
==== Front
Sci Rep
Sci Rep
Scientific Reports
2045-2322
Nature Publishing Group UK London

39227611
70298
10.1038/s41598-024-70298-6
Article
Metastable phases of Ag–Si: amorphous Si and Ag-nodule mediated bonding
Nakayama Koji S. kojisn@sanken.osaka-u.ac.jp

1
Nishijima Masahiko 1
Zhang Yicheng 1
Chen Chuantong 1
Ueshima Minoru 2
Suganuma Katsuaki 1
1 https://ror.org/035t8zc32 grid.136593.b 0000 0004 0373 3971 F3D, The Institute of Scientific and Industrial Research, Osaka University, 8-1 Mihogaoka, Ibaraki, Osaka, 567-0047 Japan
2 grid.480124.b 0000 0001 0425 4575 Daicel Corporation, 2-18-1, Konan, Minato-ku, Tokyo, 108-8230 Japan
3 9 2024
3 9 2024
2024
14 196181 6 2024
14 8 2024
© The Author(s) 2024
2024
https://creativecommons.org/licenses/by-nc-nd/4.0/ Open Access This article is licensed under a Creative Commons Attribution-NonCommercial-NoDerivatives 4.0 International License, which permits any non-commercial use, sharing, distribution and reproduction in any medium or format, as long as you give appropriate credit to the original author(s) and the source, provide a link to the Creative Commons licence, and indicate if you modified the licensed material. You do not have permission under this licence to share adapted material derived from this article or parts of it. The images or other third party material in this article are included in the article’s Creative Commons licence, unless indicated otherwise in a credit line to the material. If material is not included in the article’s Creative Commons licence and your intended use is not permitted by statutory regulation or exceeds the permitted use, you will need to obtain permission directly from the copyright holder. To view a copy of this licence, visit http://creativecommons.org/licenses/by-nc-nd/4.0/.
Metastable phases such as supersaturated solid solutions, supercooling, and amorphous phases are well-known in metallurgy. They are often composed in non-equilibrium states and can be transformed into a stable phase by overcoming an energy barrier with driving forces. Particularly, it has been widely used for material strengthening and heterogeneous nucleation of precipitates in solids is mainly induced by heat treatments for supersaturated solid solutions. However, little is known about the metastable phases of the Ag–Si alloy, although it is a well-known simple binary eutectic alloy. Here, we show that the metastable phases composed of amorphous Si and supersaturated Ag solid solution are induced by the eutectic reaction under rapid cooling of Ag–Si. Furthermore, the solute Si in the Ag matrix reacts with oxygen to precipitate Ag by-products, which grow as nodules. The Ag nodules have high crystallinity and robust interfacial structures, and the nodule growth leads to the formation of cross-links between the Ag–Si particles. We also demonstrate the Ag nodule-mediated bonding where the rapidly cooled Ag–Si ribbon is directly used as a bonding medium, indicating the possibility of using it as a high-temperature bonding material with low-temperature processes.

Subject terms

Engineering
Materials science
Nanoscience and technology
issue-copyright-statement© Springer Nature Limited 2024
==== Body
pmcIntroduction

The eutectic alloy system is important for amorphization because the liquid becomes increasingly viscous as it is cooled to immediately above the eutectic temperature and fails to crystallize1. Such a phase transition is useful for device bonding technologies because the deep eutectic depressions in the liquidus allow solidification at low temperatures. In particular, the Au–Si eutectic alloy, which is the first amorphous alloy discovered2, has attracted attention in die bonding of semiconductor packaging. The eutectic temperature (TE) of Au–Si is 363 °C3, which requires relatively high-temperature bonding processes4,5, but the properties of Au–Si, such as low coefficient of thermal expansion6, low resistivity7, and high shear strength8, meet the key requirements for the bonding devices.

Far less attention has been paid to the Ag–Si eutectic alloy as a bonding material because TE of Ag–Si is 835 °C9, which is too high for the device bonding processes. Except for the high temperature of TE, the equilibrium phase diagram of Ag–Si is similar to that of the Au–Si alloy, where the face-centered cubic (fcc) structure of Ag and the diamond structure of Si have negligible mutual solubility. The metastable states were created by a liquid quenching technique where the alloy was levitation-melted and quenched on to a Cu substrate. In this case, the solubility of Si in the Ag matrix was reported to be up to 30 at % and the metastable phases of Ag3Si or Ag2Si were suggested10. However, the amorphous Si was not obtained.

In this study, we first show that the amorphous Si appears in the eutectic texture where the phase separated due to the eutectic reaction in the Ag–Si particles. Conventionally, the underlying concept of amorphous alloys has been considered as the mixing of multi-atomic compositions, which leads to the stabilization of the liquids due to the negative values of heat mixing11. Here, we present a new concept of amorphization that induced by the significant supercooling due to the deep eutectic depression of the liquidus temperature near the eutectic composition of Ag–Si. The similar amorphization was recently found in the Al–Si eutectic system12, but the amorphous phases were expected to appear in the hypereutectic region of 25–45 at % Si in Al–Si13.

Second, we show that the isothermal oxidation of Si in the supersaturated Ag solid solution leads to the formation of Ag nodules. This is also unexpected because the previously reported growth mechanisms of Ag whiskers, nanowires, and hillocks are caused by solid-state ionic conduction14,15, sulfide deposition16,17, stress-induced migration18,19, and the difference in thermal expansion coefficients20,21. The Ag nodules are grown until they bond between Ag–Si particles, indicating the possibility of utilizing them as a bonding material.

Finally, we demonstrate that the Ag–Si ribbon, which is prepared by a single-roll melt-spinning method, can be used as a bonding film. The significance of the Ag nodule-mediated bonding is to utilize the low-temperature precipitation of Ag by-products induced by the oxidation reaction and the stability of Ag as a bonding material. This approach provides a new insight into bonding technologies, especially for future power devices based on wide bandgap semiconductors that require high temperature stability.

Results

Structural characterization of rapidly cooled Ag–Si

Figure 1a shows a low magnification scanning transmission electron microscopy (STEM) image of the Ag84Si16 alloy particle. The bright field (BF) STEM image shows the Ag–Si particle with a diameter of 1.1 μm. The bright lamellar structures are surrounded by dark features inside the particle, and they are composed of a columnar structure. The length and diameter are about 200 nm and 30 nm, respectively. At the bottom of the particle in Fig. 1a, the columns are perpendicular to the image, so that they look like round shapes. We focus on these columnar structures. Figure 1b and c are the energy dispersive X-ray spectrometry (EDS) elemental mapping obtained from the Ag-L line and the Si-K line, respectively. Figure 1d is the synthesized image of the Ag-L and Si-K lines. The EDS maps clearly indicate that the phase of Si and Ag is separated and the column structures correspond to Si. Therefore, the bright feature of lamellae in the BF-STEM image in Fig. 1a is Si. We note that the green Si dots in Fig. 1c are homogeneously distributed in the Ag matrix, while the Ag dots do not appear in the Si columns in Fig. 1b. Analysis of the EDS data (Fig. S1) shows that the average solubility of Si in the Ag matrix is about 4%. This indicates that the supersaturated Ag solid solution can be formed by the rapid cooling, but the solubility of Ag in Si is negligible.Figure 1 STEM-EDS observations for as-atomized Ag–Si particles. (a) The BF-STEM image of the Ag–Si particle at a low magnification. The bright lamellae are surrounded by the dark area and are composed of columnar structures. The length and diameter are about 200 nm and 30 nm, respectively. The EDS elemental map of (b) the Ag-L line and (c) the Si-K line is obtained from the column features, indicating that the column composition is Si. (d) The synthesis image of Ag-L and Si-K lines. (c) The high-resolution BF-STEM image reveals that the Si column has a disordered atomic structure. (f) The FFT pattern obtained from the disordered atomic region of Si shows no clear spot patterns.

Figure 1e shows a high-resolution BF-STEM image focused on one of the Si columns. The diameter is approximately 30 nm. The image clearly reveals a disordered atomic structure of Si surrounded by the Ag crystalline lattices. In this area, the overlapping of Ag and Si may not occur because the sample thickness is sufficient to be thin. Figure 1f shows the fast Fourier transformed pattern obtained from the disordered region of Si. There are no clear spots associated with crystalline phases. Therefore, the columnar Si is amorphous.

The result is important because this is the first demonstration of Si solidification in the amorphous state in the Ag–Si system. To provide firm evidence for the presence of amorphous Si, we conduct the X-ray diffraction (XRD) on a master alloy, as-atomized particles, and melt-spun ribbons of Ag–Si. Figure 2a shows the XRD result of the Ag84Si16 master alloy. All diffraction peaks belong to fcc Ag and diamond Si crystals. Figure 2b shows the XRD result of the as-atomized Ag84Si16 particles. In this case, the peaks are assigned to fcc Ag, but not to diamond Si. The absence of Si in the diffraction pattern can also be recognized in the XRD result of the melt-spun Ag89Si11 ribbons, as shown in Fig. 2c.Figure 2 XRD results on a master alloy, as-atomized particles, and melt-spun ribbons of Ag–Si. (a) For the Ag84Si16 master alloy, all diffraction peaks belong to fcc Ag and diamond Si crystals. (b) For the as-atomized Ag84Si16 particles, the peaks are assigned to fcc Ag, but not to diamond Si. (c) For the melt-spun Ag89Si11 ribbons, the absence of Si in the diffraction pattern can also be recognized.

Thermal analysis of fine Ag–Si particles

Figure 3a shows the results of differential scanning calorimetry (DSC) for the fine Ag–Si particles (diameter < 5 μm). First, the exothermic peak appears between 220 and 280 °C and the peak position depends on the O2 concentration. Second, the large exothermic peak appears at 340 °C and the peak position is independent of the O2 concentration. Finally, the small exothermic peak appears at 440 °C in addition to the first and second peaks when the O2 concentration is low at 0.1%. The calorific value of the peak at 340 °C increases with increasing the O2 concentration. The previous DSC study of the porous Si films indicated that the oxidation of Si starts at about 300 °C22. Thus, the peak at 340 °C should correspond to Si oxidation. For the peak at 440 °C, we examine the nanoporous amorphous Si that was prepared by the dealloying process from the Al–Si alloy particles12. Figure S2 shows the DSC results of the nanoporous amorphous Si, which indicates the exothermic peak at 420 °C. Therefore, we assign that the peak at 440 °C corresponds to the crystallization temperature of the amorphous Si in the Ag–Si particles.Figure 3 DSC and SEM results after the heat treatments. (a) The DSC result shows that the first exothermic peaks appear between 220 and 280 °C, the second peaks are located at 340 °C, and the third peak appears at 440 °C. The third peak is detected only when the O2 concentration is low (0.1%). (b) The sample is isothermally heated at 280 °C for 10 min under the O2 flow of 0.1% concentration. Small bright protrusions appear on the particle surfaces. (c) Same as (b) but heated for 60 min. It is evident that some particles are connected by nodules. (d) Same as (b) and (c) but heated for 180 min. The growth of nodules is significant by aging. The spaces between the particles are filled with the nodule materials. Furthermore, much darker features of Si lamellae extend circumferentially around the particles. (e) The particles were heated at 280 °C for 180 min under the N2 flow. It shows that the no nodules appear.

Ag precipitation induced by oxidation reaction

The exothermic peaks at 220–280 °C in Fig. 3a imply the existence of metastable states that may react with oxygen. To understand the origin of the exothermic peaks in the low-temperature region, the Ag–Si particles were isothermally heated at 280 °C under the O2 flow of 0.1% and the changes in the microstructure were observed by scanning electron microscope (SEM). Figure 3b shows the cross-sectional SEM image of the particles after the isothermal heating for 10 min. Fine eutectic textures can be seen inside the particles and small bright protrusions appear on the particle surfaces. The inset indicates the SEM image, showing the surface features where the protrusions of about 300 nm are visible. The images in Fig. 3c and d are obtained after heating for 60 and 180 min, respectively. The growth of the nodules is evident by aging at 280 °C. Especially, after the heating for 180 min, the spaces between the particles begin to fill with nodule materials. The fine dark eutectic features inside the particles extend circumferentially around the particles as the nodules increase (See also Fig. 4a).Figure 4 HAADF-STEM observation after the heat treatments. (a) The HAADF-STEM image for the cross section of the particle that has been isothermally heated at 280 °C for 60 min in the flow of O2 (0.1%). The image contrast of the fine Si lamellae becomes much darker compared to that of the lamellae inside the particle. (b) The atomic structure of the nodule taken from the white arrow in (a). The distance between atoms is 4.09 Å, which corresponds to the Ag lattice constant. (c) The illustration of the Ag nodules and the Ag–Si particles. The supersaturated Ag solid solution (Ag-αSi) is labeled.

To verify the nodule appearance, we heated the particles at 280 °C for 180 min under the N2 flow only. We found that the no bright protrusions were observed in the SEM image, as shown in Fig. 3e. In fact, the DSC results obtained under the flow of N2 or Ar in Fig. 3a indicate that the exothermic reaction is negligible. Therefore, the formation of nodules is associated with the oxidation process. Furthermore, we prepared the as-casted Ag–Si alloy which was slowly cooled in the quartz tube after the melt. In this case, the solidification follows the equilibrium phase diagram and the the supersaturated Ag solid solution is not expected. The SEM image in Fig. S3 after the heating at 320 °C for 180 min in the O2 flow indicates that the large Si regions are dispersed in the Ag matrix, which are very similar to the microstructures reported by Weber23. Importantly, no nodule structures are recognized in Fig. S3. Therefore, we conclude that the solute Si in the supersaturated Ag solid solution is necessary for the nodule formation.

Structural characterization of Ag nodules

To obtain the details of the structural information, we conducted the high-angle-annular-dark-field (HAADF)-STEM imaging for the cross-sectional particles. The sample was isothermally heated at 280 °C for 60 min in the flow of low O2 (0.1%). Figure 4a shows the HAADF-STEM image of the nodules that appear on the upper left and lower right of the central particle. The lower right nodule connects the two particles. The red arrows indicate that the nodule is connected to the eutectic textures. This feature suggests that the heterogeneous nucleation of Ag occurs at the end of eutectics. Furthermore, the image contrast of the fine Si lamellae becomes much darker as compared to the inside of the particle but the Ag matrix does not change. These features were also seen in the SEM image in Fig. 3d, where the darker feature of the lamellae extends circumferentially around the particles. The image contrast in HAADF provides the changes in atomic weight because the intensity of the scattered electrons from heavy atoms is stronger at high angles. Therefore, the chemical composition of the Si lamellae has been changed (see the EDS results in Fig. 6). Figure 4b shows the high-resolution HAADF-STEM image taken from the edge of the nodule marked by the white arrow in Fig. 4a. It shows the lattice structure of the nodule. The spacing of 4.09 Å corresponds to the lattice constant of crystalline Ag. The inset shows the electron diffraction pattern, indicating the [001] direction of an fcc crystal lattice. This reveals that the nodule has a high crystallinity. Figure 4c illustrates the relationship between the nodules on the particle and the dark region inside the particle.

Discussion

Amorphous Si induced by phase separation

Shao et al. reported the thermodynamic properties of the supercooled liquid of pure Si based on the comparison of Gibbs free energies for Si phases24. They predicted that a first-order phase transition from the liquid phase to the amorphous phase occurs when the liquid could be supercooled to 1169 °C. This corresponds to the supercooling temperature (ΔT) of 1414–1169 = 245 °C, where the melting point (TM) of Si is 1414 °C, as shown in Fig. 5. The relative supercooling (ΔT/TM) of pure Si is 0.17. Recently, the amorphous Si has been found in the rapidly cooled Al–Si eutectic alloy12. The eutectic reaction leads to the phase separation between Si and Al. In this case, the phase separated Si immediately undergoes the significant supercooling of ΔT = TM−TE = 837 °C, where TE = 577 °C for Al–Si. Thus, ΔT/TM of Si in Al–Si is 0.59. ΔT is well below the predicted temperature of the first-order phase transition, so that crystallization is avoided. In this study, TE = 835 °C and ΔT = 579 °C. Thus, ΔT/TM of Si in Ag–Si is 0.41. This is much higher than that of pure Si (ΔT/TM = 0.17). Thus, the phase-separated Si is significantly supercooled, which is associated with the failure of crystallization.Figure 5 Equilibrium phase diagrams of Ag–Si, Al–Si, and Au–Si, and the predicted first-order phase transition. The phase diagrams of Ag–Si, Al–Si, and Au–Si are calculated by Thermo-Calc (Thermo-Calc Software AB) using the database of SSOL8, SGTE Solution Database v8. For pure Si, the dashed line shows the predicted first-order phase transition temperature of 1169 °C by Shao et al.24. This corresponds to ΔT = 245 °C. For Au–Si, TE = 363 °C α and ΔT of 1051 °C. For Al–Si, TE = 577 °C and ΔT = 837 °C. For Ag–Si, TE = 835 °C and ΔT = 579 °C.

However, this is not the only reason for the amorphization because the amorphous Si does not appear in the as-casted Ag–Si alloy. The BF-STEM image in Fig. 1e shows very fine structures of amorphous Si with a diameter of 30 nm. The size effect may play a role in amorphization, where the nucleation of crystals is suppressed in such small domains. Furthermore, since the thermal conductivity of Ag (429 W/m K) is almost twice that of Al (237 W/m K)25, we expect the preferential solidification of the fcc Ag phase. The nanoscale Si columns are surrounded by the highly thermally conductive Ag phases. Therefore, the diffusion of the latent heat during the Si solidification also leads to the failure of crystallization.

Oxidation of Si and precipitation of Ag

We have shown the dark contrast in the HAADF-STEM image of the Si lamellae that appear circumferentially around the particles in Fig. 4a. It is assumed that the oxidation of Si in the supersaturated Ag solid solution induces the destabilization of Ag atoms. The diffusion of Ag atoms to the particle surface is promoted, leading to the growth of Ag nodules, as illustrated in Fig. 4c. In this case, the SiOx species are expected to appear near the Ag nodules. In the case of SiO2, the reaction should be Ag-αSi + O2 → Ag + αSiO2, where Ag-αSi is the supersaturated Ag solid solution with Si dissolved and α is at %. In this study, we have obtained α ~ 4 at% from the EDS data (Fig. S1).

To confirm the oxidation near the Ag nodule, we obtained the EDS compositional maps around the nodule. Figure 6a represents the HAADF-STEM image where the nodule is located next to the dark feature on the particle surface. Figure 6b is the EDS elemental map of the Ag-L line, showing the feature of the Ag nodule. Figure S4a show the EDS spectra measured on the nodule. The nodule feature disappears in the EDS elemental map of the Si-K line, as shown in Fig. 6c. These results reveal that the nodule is composed of pure Ag. Intriguingly, the high intensity of Si-K line appears on the outermost particle surface. Figure 6d shows the EDS elemental map of the O-K line. This reveals that the high-intensity area of the O–K line corresponds to the dark area in the HAADF-STEM image. We have also obtained the EDS concentration profiles in the box area of yellow dashed line in Fig. 6a. Figure 6e shows the EDS profile running diagonally across the HAADF-STEM image. The X-ray count of the Ag-L line decreases significantly across the dark feature in the HAADF-STEM. This indicates the mass transport of Ag to the nodule formation. The X-ray count of the Si-K line is nearly constant except for the outermost surface of the particle. Figure 6f is the magnified view of the O-K line. The number of X-rays count on the dark feature is doubled. Thus, the oxidation of Si occurs underneath the Ag nodule. Figure S4b shows the EDS spectra obtained on the oxidized area, indicating that the relative concentration of O is about 11%.Figure 6 STEM-EDS analysis for the Ag nodule and the particle surface. (a) The HAADF-STEM image shows the nodule next to the dark feature on the particle surface. (b) The EDS map of the Ag-L line as the same area in (a). It reveals that the composition of the nodule is Ag. (c) The EDS elemental map of the Si-K line shows that the nodule feature disappears. (d) The EDS elemental map of the O-K line reveals that the high intensity area of the O-K line corresponds to the dark feature in the HAADF-STEM image. (e) The EDS profiles along the yellow box in (a). (f) The magnified view of the O–K line clearly indicates that the number of the X-ray count on the dark feature has doubled. Thus, the oxidation of Si occurs underneath the Ag nodule.

Ag nodule-mediated bonding

Finally, the Ag89Si11 ribbons were prepared by the single roll melt-spinning method. This method allows the cooling rate about 105–106 °C/s, which is one to two orders of magnitude higher than that of gas atomization1. Therefore, metastable states are expected to be further involved in the ribbons. Figure 7a shows the SEM image of the Ag89Si11 ribbon surface after the heating at 300 °C for 60 min in air. The image indicates the appearance of many nodules grown on the ribbon surface. The density of the nodules is 0.9 μm-2, and the size of the nodules is up to 5 μm, which is about 5 times larger than the size of the nodule that appeared on the particle, as shown in Fig. 4a.Figure 7 Sintering of Ag–Si ribbon and shear tests. (a) The SEM image of the Ag–Si ribbon surface after the heating at 300 °C for 60 min in air. (b) A demonstration of the sinter bonding of Ag–Si ribbon between the Ag-coated Cu chip and the substrate. (c) The results of the shear test. Sintering is performed at 250, 300, and 350 °C for 60 min in air under a compression pressure of 20 MPa.

Figure 7b shows a demonstration of Ag89Si11 ribbon sintering between the Ag-coated Cu chip and the substrate. Sintering was performed at 250, 300, and 350 °C for 60 min in air under a compression pressure of 20 MPa. When the as-spun ribbon is sintered at 300 °C, the shear strength values are distributed in the range of 5.8–13 MPa, as plotted by the black dots in Fig. 7c. The shear strength increases to 16.6 MPa when the sintering temperature is increased to 350 °C. However, the bonding is incomplete when sintered at 250 °C and the shear strength is only 3.0 MPa, even though the Ag nodules have been grown on the ribbon surface. This is due to the poor adhesion configuration caused by the surface roughness of the ribbon, where the molten alloy is solidified by an instantaneous contact with the Cu roller that is rotating at a high speed. Therefore, we polished the ribbon surface with #400 sandpaper to obtain a smooth surface and conducted the sintered bonds at each temperature. Bonding is completed even when the polished ribbon is sintered at 250 °C. In this case, the shear strength is 11.1 MPa, as plotted by the red squares in Fig. 7c. At 300 °C, the shear strength increases to 18.2–20.5 MPa, which is about double that of the as-spun ribbon. At 350 °C, the shear strength is 20.2 MPa, which is 22% higher than that of the as-spun ribbon. The obtained shear strength is promising for practical bonding because the die shear strength of Pb95Sn5 (wt%) commercial paste is 19–24 MPa26.

Summary

Our results shed light on the metastable states that exist even in a simple binary Ag–Si system when it is rapidly cooled. There are several implications from this study as follows. First, we have shown that the amorphous Si appears on the phase-separated nanoscale domains in the fine Ag–Si eutectic texture. The high-resolution STEM image in Fig. 1e, the FFT pattern in Fig. 1f, and the XRD results in Fig. 2b and c reveal that there is no long-range order of Si. However, the short-range order should  exist due to the strong covalent bonds of Si. Second, the precipitation of crystalline Si from the supersaturated Ag solid solution should occur upon heat treatments. This would be related to the observation of the high Si-K line intensity appearing on the outermost particle surface in Fig. 6c. Third, the Ag nodule-mediated bonding offers all-solid-phase bonding without the use of organic solvents. This is an alternative to the Au–Si eutectic bonding, which avoids the high cost of Au. Finally, the DSC result in Fig. 3a shows the shift of the exothermic peaks as a function of oxygen concentration. This suggests the possibility of the Ag-nodule mediated bonding at much lower temperatures as the oxygen concentration increases.

Methods

Figure 8 shows the sample preparation of Ag–Si particles and ribbons. We used the gas atomization method to produce the fine particles of Ag84Si16 (at%) alloy. The cross-sectional SEM image for the as-atomized particles shows that the nanoscale Ag–Si eutectic textures that are formed by the rapid cooling of the gas jets during atomization. The fine darker features correspond to the Si lamellae. The Si primary crystals and Ag dendrites are not recognized while the Ag dendrites appear in the Ag89Si11 particle, as shown in Fig. S5. The surface morphology of the particles is very smooth. The classification of particle sizes was made by the air-flow driven type sieve and the particles with the diameters below 5 μm were collected. For the bonding tests, we prepared the Ag89Si11 ribbons by a single-roll melt-spinning method (NEV-A05, Nisshin Giken Co.) with the Cu roll speed of 3000 rpm, as illustrated in Fig. 8. The diameter of the Cu roller was 200 mm. The width and thickness of the ribbon were 8 mm and about 40 μm, respectively.Figure 8 Preparation of Ag–Si particles and ribbons. The image and illustration of the sample preparation for the Ag–Si particles by the gas atomization method and the Ag–Si ribbons by the single-roll quenching method. The SEM image is a cross section for the as-atomized Ag–Si particles. Very fine eutectic textures are formed by the rapid cooling during gas atomization. The dark features correspond to the Si lamellae.

The microstructures of the surfaces and the cross-section of the particles were characterized by SEM (JSM-IT100, JEOL Ltd.). For SEM observations, the particles were embedded in a conductive resin (Technovit 5000, Kulzer) and the cross-sectional surface was finally polished with Al2O3 particles of 0.06 μm in diameter. The structural characterization at the atomic scale was carried out by spherical aberration corrected STEM with EDS (JEM-ARM200F, JEOL Ltd.). The acceleration voltage was 200 kV, the beam current was 125 pA, the probe size was 1.0 nm, the convergence angle was 16–18 mrad, and the HAADF detection angle was 68 mrad. The EDS analyses were performed to obtain the characteristic X-rays of the elements, and the elemental maps were obtained from the integrated intensity of the X-ray counts. The energy resolution was 128–135 eV at the Mn-Kα line (5.9 keV). The EDS detector was the Si-drift detector with an area size of 30 mm2. The pixel sizes were 512 × 512. The dwell time was 0.1 ms/pixel, which results in 26 min for one image. We typically took about 100–200 images for EDS analyses.

For STEM-EDS observations, the cross-sectional samples were prepared by a focused ion beam (FIB) equipped with SEM (VERSA 3D, Thermo Fisher Scientific). We used to embed the particles in an epoxy resin (G2, Gatan) prior to FIB processing. For the FIB processes, a micro-surgical lift-out technique was used for sample preparation and the low acceleration voltage of 8 kV was used to minimize the damage from the Ga beam irradiation and to prevent temperature rise during the process. The sample thickness is below about 100 nm. XRD (SmartLab, Rigaku) measurements were carried out with the Cu target (45 kV, 200 mA). The X-ray detector was D/teX. The scan speed was 10 degree/minute.

Thermal analyses were performed by differential scanning calorimetry (DSC) with a heating rate of 20 K/min (DSC204H, Netzsch). The DSC instrument was also used for the isothermal heat treatments. To systematically verify the oxidation effect, the synthesis gas of N2 and O2 was generated by controlling the flow ratio between N2 and O2. For example, the O2 concentration of 0.1% indicates the synthesis of 99.9% of N2 and 0.1% of O2. We also used pure N2 or Ar gas to distinguish the oxidation reaction. The sample mass for the DSC measurements was approximately 5 mg.

A conventional bond tester (ESR-4000, Nordson) with a shear load cell (DS200) was used for the shear tests. The Cu chip size is 4 × 4 mm2 with a height of 0.5 mm. The magnetron sputtering system was used to grow the Ag layer on the Cu chip and substrate. The thickness of the Ag layer was about 1 μm. The adhesion layer of Ti or Ni with a thickness of 100 nm was grown on the Cu surface before the deposition of Ag.

Supplementary Information

Supplementary Figures.

Supplementary Information

The online version contains supplementary material available at 10.1038/s41598-024-70298-6.

Acknowledgements

We thank Y. Kodama, R. Okumura, and N. Kagami for the assistance of sample preparations. We also thank “Advanced Research Infrastructure for Materials and Nanotechnology in Japan (ARIM)” of the Ministry of Education, Culture, Sports, Science and Technology (MEXT).

Author contributions

K.S.N and M.U. conceived the original idea of this work. K.S.N. conducted the main experiments and wrote the paper. M.N. and K.S.N. performed the STEM-EDS observations. Y.Z. prepared the ribbon samples and conducted sinter bonding and shear testing. K.S. provided the experimental resources and supervised the experiments conducted at Osaka University. M.U. and C.C. supervised the corporative research project between Daicel and Osaka University. M.N. and C.C. provided extensive reviews and discussions in revising the paper. All authors have proofread the paper.

Data availability

The data that support the findings of this study are available from the corresponding author upon reasonable request.

Competing interests

The authors declare no competing interests.

Publisher's note

Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.
==== Refs
References

1. Greer AL Metallic glasses Science 1995 267 1947 1953 10.1126/science.267.5206.1947 17770105
Greer, A. L. Metallic glasses. Science 267, 1947–1953 (1995).17770105 10.1126/science.267.5206.1947
2. Klement W Willens RH Duwez P Non-crystalline structure in solidified gold–silicon alloys Nature 1960 187 869 870 10.1038/187869b0
Klement, W., Willens, R. H. & Duwez, P. Non-crystalline structure in solidified gold–silicon alloys. Nature 187, 869–870 (1960).10.1038/187869b0
3. Okamoto H Massalski TB The Au–Si (gold–silicon) system Bull. Alloy Phase Diagr. 1983 4 190 198 10.1007/BF02884878
Okamoto, H. & Massalski, T. B. The Au–Si (gold–silicon) system. Bull. Alloy Phase Diagr. 4, 190–198 (1983).10.1007/BF02884878
4. Haubold M Lin Y-C Frömel J Wiemer M Esashi M Geßner T A novel approach for increasing the strength of an Au/Si eutectic bonded interface on an oxidized silicon surface Microsyst. Technol. 2012 18 515 521 10.1007/s00542-012-1440-1
Haubold, M. et al. A novel approach for increasing the strength of an Au/Si eutectic bonded interface on an oxidized silicon surface. Microsyst. Technol. 18, 515–521 (2012).10.1007/s00542-012-1440-1
5. Wolffenbuttel RF Low-temperature intermediate Au–Si wafer bonding; eutectic or silicide bond Sens. Actuators A 1997 62 680 686 10.1016/S0924-4247(97)01550-1
Wolffenbuttel, R. F. Low-temperature intermediate Au–Si wafer bonding; eutectic or silicide bond. Sens. Actuators A 62, 680–686 (1997).10.1016/S0924-4247(97)01550-1
6. Suganuma K Kim S-J Kim K-S High-temperature lead-free solders: Properties and possibilities JOM 2009 61 64 71 10.1007/s11837-009-0013-y
Suganuma, K., Kim, S.-J. & Kim, K.-S. High-temperature lead-free solders: Properties and possibilities. JOM 61, 64–71 (2009).10.1007/s11837-009-0013-y
7. Lani S Bosseboeuf A Belier B Clerc C Gousset C Aubert J Gold metallizations for eutectic bonding of silicon wafers Microsyst. Technol. 2006 12 1021 1025 10.1007/s00542-006-0228-6
Lani, S. et al. Gold metallizations for eutectic bonding of silicon wafers. Microsyst. Technol. 12, 1021–1025 (2006).10.1007/s00542-006-0228-6
8. Abouie M Liu Q Ivey DG Eutectic and solid-state wafer bonding of silicon with gold Mater. Sci. Eng. B 2012 177 1748 1758 10.1016/j.mseb.2012.09.005
Abouie, M., Liu, Q. & Ivey, D. G. Eutectic and solid-state wafer bonding of silicon with gold. Mater. Sci. Eng. B 177, 1748–1758 (2012).10.1016/j.mseb.2012.09.005
9. Oleslnski RW Gokhale AB Abbaschian GJ The Ag–Si (Silver–Silicon) system Bull. Alloy Phase Diagr. 1989 10 635 640 10.1007/BF02877631
Oleslnski, R. W., Gokhale, A. B. & Abbaschian, G. J. The Ag–Si (Silver–Silicon) system. Bull. Alloy Phase Diagr. 10, 635–640 (1989).10.1007/BF02877631
10. Suryanarayana C A new metastable phase in the silver–silicon system J. Less Common Met. 1974 35 347 352 10.1016/0022-5088(74)90248-3
Suryanarayana, C. A new metastable phase in the silver–silicon system. J. Less Common Met. 35, 347–352 (1974).10.1016/0022-5088(74)90248-3
11. Takeuchi A Inoue A Classification of bulk metallic glasses by atomic size difference, heat of mixing and period of constituent elements and its application to characterization of the main alloying element Mater. Trans. 2005 46 2817 2829 10.2320/matertrans.46.2817
Takeuchi, A. & Inoue, A. Classification of bulk metallic glasses by atomic size difference, heat of mixing and period of constituent elements and its application to characterization of the main alloying element. Mater. Trans. 46, 2817–2829 (2005).10.2320/matertrans.46.2817
12. Okada, J., Nakayama, K. & Wada, T. Porous amorphous silicon, method for producing porous amorphous silicon, and secondary battery. Jpn. Patent No. 6896261 (2019).
13. Zhao ZK Li JC Jiang Q Hypereutectic Al–Si binary alloys prepared by melt spinning method Adv. Eng. Mater. 2004 6 303 306 10.1002/adem.200300578
Zhao, Z. K., Li, J. C. & Jiang, Q. Hypereutectic Al–Si binary alloys prepared by melt spinning method. Adv. Eng. Mater. 6, 303–306 (2004).10.1002/adem.200300578
14. Farnell GC Judd FS The formation of silver whiskers by exposure of silver bromide microcrystals Br. J. Appl. Phys 1965 16 1413 1414 10.1088/0508-3443/16/9/128
Farnell, G. C. & Judd, F. S. The formation of silver whiskers by exposure of silver bromide microcrystals. Br. J. Appl. Phys 16, 1413–1414 (1965).10.1088/0508-3443/16/9/128
15. Ohhachi T Taniguchi I Growth control of silver whiskers on α-Ag2S, α-Ag2Se and α-Ag2Te J. J. Appl. Phys. 1969 8 1062 1063 10.1143/JJAP.8.1062
Ohhachi, T. & Taniguchi, I. Growth control of silver whiskers on α-Ag2S, α-Ag2Se and α-Ag2Te. J. J. Appl. Phys. 8, 1062–1063 (1969).10.1143/JJAP.8.1062
16. Terabe K Hasegawa T Nakayama T Aono M Quantized conductance atomic switch Nature 2005 433 47 50 10.1038/nature03190 15635405
Terabe, K., Hasegawa, T., Nakayama, T. & Aono, M. Quantized conductance atomic switch. Nature 433, 47–50 (2005).15635405 10.1038/nature03190
17. Drott J Growth of silver sulphide whiskers Acta Metall. 1960 8 19 22 10.1016/0001-6160(60)90134-6
Drott, J. Growth of silver sulphide whiskers. Acta Metall. 8, 19–22 (1960).10.1016/0001-6160(60)90134-6
18. Tohmyoh H Yasuda M Saka M Controlling Ag whisker growth using very thin metallic films Scr. Mater. 2010 63 289 292 10.1016/j.scriptamat.2010.04.013
Tohmyoh, H., Yasuda, M. & Saka, M. Controlling Ag whisker growth using very thin metallic films. Scr. Mater. 63, 289–292 (2010).10.1016/j.scriptamat.2010.04.013
19. Chen C Nagao S Jiu J Zhang H Sugahara T Suganuma K Dry-growth of silver single-crystal nanowires from porous Ag structure Appl. Phys. Lett. 2016 108 263105 10.1063/1.4955019
Chen, C. et al. Dry-growth of silver single-crystal nanowires from porous Ag structure. Appl. Phys. Lett. 108, 263105 (2016).10.1063/1.4955019
20. Chen C Suganuma K Low temperature SiC die-attach bonding technology by hillocks generation on Al sheet surface with stress self-generation and self-release Sci. Rep. 2020 10 9042 10.1038/s41598-020-66069-8 32494058
Chen, C. & Suganuma, K. Low temperature SiC die-attach bonding technology by hillocks generation on Al sheet surface with stress self-generation and self-release. Sci. Rep. 10, 9042 (2020).32494058 10.1038/s41598-020-66069-8
21. Oh C Nagao S Kunimune T Suganuma K Pressureless wafer bonding by turning hillocks into abnormal grain growths in Ag films Appl. Phys. Lett. 2014 104 161603 10.1063/1.4872320
Oh, C., Nagao, S., Kunimune, T. & Suganuma, K. Pressureless wafer bonding by turning hillocks into abnormal grain growths in Ag films. Appl. Phys. Lett. 104, 161603 (2014).10.1063/1.4872320
22. Salonen J Lehto V-P Laine E Thermal oxidation of free-standing porous silicon films Appl. Phys. Lett. 1997 70 637 639 10.1063/1.118294
Salonen, J., Lehto, V.-P. & Laine, E. Thermal oxidation of free-standing porous silicon films. Appl. Phys. Lett. 70, 637–639 (1997).10.1063/1.118294
23. Weber L Equilibrium solid solubility of silicon in silver Metall. Mater. Trans. A 2002 33A 1145 1150 10.1007/s11661-002-0216-0
Weber, L. Equilibrium solid solubility of silicon in silver. Metall. Mater. Trans. A 33A, 1145–1150 (2002).10.1007/s11661-002-0216-0
24. Shao Y Spaepen F Turnbull D An analysis of the formation of bulk amorphous silicon from the melt Metall. Mater. Trans. 1998 29A 1825 1828 10.1007/s11661-998-0006-4
Shao, Y., Spaepen, F. & Turnbull, D. An analysis of the formation of bulk amorphous silicon from the melt. Metall. Mater. Trans. 29A, 1825–1828 (1998).10.1007/s11661-998-0006-4
25. Haynes WM CRC Handbook of Chemistry and Physics 2014 CRC Press
Haynes, W. M. (ed.) CRC Handbook of Chemistry and Physics (CRC Press, 2014).
26. Maruyama M Matsubayashi R Iwakuro H Isoda S Komatsu T Silver nanosintering: A lead-free alternative to soldering Appl. Phys. A 2008 93 467 470 10.1007/s00339-008-4807-5
Maruyama, M., Matsubayashi, R., Iwakuro, H., Isoda, S. & Komatsu, T. Silver nanosintering: A lead-free alternative to soldering. Appl. Phys. A 93, 467–470 (2008).10.1007/s00339-008-4807-5
