The effects of indium on the microstructural evolution, lattice characteristics, thermal stability and mechanical performance in Sn-3.0Ag-0.5Cu solder alloys

In this study, Sn-3.0Ag-0.5Cu (SAC305) solder alloys and In were carefully weighed according to their precise composition (SAC305, SAC305-5In and SAC305-15In) using a digital precision balance. The weighted samples were then placed in a graphite crucible that was coated with a thin layer of boron nitride. Then, the crucible was placed in an electric furnace with the temperature set to 350 °C. The alloy mixtures were gently stirred every 15 min for 1 h to ensure the homogeneity of the mixture and to minimize the effects of oxidation on the results. Dross was removed before casting the molten alloys into an appropriate mould, and the mould was left to cool down to room temperature until it solidified.

The as-cast bulk solder alloys (10 mm × 10 mm × 5 mm) were ground with SiC sandpaper of different grits ranging from 100 to 2000 grit. Then, the samples were polished with alumina suspension of 1.0 µm and 0.3 µm, followed by oxide polishing suspensions (OPS) of colloidal silica for final surface smoothing to achieve a mirror-like surface finish. Field Emission Scanning Electron Microscopy (FE-SEM) equipped with Energy Dispersive X-ray (EDX) was employed to perform microstructural analysis and to determine the elemental distribution of the samples. Backscattered electron (BES) mode with an accelerating voltage of 20 kV was used in this study.

The thermal properties of the solder alloys were measured using TA Instruments DSC Q10 differential scanning calorimetry (DSC). Samples weighing approximately 5 mg in weight were placed in an aluminium pan, carefully sealed with the aluminium lid and then compressed to ensure the samples were properly enclosed. An empty aluminium pan was used for the reference sample. All samples were heated from 30 °C to 250 °C at a rate of 10 °C/min. When the maximum temperature was reached, the samples were cooled from 250 °C to 30 °C at the same rate. Each sample will undergo the heating and cooling cycle three times in a Nitrogen atmosphere. From the acquired results, the melting points, undercooling, and pasty range of the solder alloys can be determined.

The In-Situ Heating Synchrotron X-ray-Diffraction (XRD) was conducted at the Synchrotron Light Research Institute (SLRI), Thailand at Beamline (BL1.1W). The measurements utilized high-intensity synchrotron radiation with an energy of 12 keV, corresponding to a wavelength of 1.0332 Å, which allowed for high-resolution structural characterization through the use of a monochromator. The XRD patterns for SAC305-5In and SAC305-15In were recorded at intervals of every 20 °C, ranging from 30 °C to 150 °C and 30 °C to 110 °C, respectively. Within this temperature range, all phases exist in solid form. The heating rate was kept constant at 10 °C/min. At every temperature step, the exposure time was set to 180 s. From the results obtained, the XRD patterns and lattice parameters were analysed using PANalytical HighScore Plus software. Using PANalytical HighScore Plus software, the lattice parameters were determined based on the Refine Unit Cell function. This method utilized a least-square refinement algorithm, which minimizes the differences between experimentally observed diffraction peak positions (2θ) and the theoretical values calculated from the selected crystal structure.

To analyze the elemental distribution in the solder alloys, Synchrotron-based micro-X-ray fluorescence (µ-XRF) was performed at beamline 7.2 of the Synchrotron Light Research Institute (SLRI), Thailand. The X-ray source was produced using a 6.5 T superconducting wavelength shifter, delivering an incident beam with energies up to 13.5 keV. A polycapillary focusing optic was mounted on a motorized stage with four degrees of freedom to concentrate the beam to approximately 30 µm in diameter. The sample was positioned perpendicular to the incoming X-ray, while the fluorescence signals were detected using a silicon drift detector (Vortex EM, Hitachi) mounted at a 45° angle relative to the specimen. The measurement was conducted in step-scan mode, covering a 0.5 mm × 0.5 mm area, using a 30 µm step size and exposure durations ranging from 15 to 30 s per point. The resulting fluorescence data were processed using PyMca 5.9.2 to generate elemental maps for each element of the solder composition examined.

An Instron Universal Testing Machine (UTM) was used to determine the mechanical performance of the solder alloys. A tool steel mould was used in the casting process to produce the tensile bars, with a final geometry that meets the standard requirements of ASTM E8/E8M. Figure 1 shows the mould that are used in this study. Prior to casting process, the mould was preheated to a temperature of 150 °C and coated with a thin layer of boron nitride (BN). The solder alloy was melted in a graphite crucible using an electric resistance furnace at 300 °C. The molten solder alloy was held for approximately 15 min to ensure complete melting and compositional homogeneity. The molten solder alloy was then poured slowly into the mould and allowed to solidify naturally at room temperature. All compositions were cast following the same procedure to ensure process consistency. A crosshead speed of 5 mm/min was set for the tensile test on all solder alloy samples. The defects, such as porosity or inclusions, are common in cast alloys, and these issues are known to contribute to reduce elongation. In this study, 10 tensile bars were tested, and the average tensile properties were calculated using the five samples that exhibited the highest elongation. Then, the fracture surface of the tensile specimen was examined using Field Emission Scanning Electron Microscopy (FE-SEM) after the tensile test was performed.

Full size image

Figure 2a and b show the phase diagrams of SAC305-xIn and Sn-In, respectively. Meanwhile, Fig. 3 shows the Scheil solidification path and volume fractions of all phases generated using Thermo-Calc Software TCSLD5: Solder Alloys v5.1. Based on the Thermo-Calc predictions, the addition of In (0, 5 and 15 wt.%) could significantly alter the solidification behaviour and the resulting phase formations. For SAC305 solder alloy as shown in Fig. 3a, the formation of β-Sn occurs at approximately 220.1 °C with β-Sn rapidly becoming the dominant phase as the cooling progresses. As the temperature decreases further, intermetallic compounds of Cu₆Sn₅ and Ag₃Sn begin to form and grow, contributing to the final microstructure of solidified SAC305 solder alloy. Notably, SAC305 solder alloy solidifies over a relatively narrow temperature range. Meanwhile, the addition of In at 5 wt.% and 15 wt.% significantly broadens the solidification temperature range. It was observed that, with 5 wt.% of In in SAC305 solder alloy (Fig. 3c), the formation of β-Sn occurs at a much lower temperature of approximately 210.9 °C. Interestingly, the addition of In leads to the formation of new phases of Ag₉In₄ and γ-InSn₄ in addition to the existing intermetallic compounds of Cu₆Sn₅ and Ag₃Sn. It is particularly noteworthy that when the In content is increased to 15 wt.% Fig. 3e, the solidification path and formation process change significantly, with the solidification starting with the formation of the γ-InSn₄ phase. The high activity of In has reduced the activity of Sn by preferentially forming γ-InSn₄ which becomes the dominant phase in the SAC305-15 wt.% In solder alloy.

Full size image

Full size image

Figure 4 demonstrates the backscattered electron FE-SEM micrographs of the SAC305, SAC305-5In and SAC305-15In bulk solder microstructures and the corresponding EDX mapping of the elements participated. According to the Scheil solidification paths as shown in Fig. 3a, the eutectic reaction occurs in SAC305 solder alloys as:

Full size image

This solidification path results in a typical microstructure, as shown in Fig. 4a, which consists of β-Sn with a network of eutectic IMCs including Ag₃Sn and Cu₆Sn₅. EDX maps in Fig. 4b–d further confirm the elemental distribution of Sn, Ag and Cu. With the addition of 5 wt.% In, the microstructure formed as in Fig. 4e, showing a coarsening of eutectic IMCs. Indium, with an atomic number of 49, is nearly similar to that of Sn, with an atomic number of 50 and shares similar atomic size and physical properties. According to the Hume Rothery rules, this similarity allows for the substitution of Sn atoms by In atoms in Ag₃Sn and Cu₆Sn₅ intermetallic particles. As a result, Sn atoms in these IMC particles are gradually replaced by In atoms, leading to the formation of Ag₃(Sn,In) and Cu₆(Sn,In)₅ phases. This substitution process is thermodynamically favorable and is further supported by EDX point analysis shown in Fig. 5, which quantitatively confirms the presence of In within the IMC particles. Furthermore, as predicted by the Scheil solidification path depicted in Fig. 3c, the addition of 5 wt.% In significantly increases the solidification temperature range by approximately 51.3 °C. The extended solidification temperature range means that the solder remains in a semi-solid state for a longer period during cooling. This event facilitates the atomic diffusion and promotes the growth of IMC particles, thereby contributing to the observed coarsening. The coarsening of eutectic IMCs observed with the addition of In in this study is consistent with the previous findings reported by Lee et al. [17], who also observed significant coarsening particularly in the Ag₃Sn IMC particles. Alongside the formation of Ag₃(Sn,In) and Cu₆(Sn,In)₅ phases, the microstructure, as shown in Fig. 4 (e), also reveals a smaller fraction of irregular polygonal Ag₉In₄ intermetallic compounds. The formation of Ag₉In₄ is attributed to the increased In content, which facilitates the reaction between silver and In during the solidification process. With further addition of In up to 15 wt.% in the SAC305 solder alloy, a significant transformation of the microstructure occurs. The primary Sn-rich phase and the typical eutectic structure, which are dominant at lower In addition, are no longer observed. Instead, as shown in Fig. 4j, the microstructure becomes dominated by the formation of the γ-InSn₄ phase (further confirmed by EDX point analysis in Fig. 5b and accompanied by a more significant presence of Ag₉In₄ IMCs. This microstructure evolution aligns well with the Scheil solidification model predictions as presented in Fig. 3. It can be concluded that a higher addition of In of 15 wt.% alters the solidification path and phase equilibria, leading to the suppression of β-Sn and eutectic IMCs appearing in the solder matrix. Then, Fig. 6 shows the Synchrotron Micro-XRF elemental mapping of Sn, Ag, Cu in SAC305 and Sn, Ag, Cu, In for SAC305-15In at room temperature. A higher intensity corresponds to a greater concentration of the specific element within the alloy. Based on Fig. 6 (d-g), the In mapping demonstrates a relatively even distribution throughout the sample, indicating a uniform distribution of In within the solder matrix. The presence of In significantly modifies the microstructure of the solder alloys. Moreover, the addition of 15 wt.% In has resulted in the formation of γ-InSn₄ as the dominant phase in the solder alloys. This is reflected in the relatively uniform distribution of In within the solder alloys.

Full size image

Full size image

In order to elucidate the phase formation and transformation in SAC305 solder alloys with In addition, In-Situ Synchrotron X-Ray Diffraction (XRD) was employed. The In-Situ Synchrotron XRD pattern presented in Fig. 7 provides critical insights into the phase stability and lattice behavior of SAC305 with the addition of 5 wt.% In and 15 wt.% In. In the case of SAC305-5In solder alloys, the analysis was primarily concentrated on elucidating the β-Sn (101) peak, while for SAC305-15In, the primary focus is on monitoring the γ-InSn₄ (100) peak. These particular peaks were selected because they provide the most direct insight into the structural evolution in the solder alloys, their ability to reflect changes in lattice parameters of a and c and also correspond to the major phase formation as predicted in the volume fraction generated from Thermo-Calc, shown in Fig. 3. Figure 7a exhibits the normalized In-Situ Synchrotron XRD pattern for SAC305-5In, measuring from 30 °C to 150 °C. As the samples of SAC305-5In were heated from 40 °C, the β-Sn (101) peak shifted to a smaller 2θ with increasing temperature due to the thermal expansion. Notably, the β-Sn (101) peak also experiences peak splitting as the samples are heated from 30 °C to 150 °C. The splitting of the peak is suggested due to the dissolution of In atoms in the matrix of β-Sn, causing slight distortions in the crystal lattice. Similar splitting phenomena have also been reported in other solutions treated with various solute concentrations [18]. As the temperature increased to 150 °C, the β-Sn (101) peak appeared to be a single peak. The merging into a single peak suggests that the phase transformations have occurred at this temperature. According to the phase diagram shown in Fig. 2a, at ~ 150 °C and above, the phase transformation from β-Sn + η-Cu₆Sn₅ + Ag₉In₄ + In to a new combination of β-Sn + η-Cu₆Sn₅ + γ-InSn₄ + In phases. This transformation explains the emergence of a single β-Sn (101) peak at elevated temperature as well as the change in the lattice behaviour observed in the In-Situ Synchrotron XRD pattern. These observations are further corroborated by the lattice parameter changes presented in Fig. 7b and c. Since β-Sn have a body-centered tetragonal crystalline structure, the axis a = b in the unit cell of β-Sn phase. As a whole, the β-Sn lattice parameters in the SAC305–5In alloy increase with temperature, reflecting the thermal expansion of its body-centered tetragonal structure during the heating process. In particular, the β-Sn lattice parameter ‘a’ shows a nearly linear increase as the temperature rises from 30 °C to 90 °C. Beyond this temperature range, the β-Sn lattice parameter ‘a’ becomes almost constant with further heating. Sn has an atomic radius of 1.45 Å while In has a larger atomic radius of 1.56 Å. When the larger In atoms dissolve substitutionally on Sn lattice sites, the expansion of β-Sn lattice would occur. As the temperature increase up to ~ 90 °C, the solubility of In in Sn increases, as indicates in the phase diagram of Sn-In in Fig. 2b. Therefore, larger In atoms incorporated into the Sn matrix. This combined effect of thermal expansion and In dissolution leads to the linearly near increasing trend in the β-Sn lattice parameter over this temperature range, and thus highlighting the role of In in modifying the anisotropic thermal expansion of β-Sn. At temperature above ~ 90 °C, the β-Sn closes to its In solubility limit and further heating results in constant expansion of lattice. This phenomenon also in a good agreement with the study done by previous researchers [19‐22]. In contrast, SAC305-15In in Fig. 7d, reveals the different behavior with the analysis mainly focused on γ-InSn₄ (100) peak as this phase constitutes the dominant phase at this composition. Across the temperature range from 30 °C to 70 °C, the γ-InSn₄ (100) peak gradually shifts to a lower degree of 2θ, which may correspond to the thermal expansion of the crystal lattice. However, as the temperature increases from 90 °C to 110 °C, the peak remains consistently a single peak reflecting the stable γ-InSn₄ phase and no phase transformation occurs as simulated by Thermo-Cal Software, as in Fig. 2. This finding was further supported by the lattice parameter of the c-axis in Fig. 7f in where the lattice remains stable from temperature 90 °C to 110 °C. It may also be suggested that the expansion of the crystal lattice c-axes occurs without any contribution from the In solute atoms’ redistribution [23].

Full size image

Figure 8 shows the thermal reactions of SAC305-xIn solder alloys during endothermic (heating) and exothermic (solidification) processes, as determined by DSC heating and cooling curves. It was observed that each SAC305-xIn solder alloys exhibit a single endothermic and exothermic peak during heating and cooling, respectively. Notably, as the In content in SAC305 solder alloys increases from 0 to 15 wt.%, the endothermic peak gradually decreases in intensity and becomes broader in shape. For pure SAC305 solder alloys, the endothermic peak is sharp and intense indicating greater heat absorption. This sharp and intense peak demonstrates that almost all of the SAC305 solder alloy melts within a narrow temperature interval, reflecting a well-defined transition. In contrast, with the addition of 15 wt.% In in SAC305 solder alloys, the peaks not only reduce in height (intensity) but also become noticeably broader and less sharp. This broadening is likely due to the higher volume fraction of the γ-InSn₄ phase in the SAC305-15In solder alloys (as indicated in Fig. 3), as the formation of this lower-melting In-containing phase results in a more gradual melting process and lower overall heat absorption during the heating process. These trends are further supported by the liquidus temperature and pasty range results as shown in Table 1. According to Table 1 the results show a marked decrease in the average liquidus temperature and an increase in the average pasty range with the higher In content. Specifically, the average liquidus temperature reduces by ~ 12% from about 223.8 °C for SAC305 to 197 °C for SAC305-15In, while the pasty range widens significantly. Apart from that, the liquidus temperature was directly compared between measured (from DSC testing) and calculated (from Thermo-Calc) as presented in Table 1 as well. The comparison between measured and calculated revealed that the measured temperature was 2 to 4 °C higher than calculated values. The measured and calculated values presented the same variation with both consisting of a decreasing trend.

Full size image

Another important parameter in elucidating the thermal behaviour in SAC305 solder alloys with In addition is the degree of undercooling, ΔT. Undercooling can be associated with the difficulty in nucleating solid phases in a liquid state. Thermodynamically, a greater degree of undercooling results in a higher driving force for IMC growth, thereby leading to the formation of larger grains in solder, anisotropic behaviour and unfavorable mechanical properties [24]. For a solder alloys with a pasty range, undercooling can be measured by the difference between T_liquidus and T_onset during cooling. Table 1 presents the average undercooling values for the solder alloys SAC305, SAC305-5In and SAC305-15In. It was worth noting that the undercooling values increase with increasing In content of 5 wt.% In, approximately from 32.0 to 34.7 °C. However, as the In content further increased to 15 wt.%, the undercooling was reduced significantly by up to 3.2 °C. This pronounced reduction in undercooling with increasing In content may be attributed to the promotion of new phase formation, particularly the new intermetallic compounds of γ-InSn_4. As revealed in Fig. 3b,d,f, the volume fraction of γ-InSn₄ constitutes a significant portion of the solid phases during cooling. The presence of γ-InSn₄ phases provides more heterogenous nucleation sites, which facilitates the earlier onset of solidification and reduces the degree of undercooling required for the nucleation to occur. The results obtained in this study are well paralleled with the findings of [25], which reported that the addition of In to Sn-2Cu solder alloys lowered the undercooling by up to 5.97 °C.

Figure 9a shows the representative plots of the stress–strain curve for SAC305, SAC305-5In and SAC305-15In. The data clearly demonstrate that the addition of In significantly affects the mechanical behaviour of the SAC305 solder alloys. As shown in Fig. 9b, incorporating 5 wt.% In into SAC305 solder alloy results in a significant improvement in the ultimate tensile strength. Specifically, the ultimate tensile strength of SAC305-5In increases by approximately 50% compared to that of pure SAC305 solder alloys. However, when the In content is further increased to 15 wt.% (SAC305-15In), the ultimate tensile strength decreases, dropping to approximately 43 MPa. While the value is lower than SAC305-5In, it still represents an improvement over the pure SAC305 solder alloys. Meanwhile, for the elongation as indicated in Fig. 9c, the data reveal a more complex trend. The SAC305-5In solder alloy exhibits an increase in the elongation compared to pure SAC305 solder alloy. Conversely, at 15 wt.% In, elongation experiences an 18% reduction relative to pure SAC305, indicating a decrease in ductility in the sample. These results suggest that while adding In to SAC305 solder alloys improves the mechanical strength, there is also an optimal content around 5 wt.% In has a balance between strength and ductility that is most effective. Increasing the In content beyond this level, up to 15 wt.% may further increase the strength to some degree, but at the expense of reduced elongation.

Full size image

The observed changes in mechanical properties with increasing In content can be attributed to microstructural evolution, particularly the formation and growth of the γ-InSn₄ phase. It is known that the γ-InSn₄ phase exhibits a hexagonal crystal structure with space group of P6/mmm and P63/mmc. This crystallographic structure has relatively few slip systems. In metal alloys, slip systems refer to specific crystallographic planes and directions within the crystal lattice that enable atoms to slide past one another under applied stress, allowing the material to undergo plastic deformation [26]. These slip systems control the motion of dislocations within the crystal lattice. Since the γ-InSn₄ phase has relatively few slip systems, it may limit the directions and planes along which dislocations can move. This restriction reduces the ability of the γ-InSn₄ phase to undergo plastic deformation, making the γ-InSn₄ phase inherently brittle. Consequently, as the volume fraction of the γ-InSn₄ phase increases in the SAC305-15In solder alloy (as indicated in Fig. 3), the overall ductility decreases. Notably, the increased presence of this brittle phase results in the solder alloys that exhibit higher strength but a reduced ability to absorb strain before fracture, thereby lowering their capacity for plastic deformation under mechanical loading. In addition, the finding from S. Tian et.al [13] demonstrates that when the In content in Sn-0.7Cu exceeds higher than 3 wt.%, the elongation of the alloys is reduced significantly. This reduction in ductility was attributed to the progressive substitution of Sn atoms by In within the matrix, which alters the deformation behavior and promotes the formation of more brittle intermetallic phases.

The evolution of the microstructure could affect the changes in the mechanical behaviour of in the SAC305-xIn solder alloys. The fracture behaviours after the tensile testing were examined using field emission scanning electron microscopy (FESEM). Figure 10 illustrates the fracture surfaces of SAC305-xIn observed under FESEM. The evolution of the fracture characteristics with increasing In content reveals a distinct transition in the failure mechanism. As indicated in Fig. 10a, the fracture surface of SAC305 solder alloys after tensile testing shows a distribution of small dimples and voids, which are characteristic of ductile failure. These features arise due to the nucleation, growth and subsequent coalescence of microvoids within SAC305 solder alloys under the tensile loading. The existence of microvoid coalescence implies that the solder alloys experienced significant plastic deformation prior to fracture. In addition, the presence of these dimples and voids reflects the ability of SAC305 solder alloys to absorb and dissipate mechanical energy through localized deformation, thereby delaying crack initiation and propagation. Meanwhile, with 5 wt.% In addition in SAC305 solder alloys, the fracture surface exhibited significantly larger and deeper dimples. The existence of these deeper dimples in the SAC305-5wt.%In solder alloys indicates that the materials experienced more extensive plastic deformation before failure, thereby enhancing the ductility of the solder alloys. As indicated in Fig. 2b, the solubility limit of In in Sn at room temperature is ~ 6 wt.%, suggesting that In is largely dissolved in the β-Sn matrix. This indicates that the addition of In may strengthen the solder alloys through solid solution strengthening. In contrast, further addition of In to 15 wt.% in SAC305 solder alloys resulted in more pronounced morphological change, with the evidence of cleavage facets and some microvoids. These characteristics indicate a transition from ductile to mixed ductile and brittle fracture mode. As indicated in Fig. 3f, the addition of 15 wt.% In results in the increased formation of γ-InSn₄ intermetallic compounds at room temperature. Since this phase is known to be brittle, its presence significantly reduced the ductility of the solder alloys.

Full size image