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thermal fatigue life of SnAgCu solder joints has been proposed. Damage and failure occurs by recrystallization of the large Sn grains across the high strain ...
Towards a Quantitative Mechanistic Understanding of the Thermal Cycling of SnAgCu Solder Joints D. Schmitz1, S. Shirazi1, L. Wentlent2, S. Hamasha1, L. Yin3, A. Qasaimeh4, and P. Borgesen1 1

Department of Systems Science & Industrial Engineering, Binghamton University, Binghamton, NY 13902 Materials Science & Engineering, Binghamton University, Binghamton, NY 13902 3 GE Global Research, Niskayuna, NY 12309 4 Tennessee Tech University, Cookeville, TN 38505 2

Abstract Microelectronics manufacturers continue to subject a wide range of products or representative test vehicles to accelerated thermal cycling tests. Most such testing is focused on comparisons, whether among alternatives or to an established requirement. However, more often than commonly recognized such comparisons may not reflect the relative performances in service. In fact, most current models have been shown to fail to account for important systematic trends as well as being inconsistent with our current understanding of the failure rate controlling damage mechanism. An alternative mechanistically justified model for the thermal fatigue life of SnAgCu solder joints has been proposed. Damage and failure occurs by recrystallization of the large Sn grains across the high strain region of the joint, followed by crack growth along the resulting network of high angle grain boundaries. The recrystallization was shown to be the damage rate controlling mechanism, except for extremely high strain assemblies and/or harsh cycling conditions, i.e. if we can predict the recrystallization we can predict the number of cycles to failure. So far the model accounts for a variety of important trends and offers guidance as to the interpretation and generalization of accelerated test results. General extrapolations towards service conditions will, however, require the specific functional dependence of the rate of recrystallization on the stress and the precipitate distributions. Another potential difficulty is that constitutive relations are extracted from single sided creep experiments while the dislocation cell structures built up under cyclic loading are certain to be different. Furthermore, the repeated ‘annealing’ during the high temperature dwell affects the hardening. Even if these effects could be ignored the ongoing evolution of the constitutive relations would still effectively prevent the extraction of the above-mentioned functional dependence from comparisons between thermal cycling results and FEM. A special experiment is ongoing in which stresses and strains on the solder joints can be controlled and measured directly, allowing the testing of individual assumptions underlying the proposed model. Preliminary results are presented and compared to results of thermal cycling across different temperature ranges. Introduction Accelerated testing would rarely be meaningful without the assumption of at least a qualitative correlation with performance of the product under realistic service conditions. Such a correlation is however far from always obvious. At the

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extreme, cases have been identified where a mechanistic understanding would suggest that the alternative surviving the longest in testing might not do so in actual use. Figure 1 shows two comparisons between four different SnAgCu alloys. Model BGA assemblies with 20 mil diameter solder balls were cycled between 0oC and 100oC with 10 minute heating and cooling ramps. The solid diamonds reflect the characteristic life, N63, for each alloy in cycling with 10 minute dwells at both temperature extremes. These are very common accelerated test parameters in the microelectronics industry, and taken by themselves the results would suggest that alloy IV (a Zn doped alloy) is the best. However, the open circles show what happens when the dwell times were extended to 30 minutes each. Suddenly, alloy IV is the worst performer. The question would seem to be which test best represents realistic service conditions, but the fact is that either would almost certainly require major extrapolations.

Figure 1: Characteristic number of thermal cycles to failure in 0/100oC cycling with two different dwell times at the temperature extremes [21]. Figure 2 shows two comparisons between four different designs of model CSP assemblies with SAC305 solder joints. The designs differed in terms of pitch, solder volume and pad size. Cycling between -40oC and 125oC suggested that designs C and D (the largest joints and pitches) would be the most reliable, but the strain ranges were clearly high. Cycling instead between 0oC and 80oC design A (the smallest joints and pitch) is clearly superior. It may be tempting to assume that the latter is closer to most realistic service conditions, but if we do not understand the reason for the different performance, we cannot be sure that the trend continues.

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Figure 2: Characteristic number of cycles to failure for four different model CSP designs in -40/125oC cycling (left) and 0/80oC cycling (right). In both cases with 10 oC/min ramp rates and 15 minute dwells [1]. The above examples illustrate at the very least the need for a mechanistic understanding of the effects of cycling parameters on damage evolution and failure. Only this will allow us to predict when accelerated testing may lead to the wrong conclusion. As for a quantitative model, which would allow us to extrapolate accelerated test results to life in service, attempts at this are commonly based on a 2-stage approach: 1) the calculation of stresses, strains and parameters that can be derived from those versus temperature and time 2) the prediction of damage evolution based on this. In the case of SnAgCu solder joints modeling is complicated by the ongoing changes in the constitutive relations during storage [2] and cycling [3]. Dutta et al. have developed microstructurally adaptive constitutive models to account for this [4], but accurate Finite Element Modeling (FEM) based on these would be computationally prohibitive without significant simplifying approximations. A different question is to which extent these or any other constitutive relations allow for simulations of the most relevant parts of each thermal cycle. We shall address this briefly below. As for a damage function, Borgesen et al. relied on an unusually comprehensive set of thermal cycling test results to argue that any of the current models in the literature could be strongly misleading [5]. They furthermore noted [1] that almost all the models were inconsistent with our mechanistic picture (below). Certainly, none of them seem able to explain the size dependent acceleration factors for SAC305 joints evident in Figure 2. The present paper briefly outlines the model proposed by Borgesen et al. and addresses some of the underlying assumptions through both thermal cycling and a unique experiment simulating thermal cycling by alternating between isothermal loading and annealing. Experimental Samples for isothermal cycling were prepared by reflow soldering 30 mil (0.75mm) diameter spheres onto 22 mil (0.55mm) diameter Cu pads on typical BGA component substrates. Soldering was accomplished by first printing a tacky flux through a stencil onto the substrate pads and then placing individual spheres through apertures in a separate ‘bumping’ stencil. Reflow was done in a nitrogen ambient with less than

50 ppm O2 using a Vitronics-Soltec 10-zone full convection oven and a profile with a 245oC peak and 45-60 seconds above liquidus. A row of 18 joints was first annealed for 48 hours at 100oC in order to coarsen and stabilize the secondary precipitates, then loaded simultaneously in an Instron tensile tester. The joints were loaded to the designated creep load at a rate of approximately 0.3MPa/s after which the load was maintained for 5 minutes. The loading was then reversed to bring the joints back to the original position over a period of about 5 minutes. This was followed by a one hour anneal at 100oC, after which the loading sequence was repeated. Figure 3 shows the loading-unloading part of a typical cycle.

Figure 3: Load per joint vs. displacement during the first part of a typical isothermal cycle. All thermal cycling experiments employed fully balanced model CSP components each consisting of a 20 mil thick silicon ‘die’ sandwiched between two 16 mil thick FR4 substrates using thin layers of a rigid flip chip underfill to fully encapsulate the die. In all cases the silicon was sized so that all solder pads were located under it. 16 mil (0.4mm) diameter SAC305 balls were soldered to solder mask defined Cu pads on the component. The corresponding 62 mil thick four layer board had OSP coated non-solder mask defined Cu pads. One set of thermal cycling experiments employed components with 8 x 8 full arrays of 0.38mm diameter pads with five different pitches. The resulting assemblies were cycled in either 0/100oC or -20/100oC with 10oC/minute ramps and 10 minute dwells at either extreme. Another set of thermal cycling experiments employed components with 12 mil (0.3mm) diameter pad openings arranged in a conventional partially depleted array of 256 pads on a 0.8mm pitch. The resulting assemblies were cycled in either 0/100oC or 0/80oC with 10oC/minute ramps and 15 minute dwells. After cycling, solder joints were cross sectioned and the microstructure first characterized by cross polarizer microscopy. Selected samples were further analyzed by Electron Backscatter Diffraction (EBSD).

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Understanding & Model Realistic SnAgCu solder joints are invariably the results of single solidification events during cool-down from reflow, leading to either one or three distinct Sn grain orientations in each. In the latter case these orientations are a result of cyclic twinning [6]. Small solder volumes and pad sizes may allow for such strong undercooling that the twinned grains form an interlaced structure, but typical BGA and CSP scale joints end up with large Sn grains such as those shown in Figure 4.

and the density of importance is the one established by the end of the dwell. The authors argue that this must scale with the creep rate at that point. To the extent that the continuous addition of dislocations to this structure during the reversed loading at high temperature contributes to the recrystallization, this part can be considered dynamic. The authors argue that this contribution may be calculated based on the creep or work during the high temperature dwell only. Neither alternative is compatible with a scaling of life with the total work done in each cycle. Based on such considerations and a number of simplifications Borgesen et al. proposed a first very simple approximation for the life in thermal cycling: NF = ψ / {1 + ξ * tdwell}

Figure 4: Cross polarizer image of typical SAC305 solder joint [7]. It has long been recognized [8 - 12] that the failure of such joints in thermal cycling tends to involve recrystallization within the Sn grains (Fig. 5). Qasaimeh et al. [13] showed crack growth to remain relatively slow until a continuous network of grain boundaries had been established across the high strain region of the joint, and Yin et al. [14] found that the completion of this network tended to take about 1/3 of the total life, independently of cycling parameters. This led to the suggestion that the life could be predicted based on a prediction of the rate of recrystallization.

Figure 5: SAC305 solder with crack through recrystallized region near component pad after thermal cycling. Korhonen et al. [15] showed that reproducing thermal cycling stresses and dwell times in isothermal cycling at any temperature between -25oC and 125oC was not sufficient to cause significant recrystallization. Borgesen et al. [16, 1] argued that the level of recrystallization shown in Figure 5 could only be achieved through the repeated alternations between the establishment of a dislocation cell structure at low temperature and the coalescence and rotation of these cells at a higher temperature characteristic of thermal cycling. Based on this, Borgesen et al. [1] have proposed a model according to which the rate of damage in thermal cycling varies with the density of dislocations and the high temperature dwell time. As far as the dislocation structure built up during the low temperature dwell is concerned, the subsequent recrystallization cannot be considered dynamic

(1)

where ψ accounts for the dislocation density and the term in the parenthesis describes the coalescence and rotation during the high temperature dwell time, tdwell [1]. While this approximation is practical the authors caution that it must be used with considerable care. Notably, ignoring initial precipitate distributions and subsequent coarsening the expression cannot account for effects of solder volume and pad size (Fig. 1), although these are predicted by the general model [1]. Results and Discussion The results in Fig. 1 offer a particularly obvious example of the need for a quantitative mechanistic understanding of the behavior of SnAgCu solder joints in thermal cycling. The general picture outlined above allows us to rationalize a variety of systematic trends otherwise confounding our interpretation of accelerated test results [14, 1]. However, two important issues remain unresolved as far as the model is concerned. The model relates the number of cycles to failure to the dislocation density established during the low and/or high temperature dwell, and thus to the near-steady state creep rate. The question remains as to what constitutive relations might be appropriate for the calculation of this creep rate. Constitutive relations are usually extracted from single sided creep experiments, while the dislocation cell structures built up under reversed cyclic loading are certain to be different. Furthermore, the repeated ‘annealing’ during the high temperature dwells affect the hardening, but as evidenced by the eventual recrystallization they do not completely eliminate the dislocation cell structures and the hardening associated with them. It is therefore not clear how best to account for contributions of primary creep in each cycle. We address this for use in the Borgesen model below. The other issue is to which extent the recrystallization is in fact controlled by an ongoing production of dislocations during the high temperature dwell, i.e. whether the recrystallization process is dynamic. The relative importance of the high and the low temperature dwells is addressed in two different ways. Isothermal Cycling Experiments

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density. For dwell times longer than a few minutes these rates can be reasonably well approximated based on steady-state creep. Dutta et al. have developed microstructurally adaptive constitutive relations to account for the ongoing coarsening of the secondary precipitates during thermal cycling and aging [4]. We propose that these will be well suited for our purposes. Nevertheless, the extraction of functional dependencies in the model from comparisons between thermal cycling results and FEM are still expected to be very sensitive to the approximations still needed in the latter. Figure 6: Creep rate vs. time under a constant load of 8MPa per joint in a typical cycle. Our isothermal cycling experiments allow for the direct measurement of creep vs. time under conditions resembling those encountered in thermal cycling. Figure 6 shows an example for 0.75mm diameter SAC305 joints on Cu pads after a number of preceding cycles. The samples were loaded to 8MPa per joint at a rate of approximately 0.3MPa/s right after cool-down from annealing, and the load then kept constant for 5 minutes. Starting right after the 8MPa/joint is reached the room temperature creep rate is seen to drop rapidly over the first few seconds, but to level off substantially within less than 5 minutes.

Figure 7: Creep rate at the end of 5 minutes under a constant load of 8MPa/joint vs. number of cycles. Figure 7 shows the creep rate at the end of the 5 minute dwell as a function of number of cycles. For reasons associated with the current experimental procedure this measurement shows considerable scatter, but there is no long term trend. The same is true for the total displacement during each dwell (Fig. 8). We conclude that the initial 48 hour anneal tended to stabilize the precipitate distributions after which it only took a few cycles for the behavior to become largely repeatable. This is good news for FEM calculations as it may not be necessary to model every cycle after an initial anneal. Calculations of, for example, the total work per thermal cycle would still require accurate accounting for major contributions from primary creep during both ramps and the early stages of each dwell. The Borgesen model does, however, rely only on calculations of the creep rate near the end of either the low temperature or the high temperature dwell, as this is expected to represent the relevant dislocation

Figure 8: Total displacement at the end of 5 minutes under a constant load of 8MPa/joint vs. number of cycles. Isothermal cycling is therefore ongoing to address dependencies on strain rates and temperatures, and if relevant work, directly. These experiments are extremely labor intensive and time consuming, but so far they appear to offer preliminary insight into the question of whether the recrystallization is in fact predominantly dynamic. The experiments allow for the isolation of effects of individual parameters. Thus, an experiment was first conducted in which 0.75mm diameter SAC305 joints were loaded to 12MPa per joint for 5 minutes, and then loading was reversed and a displacement rate used that returned the sample to the original position within 5 minutes. This was then followed by a one hour anneal at 100oC before the loading sequence was repeated.

Figure 9: Cross polarizer images of near-pad region in two joints after 50 cycles to 12 MPa/joint. Figure 9 shows cross polarizer images of two joints after 50 such cycles. Recrystallization is evident but still much less than what is commonly observed in thermal cycling (Fig. 5). This is, however, believed to be a result of competing damage mechanisms. Isothermal cycling of SAC305 usually leads to failure by transgranular crack growth before much recrystallization [15, 17, 18]. Similarly, in thermal cycling transgranular crack growth competes with the formation of a

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continuous network of grain boundaries across the high strain region in the joint. Joints may therefore also fail by transgranular crack growth in thermal cycling if the cyclic strain range is high enough. Indeed, Figure 10 shows a cross polarizer image of a typical SAC305 joint in our highest strain CSP assemblies after failing in -40/125C cycling. The joint failed after only 149 cycles. The image shows a very limited amount of recrystallization compared to the level shown in Figure 5. Clearly, the sample failed before an obvious network of high angle grain boundaries had evolved across the joint.

Figure 10: Cross polarizer image of 20 mil diameter SAC305 solder joint with crack after 149 cycles of -40/125oC.

Figure 11: EBSD of the region circled in Figure 10. Top – color map showing different Sn orientations; middle – lines indicating boundary misorientations (red 2-5o, green 5-10o, blue >10o). This does not mean that recrystallization was not involved. Figure 11 shows an EBSD image of the region circled in Figure 10. A large density of boundaries with misorientations in excess of 10o is observed in the area of the crack. Other

joints surviving a similar number of cycles showed more or less recrystallization. Work is in progress to assess whether this variability is consistent with expected effects of Sn grain orientations. According to the current version of the Borgesen model [1] reproducing this in our isothermal cycling experiment will require a similar number of cycles and a similar dislocation density in each cycle. If controlled by transgranular crack growth the failure of a SnAgCu joint in isothermal cycling occurs upon the accumulation of a given amount of total work [19, 20]. This is only an approximation, the total work to failure increasing systematically with loading rate, and recent work suggests that a limited variation with cycling amplitude is associated with an increase in the amount of work going to heat due to friction between the opposing surfaces of large cracks toward the end of life [20]. Otherwise, the work to failure did not vary much with temperature, dwell times, or repeated variations in amplitude. Overall, as long as they are controlled by transgranular crack growth experimental life times can thus be extrapolated quite well to lower stresses based on the assumption of a constant work to failure. In the experiment above cycling to 12MPa per joint led to an average work of 1.4x10-4J per cycle and a total life of 50 cycles. Cycling to 6MPa per joint led to an average work of 1.4x10-5J per cycle, suggesting a life of about 500 cycles, while cycling to 8MPa per joint led to an average of 2.7x105J per cycle, and thus an estimated life of 260 cycles. These estimates are upper limits based on the assumption of failure by transgranular crack growth at the lower stresses as well. Recrystallization may cause faster failure than that. Cycling is ongoing at both loads to address that. The predicted life at 8MPa is about twice that of the joints in Figure 10 and 11. This may suggest that the low temperature strain rate is lower, but there may also be time for more recrystallization before failure.

Figure 12: Cross polarizer image of SAC305 solder joint after 100 cycles of ‘simulated thermal cycling’ (alternating isothermal cycling and annealing). So far, cycling to 100 cycles with 8MPa per joint led to a level of recrystallization near the pad that is very limited (Figure 12), but seemingly of the same general magnitude as in the thermal cycling sample in Figure 10 even though the joints are expected to last longer. This would seem to support a picture in which the dislocation density established during

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the low temperature part of the thermal cycle controls the degree of coalescence and rotation during the high temperature dwell that eventually leads to recrystallization.

of a competition between transgranular crack growth and recrystallization, thermal cycling also often leading to a limited amount of the former before completion of a network of high angle grain boundaries by the latter [13, 14].

Figure 15: EBSD of the small crack at the right side of the joint in Figure 14.

Figure 13: EBSD of the near-pad region in Figure 12. Top – color map showing different Sn orientations; middle – lines indicating boundary misorientations (red 2-5o, green 5-10o, blue >10o). Figure 13 provides for a more accurate comparison to the EBSD image in Figure 11. At the top different Sn grain orientations are indicated by different colors. The center image shows small, medium and large boundary misorientations, and the image at the bottom shows a graphic representation of some of the grain orientations. Most misorientations are still small, meaning that this is still early in the dynamic recrystallization process, but the overall picture leaves little doubt that this represents the early stages of the kind of recrystallization seen in Figure 5.

The crack on the left is considerably longer, almost 50µm, and here there are a significant number of high angle boundaries near the tip (Figure 16). In fact, angles approach 60o right at the tip. This might support the assumption that the apparent transgranular crack growth mechanism also involves recrystallization [1]. Since it is not controlled by the formation of an entire network of high angle boundaries modeling of this, and thus prediction of life in isothermal cycling, will however still need to be different. As reflected in the lower parts of Figures 13 and 14 the Sn grain orientations are quite similar. Work is ongoing to assess whether the differences in recrystallization are associated with the small difference in orientation or with the local distributions of secondary precipitates.

Figure 14: Cross polarizer and EBSD images of SAC305 solder joint after 100 cycles of ‘simulated thermal cycling’. Figure 14 finally shows another joint also subjected to 100 cycles. Unlike the one above this joint shows no indication of the beginning of ‘global’ recrystallization across the high strain region of the joint. Instead, it shows the onset of two cracks at either side of the joint. The crack on the right side is very small, less than 10µm in length, and there is no significant misorientation near the tip (Figure 15). This would be consistent with the assumption

Figure 16: EBSD of the larger crack at the left side of the joint in Figure 14. Thermal Cycling Experiments The above was complemented by conventional thermal cycling of model assemblies. Figure 17 shows first a Weibull

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plot reflecting the effect of the high temperature dwell on the life of model CSP assemblies with 0.4mm diameter SAC305 solder joints. Lowering the dwell temperature from 100 oC to 80oC while keeping the low temperature dwell (0oC), dwell times and ramp rates the same led to a slight broadening of the failure distribution and an increase in characteristic life, N 63, by a factor of 2.2.

Figure 17: Cumulative failure distributions for model CSP assemblies in 0/100oC and 0/80oC thermal cycling with 15 minute dwells at both temperature extremes. The effect of the low temperature dwell was much weaker. Figure 18 shows the life of 0.4mm diameter SAC305 solder joints in -20/100oC cycling plotted against the corresponding life in 0/100oC cycling for each of five different model CSP designs. The combined uncertainties of two independently measured N63 values leads to significant scatter, but there is no systematic trend with respect to cyclic strain range. On average, lowering the minimum temperature to raise the temperature range (ΔT) by 20% lead to a 10% reduction in life.

Figure 18: Characteristic life, N63, in -20/100oC vs. life in 0/100oC for model CSP assemblies with 10 minute dwells at both temperature extremes. Broken line represents the life in the two cycles being equal.

The stronger sensitivity to the high temperature dwell is not a result of more effective coalescence and rotation of the dislocation cells there. The dwell time parameter ξ in Equation 1 was found to vary slowly, or not at all, with the dwell temperature. This would suggest that the dislocation density determining the rate of recrystallization during a given cycle, and thus the parameter ψ, depends significantly on the maximum temperature. While the dislocation cell structure established at low temperature is important, the contribution of dislocations created during the high temperature dwell may be dynamic, i.e. according to Borgesen et al. [1] this may actually depend on the work there. The stresses at the high temperature are of course sensitive also to the difference between this and the low temperature extreme, especially at the beginning of the dwell, affecting the creep rates as well. Nevertheless, it seems worth noting that according to the constitutive relations proposed by Dutta et al. [4] the steady state creep rate at the same stress is a factor of 2.2 higher at 100oC than at 80oC, in excellent agreement with the results in Figure 17. Conclusion Except in very high strain assemblies and/or harsh cycling conditions in which life is less than a couple of hundred cycles, it has been shown that the life of a SnAgCu solder joint in thermal cycling is controlled by the evolution of a continuous network of high angle grain boundaries across the high strain region of the joint. Interpretations of thermal cycling experiments are hampered by the inability to vary one important parameter at a time, and by the lack of direct knowledge of the stresses and/or strain rates involved. Not surprisingly, current constitutive relations do not allow for calculation of the latter except perhaps after the first few minutes in each dwell. Solder joints were loaded alternately with thermal annealing while measuring loads and creep rates directly. This allowed for simulation of thermal cycling, except that no load was applied at the high temperature. Preliminary results support a picture in which continuous recrystallization is a result of the build-up of dislocation cell structures during the low temperature dwell and the coalescence and rotation of these during the high temperature dwell. This would mean that the recrystallization cannot be dynamic. Accelerated thermal cycling results do, however, show life in thermal cycling to be much more sensitive to the maximum than the minimum temperature. This would seem to suggest that contributions from dynamic recrystallization during the high temperature dwell are still important. Our mechanical loading and annealing experiments are being extended to include a reversed load at high temperature in order to assess this directly. Acknowledgments This research was supported by the National Science Foundation under Grant No. DMR 1206474. The assistance of Michael Meilunas, Universal Instruments, with sample design and preparation is gratefully acknowledged. References

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