electron field; (b) Cross-section view in secondary ion field. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 17. Cross-section View of the Fracture Surface (200 h aging) (a) FIB Cutting Location in Secondary Electron Field; (b) Cross-section View in Secondary Ion Field. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
of the fracture surface were Cu3Sn. It means that for these two cases, the fractures are located inside the Cu6Sn5 layer and the Cu3Sn layer, respectively. For the sample after aging for 250 h, as displayed in Fig. 21, the presence of Cu3Sn on one side of the fracture surface can be confirmed by both the nodular morphology with very small grain size and the EDX spectrum. While on the opposing side, many dents appear at the bottom of the Cu6Sn5 grains, which perfectly correspond to the Cu3Sn
nodules. This indicates that the fracture location is at the interface between the Cu6Sn5 and Cu3Sn. Since the impact test sample under each aging condition demonstrated a unique dominant fracture location, the variation of the fracture energy not only represents the thermal aging effect on the fracture energy of the entire Cu–Sn IMC system, but also more importantly, is the direct measurement of the fracture resistance capability of the
Fig. 18. Cross-section Observation of the Fracture Surface (1000 h aging) (a) FIB Cutting Location in Secondary Electron Field; (b) Cross-section View in Secondary Ion Field. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
10
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
fracture of the solder joint not occur inside the Cu3Sn layer in the first place, but only after a certain period of thermal aging? This question could be answered by reviewing on the microstructure evolution of Cu3Sn layer during the aging process. As the Cu3Sn layer exhibits an intergranular fracture mode, the fracture actually happens at the Cu3Sn grain boundary. However, after the soldering process, only a singlestack of Cu3Sn grains is formed between the Cu6Sn5 layer and the Cu substrate. In this case, when the solder joint is subjected to external loading, the fracture has no choice but to be generated either in the Cu6Sn5 layer or along the interface between these two layers. Under the thermal aging conditions, new Cu3Sn grains are nucleated and developed. As a result, the Cu3Sn layer gradually becomes a multi-stack structure and more Cu3Sn grain boundaries are exposed. When the Cu3Sn grain boundaries become a connected path, the fracture will take place in the Cu3Sn layer and the strength of the entire IMC system is consequently reduced. This process is schematically depicted in Fig. 23. 4. Conclusions Fig. 19. Comparison of impact fracture energy among different thermal aging cases.
three specific locations. In other words, the Charpy impact test result convincingly shows that the Cu6Sn5 and the Cu6Sn5/Cu3Sn interfaces have similar fracture energy, and Cu3Sn is the weakest among the three. 3.4. Cu3Sn-controlling thermal degradation mechanism Since Cu3Sn layer is the weakest link in the Cu–Sn IMC system, then this question would be raised very simultaneously: why does the brittle
The growth of the Cu3Sn layer is driven by two different mechanisms, the nucleation of new Cu3Sn grains at the triple-junction sites near the Cu3Sn/Cu interface, and the oriented elongation of the existing Cu3Sn grains. As a result, the Cu3Sn layer is gradually converted from a single-stack structure at the initial condition to a multi-stack structure after a prolonged aging process. After the high speed ball pull tests, three fracture modes were identified on the fracture surface: the transgranular cleavage of the Cu6Sn5 grains, the interfacial fracture between the Cu6Sn5 and Cu3Sn layers and the Cu3Sn intergranular fracture. At the initial stage of the aging period, including the as-reflow condition, only the first two types of
Fig. 20. Fractographic inspections on the fracture surfaces of a modified Charpy test sample before thermal aging (a) One side of fracture surface; (b) The matching side of fracture surface.
Fig. 21. Fractographic inspections on the fracture surfaces of a modified Charpy test sample subjected to 250 h of thermal aging at 150 °C (a) One side of fracture surface; (b) The matching side of fracture surface.
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
11
Fig. 22. Fractographic inspections on the fracture surfaces of a modified Charpy test sample subjected to 500 h of thermal aging at 150 °C (a) One side of fracture surface; (b) The matching side of fracture surface.
Fig. 23. Schematic demonstration of the Cu3Sn-controlling thermal degradation mechanism.
fracture mode can be observed on the fracture surface. Along with the thermal aging process, in addition to the reduction of the pull force, the fracture location in the Cu–Sn IMC layer gradually shifted from the Cu6Sn5 layer and the Cu6Sn5/Cu3Sn interface to inside the Cu3Sn layer. When the solder joint experiences a long thermal aging period (1000 h in our study), the intergranular fracture of the Cu3Sn layer becomes dominant. The Charpy impact test result quantitatively proved that the Cu3Sn layer had the lowest fracture energy. Based on all the investigations included in this study, a Cu3Sn-controlling thermal degradation mechanism was proposed, that behind the thickness values, the true factor that determines the strength of the Cu–Sn IMC is the microstructure of the Cu3Sn layer. References [1] T. Laurila, V. Vuorinen, J. Kivilahti, Interfacial reactions between lead-free solders and common base materials, Mater. Sci. Eng. R. Rep. 49 (2005) 1–60. [2] C. Ho, S. Yang, C. Kao, Interfacial reaction issues for lead-free electronic solders, Lead-Free Electronic Solders, Springer 2007, pp. 155–174. [3] R. Darveaux, C. Reichman, Ductile-to-brittle transition strain rate, 8th Electronics Packaging Technology Conference, Singapore 2006, pp. 283–289. [4] S.-S. Ha, J.-K. Jang, S.-O. Ha, J.-W. Kim, J.-W. Yoon, B.-W. Kim, et al., Mechanical property evaluation of Sn–3.0Ag–0.5Cu BGA solder joints using high-speed ball shear test, J. Electron. Mater. 38 (2009) 2489–2495. [5] K. Newman, BGA brittle fracture—alternative solder joint integrity test methods, Proceedings. 55th Electronic Components and Technology Conference, Lake Buena Vista, Fl, USA 2005, pp. 1194–1201. [6] F. Song, S.R. Lee, K. Newman, B. Sykes, S. Clark, Brittle failure mechanism of SnAgCu and SnPb solder balls during high speed ball shear and cold ball pull tests, Proceedings. 57th Electronic Components and Technology Conference 2007, pp. 364–372.
[7] Y. Su, L. Tan, T. Tee, V. Tan, Rate-dependent properties of Sn–Ag–Cu based lead-free solder joints for WLCSP, Microelectron. Reliab. 50 (2010) 564–576. [8] T.-C. Chiu, K. Zeng, R. Stierman, D. Edwards, K. Ano, Effect of thermal aging on board level drop reliability for Pb-free BGA packages, 54th Electronic Components and Technology Conference, Las Vegas, NV, USA 2004, pp. 1256–1262. [9] F. Song, S.R. Lee, K. Newman, H. Reynolds, S. Clark, B. Sykes, Effect of thermal aging on high speed ball shear and pull tests of SnAgCu lead-free solder balls, 9th Electronics Packaging Technology Conference, Singapore 2007, pp. 463–470. [10] G. Zeng, et al., A review on the interfacial intermetallic compounds between Sn–Ag– Cu based solders and substrates, J. Mater. Sci. Mater. Electron. 21 (5) (May 2010) 421–440. [11] C.E. Ho, et al., Interfacial reaction issues for lead-free electronic solders, J. Mater. Sci. Mater. Electron. 18 (1–3) (Mar. 2010) 155–174. [12] C. Yang, F. Song, S.W.R. Lee, Impact of Ni concentration on the intermetallic compound formation and brittle fracture strength of Sn–Cu–Ni (SCN) lead-free solder joints, Microelectron. Reliab. 54 (2) (2014) 435–446. [13] K.J. Zeng, et al., Kirkendall void formation in eutectic SnPb solder joints on bare Cu and its effect on joint reliability, J. Appl. Phys. 97 (2005) 024508. [14] L. Xu, J.H.L. Pang, Effect of imtermetallic and Kirkendall voids growth on board level drop reliability for SnAgCu lead-free BGA solder joint, Proc. 56th Electronic Components & Technology Conference, San Diego, CA June 2006, pp. 275–282. [15] Z. Mei, M. Ahmad, M. Hu, G. Ramakrishna, Kirkendall voids at Cu/solder interface and their effects on solder joint reliability, Proc. 55th Electronic Components & Technology Conference, Orlando, FL June 2005, pp. 415–420. [16] J.Y. Kim, J. Yu, Effects of residual impurity in electroplated Cu on the Kirkendall void formation during soldering, Appl. Phys. Lett. 92 (2008). [17] L. Yin, P. Borgesen, On the root cause of kirkendall voiding in Cu3Sn, J. Mater. Res. 26 (2011) 455–466. [18] J. Zhou, et al., Effect of Cu substrate and solder alloy on the formation of Kirkendall voids in the solder joint during thermal aging, Proc. 11th ICEPT-HDP 2011, pp. 944–948. [19] W.K. Choi, H.M. Lee, Effect of soldering and aging time on interfacial microstructure and growth of intermetallic compounds between Sn–3.5Ag solder alloy and Cu substrate, J. Electron. Mater. 29 (2000) 1207–1213.
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
Contents lists available at ScienceDirect
Microelectronics Reliability journal homepage: www.elsevier.com/locate/mr
Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints Chaoran Yang, Fuliang Le, S.W. Ricky Lee ⁎ Electronic Packaging Laboratory, Center for Advanced Microsystems Packaging, Hong Kong University of Science & Technology, Clear Water Bay, Kowloon, Hong Kong
a r t i c l e
i n f o
Article history: Received 8 May 2015 Received in revised form 14 March 2016 Accepted 14 March 2016 Available online xxxx Keywords: Cu–Sn IMCs Brittle failure Fractographic analysis Charpy impact test Cu3Sn-controlling thermal degradation mechanism
a b s t r a c t A detailed experimental study on the fracture mechanism of Cu–Sn intermetallic compounds (IMCs) in the Pb-free solder was presented in this paper. The growth behaviors of the Cu6Sn5 and Cu3Sn IMCs were inspected and the respective evolution pattern of their microstructures was investigated. Then, a detailed fractographic analysis on brittle fractured solder joints was conducted after the high speed ball pull test. The fracture locations in the Cu–Sn IMC layers during different periods of aging process were identified. The fracture modes of Cu6Sn5 and Cu3Sn were determined as well. Afterwards, the fracture energies of different Cu–Sn IMC materials were directly compared using the Charpy impact test with a specially designed specimen. It was found that the grain boundary of Cu3Sn is the weakest link in the Cu–Sn IMC system. Finally, based on these three parts of study, a mechanism to explain the thermal degradation of Cu–Sn IMCs was proposed. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction The intermetallic compound (IMC) layer plays a significant key role in the solder joint of microelectronics. The presence of a continuous IMC layer is a must to form a physical bonding between the solder and the substrate metal [1,2]. Yet on the other hand, the IMC material is brittle in nature. It tends to fracture when the solder joint is subjected to a high strain-rate stress condition, causing a brittle failure of the solder joint [3–7]. With the explosive popularity of portable electronics and the introduction of more rigid lead-free solder alloys since the beginning of this century, this issue has become an even greater concern. As a result, improving the mechanical strength of the IMC layer instead of losing it becomes a natural demand. However, when the solder joint is in an elevated temperature environment, the IMC strength will gradually be reduced, and the solder joint is more vulnerable to generate brittle failure [8,9]. This phenomenon is widely known as the thermal aging effect. Among numerous IMC systems in microelectronics, Cu–Sn IMCs, consisting of a layer of Cu6Sn5 (η-phase) and a layer of Cu3Sn (ε-phase), suffer most from this aging issue. Due to the reason that the Cu–Sn IMCs are directly contacted with the Cu substrate without any barrier layer in between, the growth rate of the Cu–Sn IMCs is higher than the other IMC systems, such as Ni–Sn IMCs, and the deterioration of the IMC strength during the aging process is also more severe.
⁎ Corresponding author. E-mail address: [email protected] (S.W.R. Lee).
Even though that the thermal degradation of Cu–Sn IMCs brings a great concern about the long-term reliability of the solder joint, the mechanism behind is still not clear. Peoples just simply attribute the strength reduction to the excessive growth of the IMC layer, for it is the most salient difference during the aging period [10,11]. Most of the research reports do support the statement that “if the IMC layer is thicker, then its strength would be lower” [8,9], but obviously this is only a phenomenological description without considering the root cause. Even this statement should also be carefully applied. Our previous study showed that when the thickness of the Cu–Sn IMC layers was affected by the element doping inside the solder matrix, this “thickness-strength” relation would not be valid any more [12]. Zeng observed [13] a voiding phenomenon inside the Cu3Sn layer when the solder joint was subjected to thermal aging. The void density increased with age. He proposed that this void formation was due to the different diffusion rates of Cu and Sn atom, also known as the Kirkendall effect, causing the deterioration of the solder joint. Xu and Pang [14] and Mei [15] have reported similar results. However, they provided no evidence that the fracture was either initiated at or propagated along those defects. Some recent studies have shown that voids are actually caused by impurities in the electroplated Cu layer instead of by the Kirkendall effect [16–18]. The void density strongly depends on the quality of the Cu foil, and it can drastically be reduced if the Cu foil has an optimized annealing process after electroplating. Therefore, the void formation should not be the intrinsic reason of the thermal degradation of Cu–Sn IMC layers. The present research work is aimed to unveil the thermal degradation mystery of the Cu–Sn IMCs, by starting from the very basic studies:
http://dx.doi.org/10.1016/j.microrel.2016.03.021 0026-2714/© 2016 Elsevier Ltd. All rights reserved.
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
2
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
the microstructure characteristics and the fracture morphology of the Cu–Sn IMC layers during different periods of the aging process. The growth behavior, and more importantly, the microstructure evolution of the Cu–Sn IMC layers during the aging process were inspected at first. High magnification observation technique, including scanning electron microscope (SEM) and transmission electron microscope (TEM), was employed. Then, the brittle failures of solder joints were generated using the high speed ball pull test and the fractographic analysis was conducted on both sides of the fracture surface to determine the fracture location. A quantitative evaluation was performed as well using a specially designed Charpy impact specimen so that the fracture energy of different Cu–Sn IMC materials can be directly compared. Based on these three parts of study, a mechanism to explain the thermal degradation of Cu–Sn IMC layers is proposed at the end of this paper. 2. Experimental procedures 2.1. SEM/TEM sample preparation The first step to understand the degradation mechanism of Cu–Sn IMCs is to investigate their microstructure evolution during the aging process. Due to the different grain sizes of Cu6Sn5 and Cu3Sn, a Scanning Electron Microscope (SEM) technique was used to investigate the growth behavior of the Cu6Sn5 layer, while the inspection of Cu3Sn grains was conducted with the help of a Transmission Electron Microscope (TEM). Accordingly, different sample preparation methods were adopted for these two microscopic observation techniques. The samples for SEM observation were prepared through the conventional reflow soldering process. In the current investigation, Sn3.0 wt%Ag0.5 wt%Cu (SAC305) lead-free solder ball was selected for the sample preparation. The solder balls were attached to a conventional PCB through reflow process soldering. The pad finish of the PCB was Organic Solderability Preservative (OSP) to make sure that the Cu–Sn IMCs were formed during the chemical reaction between the solder and base metal. An industrial-level eight-zone reflow oven was used, and the reflow temperature profile is plotted in Fig. 1. The peak temperature was around 243 °C. After the solder ball attachment, the samples, except for those reserved for the as-reflow condition inspection, were put into a high temperature oven for the isothermal aging test. The aging temperature was maintained at 125 °C for a total of 1000 h. The samples under the as-reflow condition and after aging for 600 h and 1000 h were used for IMC observation and characterization. Some of the samples were molded using transparent epoxy for the subsequent cross-section preparation. Conventional mechanical grinding and polishing techniques were employed. Afterwards, the cross-sectional area was slightly etched with a 2%HCl–98%C2H5OH solution for a few seconds to expose the IMC morphology. To fabricate the TEM specimen, a different approach was used. A polycrystalline copper plate was cut to a size of 10 × 10 × 1.5 mm3. The copper surfaces were ground and carefully polished with 0.5 μm
Fig. 1. Reflow temperature profile.
Fig. 2. Setup of the Charpy impact test.
diameter alumina soliquoid and then rinsed in acetone, methanol and distilled water in an ultrasonic bath. Afterwards, two copper plates were immediately covered by a commercial Sn0.7Cu solder paste and clamped together. Instead of the reflow oven, a hot plate was used to melt the solder and the surface temperature of the hot plate was maintained at around 243 °C. The sample was put on the hot plate, and after the solder started melting, the sample was kept on the hot plate for about 4 ~ 5 s before being removed and air cooled to room temperature. Thus, the TEM specimen was actually a Cu–Sn–Cu sandwich-like structure, and the same aging condition was applied to promote the growth of the IMC layer. Before doing the cross-section, the TEM specimen was first cut into a thin slice with a thickness of around 1 mm. This slice was temporarily attached to a TEM holder using an instant wax so that it could be thinned to less than 20 μm using mechanical grinding and polishing techniques. After that, a very thin Cu–Sn–Cu specimen slice
Fig. 3. Modified Charpy impact test sample.
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
Fig. 4. IMC growth in the modified Charpy impact test samples (a) After Aging for 250 h; (b) After Aging for 500 h.
Fig. 5. Effect of thermal aging on growth of Cu6Sn5 layer (a) 0 h aging; (b) 600 h aging; (c) 1000 h aging.
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
3
4
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
also investigated along with the mechanical test, to identify the facture location during different aging period. Therefore, the finished PCB coupons and the separated solder balls were carefully collected during the test, and they were put in a Scanning Electron Microscope (SEM) for fractographic observation. 2.3. Modified Charpy impact test
Fig. 6. Microstructure of the Cu3Sn layer at as-reflow condition (before thermal aging).
was transferred to an ion milling system (Gatan 600 CTMP) for the final thinning. The milling parameter was 5.0 keV, 5 μA and low milling angle at 5°. TEM JEM 2010 from JOEL were used in this study. In order to capture the microstructure evolution of the Cu3Sn grains at the early stage of thermal aging, samples after 100 h and 200 h of aging were used to get the TEM image.
Although the ball pull test result can provide an overall assessment of the mechanical integrity of the IMC layers, it cannot represent the fracture toughness of a single IMC material, due to their very complex microstructure and fracture behavior. In this study, a Charpy impact test was performed to quantitatively characterize the relevant toughness of the IMC materials separately. A typical setup of the Charpy impact test is shown in Fig. 2. Although the present study followed the ASTM E23 Standard in principle, there was a modification in the test samples. Instead of making a V-notch on a whole piece of Cu specimen, the present samples consisted of two identical half-pieces (with one corner trimmed) which were soldered together as shown in Fig. 3 (all dimensions followed ASTM A370). The solder joint between the two half-pieces was formed by electro-plated Sn with thermal compression at 400 °C. After 10 min of heating, the whole bond line became a layer full of Cu6Sn5 IMC. The layer thickness was about 5 μm. Some of the prepared samples were subject to subsequent thermal aging at 150 °C for up to 500 h. Fig. 4 shows further IMC growth during thermal aging. It was observed that the solder joint between the two half-pieces of Cu became a mixture of Cu3Sn and Cu6Sn5 first and then turned into a full layer of Cu3Sn after 500 h of thermal aging. 3. Results and discussion
2.2. High speed ball pull test 3.1. Growth behavior of Cu–Sn IMCs under thermal aging conditions Joint-level ball pull and ball shear tests are widely used in industry to evaluate the integrity of the solder joint. In this study, high speed solder ball pull tests were performed with a commercial bond tester (DAGE 4000HS), following JEDEC Standard JESD22-B115. The same solder ball composition and pad finish were used and the thermal aging condition (125 °C, 1000 h) was applied as well to accelerate the growth of Cu–Sn IMC layers. The nominal diameter of solder balls was 0.76 mm and the solder pads were solder mask defined (SMD). Based on the results of previous research [6,9], the ball pull speed was set at 500 mm/s, which can guarantee brittle fracture of solder joints at the IMC area and has a good correlation to the failure mode in mechanical drop tests. In addition to the fracture strength data, the fracture surface was
3.1.1. Growth behavior of Cu6Sn5 during thermal aging The morphologies of the Cu6Sn5 layer during different aging periods observed from the top and the cross-section are shown in Fig. 5. For the as-reflow condition, the Cu6Sn5 layer appeared to be a scallop-like morphology, leading to a very rough and non-uniform IMC layer. For the top view observation, it can be found that the grain size of the Cu6Sn5 is around 1 ~ 2 μm. With the increase of the aging hours, the thickness of the IMC layer also increased. The top view observation shows the coarsening of the Cu6Sn5 grains and that the initial scalloped intermetallic layer became planar. This planarization process is due to the shorter diffusion distance between the scallop valleys and the Cu substrate.
Fig. 7. Microstructure of the Cu3Sn layer at the early stage of aging process (a) After 100 h of thermal aging; (b) After 200 h of thermal aging.
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
5
Thus, the Cu will diffuse faster to the scallop valley than to the scallop peak, leading to a faster growth rate at the valley and subsequent planarization of the layer. This process, as explained by Choi and Lee [19] from a thermodynamic point of view, is that the larger surface area of a scallop-like morphology, versus that of a planar morphology with a similar thickness, results in a large driving force for the reduction in surface-free energy of the layer. 3.1.2. Growth behavior of Cu3Sn during thermal aging The TEM image of the Cu3Sn layer at the as-reflow condition is shown in Fig. 6. The Cu3Sn grain size is around 150 nm and demonstrated a columnar structure. One important feature is that only one stack of Cu3Sn grains appeared between the Cu6Sn5 and the Cu substrate. Therefore, compared with the Cu6Sn5 layer, the Cu3Sn layer was relatively uniform. Different from Cu6Sn5, two kinds of growth mechanisms can be observed in Cu3Sn, as shown in Fig. 7. After 100 h of thermal aging, some of the Cu3Sn grains became elongated towards the direction perpendicular to the Cu substrate, as indicated by the yellow dots in Fig. 7(a). In addition, at some of the triple-junction sites along the interface between the Cu3Sn grains and Cu, the new nucleated Cu3Sn grains can be detected, as marked with the white dots. Since these two processes happen simultaneously, after the sample was solid state aged for 200 h, the grain size of the Cu3Sn has become non-uniform. At some sites, the Cu3Sn layer was converted to a multi-stack structure. After 1000 aging hours, as shown in Fig. 8, the Cu3Sn layer becomes a complete multi-stack structure.
Fig. 9. Ball pull force vs. thermal aging time.
Figs. 9 and 10 present the results of high speed solder ball pull tests in terms of ball pull force versus thermal aging time and IMC thickness, respectively. Note that the IMC thickness was evaluated by an area averaging approach based on cross-section pictures similar to those typical ones given in Fig. 2. The trends in Figs. 9 and 10 indicate that the ball pull strength decreases with increasing the aging time and the IMC thickness. To understand the mechanism of these two figures, especially Fig. 10, the corresponding fracture surfaces on top of the copper pads and at the bottom of the solder balls were both inspected to determine the fracture location. The observation results of samples under three thermal aging conditions are presented here: without aging, after aging for 200 h and after aging for 1000 h, and the SEM observation results are demonstrated from Figs. 11–13.
To determine the fracture location, not only the fracture morphology inspection is needed, but also the elemental characterization. The fracture surface of the solder joints presented a very complicated morphology. However, through high magnification SEM observation, it can be found that for the sample without thermal aging, two types of morphology can be distinguished on the fracture surface at the pad side. One is relatively flat, demonstrating a plateau-like structure with the outline of grain boundaries. The other is like bunches of small particles randomly distributed among those plateaus. The EDX result showed that those plateaus were Cu6Sn5 and the small particles were Cu3Sn grains. The fracture surface at the bottom of solder ball also presented two types of morphology that respectively matched very well with the two morphologies at the pad side. One was also a plateau-like morphology with a flat surface and the other was just like bunches of dents embedded on the fracture surface. The element analysis showed that the alloy compositions at the bottom of the solder balls were all Cu6Sn5. It can be determined that for as-reflow solder joint samples, there are two dominant fracture locations: (1) inside the Cu6Sn5 grains, where the fracture surface shows flat morphology at both the pad side and the ball side; (2) along the interface between the Cu6Sn5 and Cu3Sn layer, where the fracture surface shows the exposed small Cu6Sn5 particles at the pad side and the dents embedded at the bottom of the Cu6Sn5 grains.
Fig. 8. Microstructure of the Cu3Sn layer after 1000 h of thermal aging.
Fig. 10. Ball pull Force vs. IMC thickness.
3.2. Thermal aging effect on fracture location after ball pull test
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
6
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
Fig. 11. Fracture Surface of the Solder Ball after Pull Test (without aging) (a) Overview of the Fracture Surface; (b) Close-up View of Fracture Surface of the Solder Ball shown in (a); (c) Close-up View of the Fracture Surface shown in (b).
For the samples after 200 h of thermal aging, it seems that mild thermal aging would not significantly alter the dominant fracture location inside the Cu–Sn IMC layers, but after the solder joint went through 1000 h of thermal aging, a remarkable change in the fracture surface can be observed. Another type of morphology appeared on the fracture surface at both the pad side and solder ball side, and covered most of the area. This morphology showed a relatively fine texture due to the small grain size. The EDX spectrum confirms that the composition of those fine textures was Cu3Sn. This result suggests that when solder joints endure a long period of thermal aging, the dominant fracture location shifts to the Cu3Sn layer. Yet the other two kinds of aforementioned morphology (the flat plateau and the small particles/dents) still can be observed at the fracture surface, meaning that the fracture location is either inside the Cu6Sn5 or along the Cu6Sn5/Cu3Sn interface at those specified areas. A close-up observation of the fractured Cu6Sn5 grains and Cu3Sn grains shows that they have completely different fracture mechanisms. It was found that the fracture of the Cu6Sn5 occurred by transgranular cleavage and the river line between different cleavage planes is indicated
in Fig. 14. The fracture of the Cu3Sn layer, as shown in Fig. 15, appeared an intergranular fracture mode as the crack path follows the Cu3Sn grain boundaries. Further microscopic inspections were performed on the fractured pad surfaces cross-sectioned by Focused Ion Beam (FIB) after the high speed ball pull tests. Two benefits can be obtained by utilizing FIB technology. Before etching (this is how FIB ‘cuts’ the sample), the FIB will deposit an extra thin layer on the surface of the target area at first. Usually this thin layer of deposition is tungsten or platinum, and it protects the material underneath it from etching damage. This ensures that no information about the fracture surface will be lost due to the crosssectioning. Moreover, with a dual beam system, the FIB is also equipped with a SEM imaging unit. Therefore, the cross-sectioning can be very site-specific. Cross-section observations of the fracture surfaces (pad side) during different aging periods are shown in Figs. 16–18. The top view images from the secondary electron field are provided as well to indicate the cross-sectioning location, which is marked by the red dash line. In order to be aligned with previous fractographic analysis, samples
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
7
Fig. 12. Fracture Surface of the Solder Ball after Pull Test (200 h aging) (a) Overview of the Fracture Surface; (b) Close-up View of Fracture Surface of the Solder Ball shown in (a); (c) Closeup View of the Fracture Surface shown in (b).
without thermal aging, after aging for 200 h and for 1000 h were selected. From the cross-section observation, the layered structure below the fracture surface can be displayed. Taking the as-reflow condition sample as the example, as shown in Fig. 16, the cross-sectioning area intentionally runs through from the flat area (left-hand side of the SE image) to the rough area (right-hand side of the SE image). In the secondary ion field, it was recognized that the flat area was the bottom part of a Cu6Sn5 grain, underneath which there was a thin layer of Cu3Sn grains. In the right part of the figure, this Cu3Sn layer became exposed at the top of fracture surface. This result further confirms that those small particles on top of the fracture surface are indeed Cu3Sn grains. The 200-h aging sample displayed a very similar fracture surface morphology to the sample without thermal aging, yet a much thicker Cu3Sn layer was inspected in the cross-section observation. While for the sample aged for 1000 h, the Cu6Sn5 disappear and only Cu3Sn remained on the pad side. These results provide a strong indication that a change in the IMC related failure mechanism exists when the SAC solder joint is subject to prolonged thermal aging. Most likely, such a variation also leads to a reduction in the solder joint strength. In other words, the decrease of IMC strength is not simply
a result of the increase in IMC thickness, but due to the shift of fracture location in different IMCs. 3.3. Fracture energy comparison of Cu–Sn IMCs The Charpy impact test result is presented in Fig. 19. Compared with the result at time zero, the samples after 250 h of thermal aging exhibited a slight decrease in the fracture energy. After 500 h of thermal aging, the fracture energy dropped substantially. Note that the failure modes of all samples in the Charpy impact tests were brittle fracture. However, there existed various possibilities of fracture locations. Therefore, more detailed fractographic inspections were required. After the Charpy impact test, similar to the previous ball pull test, the observation on both sides of the fractured impact test sample was carried out. During the inspection, multiple sites were observed to make sure the dominant fracture location for each aging condition, but only the typical fracture morphology is presented in this paper, shown from Figs. 20 to 22. It was found that at time zero, both sides of the fracture surface were Cu6Sn5, while after 500 h of thermal aging, both sides
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
8
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
Fig. 13. Fracture surface of the solder ball after pull test (1000 h aging) (a) Overview of the fracture surface; (b) Close-up view of fracture surface of the solder ball shown in (a); (c) Closeup view of the fracture surface shown in (b).
Fig. 14. Transgranular fracture of Cu6Sn5 grains.
Fig. 15. Intergranular fracture of Cu3Sn grains.
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
9
Fig. 16. Cross-section observation of the fracture surface (without thermal aging) (a) FIB cutting location in secondary electron field; (b) Cross-section view in secondary ion field. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 17. Cross-section View of the Fracture Surface (200 h aging) (a) FIB Cutting Location in Secondary Electron Field; (b) Cross-section View in Secondary Ion Field. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
of the fracture surface were Cu3Sn. It means that for these two cases, the fractures are located inside the Cu6Sn5 layer and the Cu3Sn layer, respectively. For the sample after aging for 250 h, as displayed in Fig. 21, the presence of Cu3Sn on one side of the fracture surface can be confirmed by both the nodular morphology with very small grain size and the EDX spectrum. While on the opposing side, many dents appear at the bottom of the Cu6Sn5 grains, which perfectly correspond to the Cu3Sn
nodules. This indicates that the fracture location is at the interface between the Cu6Sn5 and Cu3Sn. Since the impact test sample under each aging condition demonstrated a unique dominant fracture location, the variation of the fracture energy not only represents the thermal aging effect on the fracture energy of the entire Cu–Sn IMC system, but also more importantly, is the direct measurement of the fracture resistance capability of the
Fig. 18. Cross-section Observation of the Fracture Surface (1000 h aging) (a) FIB Cutting Location in Secondary Electron Field; (b) Cross-section View in Secondary Ion Field. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
10
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
fracture of the solder joint not occur inside the Cu3Sn layer in the first place, but only after a certain period of thermal aging? This question could be answered by reviewing on the microstructure evolution of Cu3Sn layer during the aging process. As the Cu3Sn layer exhibits an intergranular fracture mode, the fracture actually happens at the Cu3Sn grain boundary. However, after the soldering process, only a singlestack of Cu3Sn grains is formed between the Cu6Sn5 layer and the Cu substrate. In this case, when the solder joint is subjected to external loading, the fracture has no choice but to be generated either in the Cu6Sn5 layer or along the interface between these two layers. Under the thermal aging conditions, new Cu3Sn grains are nucleated and developed. As a result, the Cu3Sn layer gradually becomes a multi-stack structure and more Cu3Sn grain boundaries are exposed. When the Cu3Sn grain boundaries become a connected path, the fracture will take place in the Cu3Sn layer and the strength of the entire IMC system is consequently reduced. This process is schematically depicted in Fig. 23. 4. Conclusions Fig. 19. Comparison of impact fracture energy among different thermal aging cases.
three specific locations. In other words, the Charpy impact test result convincingly shows that the Cu6Sn5 and the Cu6Sn5/Cu3Sn interfaces have similar fracture energy, and Cu3Sn is the weakest among the three. 3.4. Cu3Sn-controlling thermal degradation mechanism Since Cu3Sn layer is the weakest link in the Cu–Sn IMC system, then this question would be raised very simultaneously: why does the brittle
The growth of the Cu3Sn layer is driven by two different mechanisms, the nucleation of new Cu3Sn grains at the triple-junction sites near the Cu3Sn/Cu interface, and the oriented elongation of the existing Cu3Sn grains. As a result, the Cu3Sn layer is gradually converted from a single-stack structure at the initial condition to a multi-stack structure after a prolonged aging process. After the high speed ball pull tests, three fracture modes were identified on the fracture surface: the transgranular cleavage of the Cu6Sn5 grains, the interfacial fracture between the Cu6Sn5 and Cu3Sn layers and the Cu3Sn intergranular fracture. At the initial stage of the aging period, including the as-reflow condition, only the first two types of
Fig. 20. Fractographic inspections on the fracture surfaces of a modified Charpy test sample before thermal aging (a) One side of fracture surface; (b) The matching side of fracture surface.
Fig. 21. Fractographic inspections on the fracture surfaces of a modified Charpy test sample subjected to 250 h of thermal aging at 150 °C (a) One side of fracture surface; (b) The matching side of fracture surface.
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021
C. Yang et al. / Microelectronics Reliability xxx (2016) xxx–xxx
11
Fig. 22. Fractographic inspections on the fracture surfaces of a modified Charpy test sample subjected to 500 h of thermal aging at 150 °C (a) One side of fracture surface; (b) The matching side of fracture surface.
Fig. 23. Schematic demonstration of the Cu3Sn-controlling thermal degradation mechanism.
fracture mode can be observed on the fracture surface. Along with the thermal aging process, in addition to the reduction of the pull force, the fracture location in the Cu–Sn IMC layer gradually shifted from the Cu6Sn5 layer and the Cu6Sn5/Cu3Sn interface to inside the Cu3Sn layer. When the solder joint experiences a long thermal aging period (1000 h in our study), the intergranular fracture of the Cu3Sn layer becomes dominant. The Charpy impact test result quantitatively proved that the Cu3Sn layer had the lowest fracture energy. Based on all the investigations included in this study, a Cu3Sn-controlling thermal degradation mechanism was proposed, that behind the thickness values, the true factor that determines the strength of the Cu–Sn IMC is the microstructure of the Cu3Sn layer. References [1] T. Laurila, V. Vuorinen, J. Kivilahti, Interfacial reactions between lead-free solders and common base materials, Mater. Sci. Eng. R. Rep. 49 (2005) 1–60. [2] C. Ho, S. Yang, C. Kao, Interfacial reaction issues for lead-free electronic solders, Lead-Free Electronic Solders, Springer 2007, pp. 155–174. [3] R. Darveaux, C. Reichman, Ductile-to-brittle transition strain rate, 8th Electronics Packaging Technology Conference, Singapore 2006, pp. 283–289. [4] S.-S. Ha, J.-K. Jang, S.-O. Ha, J.-W. Kim, J.-W. Yoon, B.-W. Kim, et al., Mechanical property evaluation of Sn–3.0Ag–0.5Cu BGA solder joints using high-speed ball shear test, J. Electron. Mater. 38 (2009) 2489–2495. [5] K. Newman, BGA brittle fracture—alternative solder joint integrity test methods, Proceedings. 55th Electronic Components and Technology Conference, Lake Buena Vista, Fl, USA 2005, pp. 1194–1201. [6] F. Song, S.R. Lee, K. Newman, B. Sykes, S. Clark, Brittle failure mechanism of SnAgCu and SnPb solder balls during high speed ball shear and cold ball pull tests, Proceedings. 57th Electronic Components and Technology Conference 2007, pp. 364–372.
[7] Y. Su, L. Tan, T. Tee, V. Tan, Rate-dependent properties of Sn–Ag–Cu based lead-free solder joints for WLCSP, Microelectron. Reliab. 50 (2010) 564–576. [8] T.-C. Chiu, K. Zeng, R. Stierman, D. Edwards, K. Ano, Effect of thermal aging on board level drop reliability for Pb-free BGA packages, 54th Electronic Components and Technology Conference, Las Vegas, NV, USA 2004, pp. 1256–1262. [9] F. Song, S.R. Lee, K. Newman, H. Reynolds, S. Clark, B. Sykes, Effect of thermal aging on high speed ball shear and pull tests of SnAgCu lead-free solder balls, 9th Electronics Packaging Technology Conference, Singapore 2007, pp. 463–470. [10] G. Zeng, et al., A review on the interfacial intermetallic compounds between Sn–Ag– Cu based solders and substrates, J. Mater. Sci. Mater. Electron. 21 (5) (May 2010) 421–440. [11] C.E. Ho, et al., Interfacial reaction issues for lead-free electronic solders, J. Mater. Sci. Mater. Electron. 18 (1–3) (Mar. 2010) 155–174. [12] C. Yang, F. Song, S.W.R. Lee, Impact of Ni concentration on the intermetallic compound formation and brittle fracture strength of Sn–Cu–Ni (SCN) lead-free solder joints, Microelectron. Reliab. 54 (2) (2014) 435–446. [13] K.J. Zeng, et al., Kirkendall void formation in eutectic SnPb solder joints on bare Cu and its effect on joint reliability, J. Appl. Phys. 97 (2005) 024508. [14] L. Xu, J.H.L. Pang, Effect of imtermetallic and Kirkendall voids growth on board level drop reliability for SnAgCu lead-free BGA solder joint, Proc. 56th Electronic Components & Technology Conference, San Diego, CA June 2006, pp. 275–282. [15] Z. Mei, M. Ahmad, M. Hu, G. Ramakrishna, Kirkendall voids at Cu/solder interface and their effects on solder joint reliability, Proc. 55th Electronic Components & Technology Conference, Orlando, FL June 2005, pp. 415–420. [16] J.Y. Kim, J. Yu, Effects of residual impurity in electroplated Cu on the Kirkendall void formation during soldering, Appl. Phys. Lett. 92 (2008). [17] L. Yin, P. Borgesen, On the root cause of kirkendall voiding in Cu3Sn, J. Mater. Res. 26 (2011) 455–466. [18] J. Zhou, et al., Effect of Cu substrate and solder alloy on the formation of Kirkendall voids in the solder joint during thermal aging, Proc. 11th ICEPT-HDP 2011, pp. 944–948. [19] W.K. Choi, H.M. Lee, Effect of soldering and aging time on interfacial microstructure and growth of intermetallic compounds between Sn–3.5Ag solder alloy and Cu substrate, J. Electron. Mater. 29 (2000) 1207–1213.
Please cite this article as: C. Yang, et al., Experimental investigation of the failure mechanism of Cu–Sn intermetallic compounds in SAC solder joints, Microelectronics Reliability (2016), http://dx.doi.org/10.1016/j.microrel.2016.03.021