Reversed Phase Transformation of β ฀ α -Sn at Elevated Temperatures towards Quantum Material Integration on Silicon

α -Sn and SnGe alloys have recently attracted much attention as a new family of topological quantum materials. However, bulk α -Sn is thermodynamically stable only at < 13°C. Moreover, scalable integration of α -Sn quantum materials/devices on Si has been hindered by a large lattice mismatch. To address these challenges, we demonstrate compressively strained α -Sn doped with ∼ 2-4 at.% Ge on a native oxide layer on Si, grown at 300-500°C through a reversed β -to- α -Sn phase transformation without relying on epitaxy. The size of α -Sn microdots reaches up to 200 nm, ∼ 10x larger than the upper size limit for α -Sn formation reported before. Furthermore, the compressive strain makes it one of the few candidates for 3D topological Dirac semimetals with interesting applications in spintronics. We ﬁnd that Ge-rich GeSn nanoclusters in the as-deposited materials seeded the reversed β -to- α -Sn transition at elevated temperatures. This process can be further optimized towards optically tunable SnGe quantum material/device integration on Si. that stabilize α -Sn at elevated temperatures. α -Sn containing 0.75 wt.% Ge or 0.6 wt.% were to be 60°C or 90°C, respectively. Impurities the β → α -Sn transition at low previous studies on the effect of Ge gave inconsistent observations. One work showed that Ge could act as secondary nuclei for α -Sn and help the preparation of compact α -Sn specimens from β -Sn. Other studies showed that impurities such as Ge and Si -Sn transition upon annealing at elevated temperature. The average Sn composition in GeSn nanoclusters is derived from the Ge-Ge and/or Ge-Sn peak positions in the Raman spectra.


Introduction
The phase transformation from ductile metallic β-Sn to brittle α-Sn 1 at low temperatures (<13°C) 2 has long been known in history. Commonly known as "tin pest" or "tin plaque" in early literature, legends of this transformation range from the crumbled tin buttons of Napoleon's Grande Armée in Russia 3 to the leaky tin cans of Robert Scott's Antarctica expedition 4 . In the 21 st century, the old "pest" has made a "quantum leap" in that α-Sn has been identified as a very interesting topological quantum material 5 . Relaxed or compressively strained α-Sn is a topological Dirac semimetal (TDS) 5 , a 3D counterpart of graphene with the Dirac cone protected by crystal symmetry and spin-orbit coupling. Recently, topological surface states in α-Sn TDS has been utilized to demonstrate current-induced magnetization switching and damping enhancement for spintronics. 6,7 Under a moderate tensile strain of 0.7%, on the other hand, α-Sn turns into a topological insulator with a topological gap of ∼50 meV. Recently, theoretical predictions have also shown that SnGe alloys with >40% Sn composition can be a TDS under compressive strain 8 . If verified experimentally, this wide range of SnGe alloys could greatly enrich the family of 3D TDS since only a few of them have been discovered so far 5,9,10 . They are also naturally compatible with Si-based integrated circuits as Group IV alloys. However, integrating α-Sn and SnGe quantum materials on Si poses a significant challenge. Limited by lattice matching for hetero-epitaxy, so far most of the α-Sn thin films have been grown on InSb instead of Si because the latter has a large lattice mismatch of ∼20%. Ge-rich GeSn alloys on Si have been extensively explored in optoelectronics community [11][12][13][14] , but high Sn composition SnGe alloys have been rarely studied experimentally since they are well beyond the solubility limit 2 . Furthermore, the thermal stability is also limited by the phase transition to β-Sn at elevated temperatures. For example, the thickness of α-Sn grown on Si is limited to ∼6 nm at room-temperature before it transforms completely to β-Sn 6,7 . Therefore, it would be ideal to find a solution to grow thermally stable α-Sn and SnGe TDS on Si substrates.
To obtain metastable α-Sn, all previous literature adopted the strategy of low temperature growth and introduced different mechanisms to retard its transition to β-Sn at elevated temperatures. For example, previous studies have shown that impurities could help stabilize α-Sn at elevated temperatures. α-Sn containing 0.75 wt.% Ge or 0.6 wt.% Si were observed to be stable at 60°C or 90°C, respectively. 15,16 Impurities could also affect the β→α-Sn transition at low temperatures [17][18][19][20][21][22][23] , but previous studies on the effect of Ge gave inconsistent observations. One work showed that Ge could act as secondary nuclei for α-Sn and help the preparation of compact α-Sn specimens from β-Sn. 24 Other studies showed that impurities such as Ge and Si inhibit the β→α-Sn transition at low temperatures. 25,26 α-Sn thin films grown on InSb by epitaxy could be stable above room temperature up to ∼180°C, yet the thickness is limited to 8 nm. [27][28][29] It was thought that the epitaxial interface helps to stabilize α-Sn, and the α→β-Sn transition temperature depends on the thickness of the film: the thicker the film, the lower the α→β-Sn transition temperature. The slight lattice mismatching between α-Sn and InSb leads to an in-plane compressive strain which could enhance the thermal stability of the α-Sn films. 7,30,31 However, another study showed that negative pressure, or tensile strain, is required to stabilize α-Sn beyond 13°C. 32 Some recent studies on Sn nanoparticles have revealed that the stability of α-Sn is size-dependent. The critical size beyond which α→β-Sn transition starts is reported to be 8, 17 and 1.6 nm for Sn nanoparticles fabricated by epitaxy, chemical reduction and microplasma methods, respectively. [33][34][35] Furthermore, none of the previous literature observed any reversed transformation from β-Sn to α-Sn at >13°C.
Here we report a remarkable finding that α-Sn alloyed with ∼2-4 at.% Ge can be grown and stabilized on native oxide layers on Si substrates without hetero-epitaxy from the substrate. The size of α-Sn microdots reaches up to 280 nm, which is ∼10x larger than the reported critical size mentioned above. Completely unexpected and exactly opposite to the bulk phase diagram, we find that the as-deposited β-Sn reversely transforms into α-Sn upon a rapid thermal annealing (RTA) at 300-500°C, while such transformation does not occur at all at low temperatures, in sharp contrast to pure Sn dots. Based on detailed microstructure analyses, we propose that the Ge-rich GeSn nanoclusters in the as-deposited materials have served as the seeds for α-Sn nucleation and growth at elevated temperatures. To our knowledge, this work is the first demonstration of compressively strained α-SnGe grown at >300°C without epitaxial interface stabilization towards 3D TDS integration on Si. The TDS Sn microdots could potentially be coupled through graphene or quantum dots towards quantum device applications. [36][37][38] . They also demonstrated optical absorption at ∼1250-1750 nm wavelength, which may enable quantum state switching by optical excitation 39 at 1310-1550 nm telecommunication wavelengths, potentially through integration with silicon photonics. This reversed β→α-Sn phase transition can potentially be scaled up using selective epitaxy process for quantum device integration on Si.

Results
Nominally, 50 nm-thick SnGe thin films containing ∼5 at.% Ge are deposited on single-side polished Si(001) substrates by co-evaporating Sn and Ge using a physical vapor deposition (PVD) system. There is a thin layer of native oxide (∼10 nm) on the surface of the Si substrate, as will be detailed later. Note that the Sn-rich SnGe dewets the surface of the native oxide, hence the 50 nm thickness based on the deposition rate is nominal. After the deposition, the as-deposited SnGe thin films are annealed for 2 minutes at a variety of temperatures ranging from 300 to 500°C by a RTA process.
Scanning electron microscope (SEM) images of the SnGe film annealed at 350°C are shown in Figure 1(a),(b). It is clear that the film is not continuous, showing a textured surface with the presence of microdots of different sizes up to ∼500 nm in diameter. Figure 1(c) shows a histogram of the diameters of the dots in Figure 1(a). A two-peak (bimodal) Gaussian fit is applied to the histogram, giving the average diameter of the small nanodots and large microdots to be 34 nm and 223 nm, respectively. The bimodal distribution indicates Oswald ripening process, i.e. bigger SnGe dots grow at the cost of shrinking smaller dots due to surface energy minimization. Further analysis using energy-dispersive X-ray spectroscopy (EDS) shows that the big microdots contain mostly Sn, while the Ge is more abundant in the gaps between the big microdots or in the small nanodots. An EDS line-scan is shown in Figure 1(d). On average there is 2.1 at.% Ge within the big microdots and 5.9 at.% Ge in the gaps or small nanodots. This observation is also consistent with Oswald ripening since Sn is likely the dominant surface-diffusion species from small nanodots to large microdots due to the low melting point, thereby leading to a higher Sn composition (lower Ge composition) in big microdots.
The transition from β-Sn to α-Sn upon RTA were studied via Raman spectroscopy and X-ray diffraction (XRD). Figure 2(a) shows the β→α-Sn transition after RTA annealing at 300°C and at 350°C for 2 minutes in comparison with the as-deposited sample. Interestingly, in addition to the Sn-Sn vibration mode in β-Sn at 125.6 cm -1 , 40 the as-deposited sample actually showed characteristic features of Ge-rich GeSn alloys 41 , with Ge-Ge and Ge-Sn vibration modes clearly visible at 288.2 cm -1 and 250.2 cm -1 , respectively. The Ge-Ge and Ge-Sn peaks are both relatively broad, indicating tiny GeSn nanoclusters in the as-deposited samples. Note that the diffraction peaks from nanoclusters are usually too weak to be observed in XRD data. On the other hand, the cross-sectional EDS mapping and high-resolution transmission electron microscopy (HRTEM) analyses in panels (c) and (d) clearly revealed these Ge-rich nanoclusters at the microdot/oxide interface of the as-deposited samples. From the Ge-Ge and Ge-Sn Raman peak positions, the average Sn composition in these GeSn nanoclusters is estimated to be ∼17 at.%. 41 This is also supported by the EDS and HRTEM analyses in (c) and (d) showing GeSn nanoclusters with ∼10-30 at.% Sn. These Ge-rich GeSn nanoclusters are formed spontaneously via Ge-Sn intermixing during the deposition, indicating that the solubility of Sn in Ge in nanostructures may well exceed that of the bulk phase diagram (∼1 at.%). After annealing at 300°C, a peak corresponding to the Sn-Sn vibration mode in α-Sn is observed at ∼195.9 cm -1 in addition to the β-Sn, Ge-Sn and Ge-Ge peaks. 40,41 . Note that the Ge-Ge peak is also blue-shifted after RTA at 300°C, indicating that the Sn composition in the GeSn nanoclusters is decreased significantly upon the formation of α-Sn. Annealing at 350°C causes a further increased intensity of the α Sn-Sn peak as compared to lower temperature annealing, as well as further blue-shift of the Ge-Ge peak. This is in good agreement with the corresponding XRD data in Figure 3, which also showed the transition of β to α-Sn starting at an annealing temperature of 300°C as well as a better crystallization/larger fraction of α-Sn after annealing at 350°C. It is important to note that the α Sn-Sn peak is present only after annealing, so is the blue-shift of the Ge-Ge peak with increasing temperature. This indicates that there is a phase separation of the GeSn nanoclusters upon RTA, allowing for Sn-rich α-SnGe to nucleate on Ge-rich nanoclusters and initiate β→α-Sn transition at elevated temperatures. Based on the experimental data and analyses discussed above, Figure 2(e) summarizes the mechanism of β→α-Sn transition. The Ge-rich GeSn nanoclusters in the as-deposited samples initiate the nucleation and growth of α-Sn upon RTA, which leads to the reversed β→α-Sn transition, exactly opposite to the bulk phase diagram. Since Ge has a smaller lattice constant than α-Sn, such a nucleation and growth process also leads to compressive strain in α-Sn towards a 3D TDS, as will be detailed later. Furthermore, contrary to the bulk phase diagram again, we note that the β-to-α-Sn transition cannot be obtained by lowering the temperature in this case (i.e. when Ge is incorporated). As shown in Figure 2(b), while pure β-Sn readily started to transform into α-Sn at -20°C (see Supporting Information), with a few percent Ge incorporation we do not see such a transition any more even after cooling down the sample to -50°C and waiting for 45 min. Note that the absence of β→α-Sn transition at low temperature and its presence after RTA at 300-500°C is completely unexpected, given that β-Sn is known to transform into α-Sn at temperatures below 13°C and α-Sn is rarely observed at above 150°C even with coherent epitaxial interface to enhance its stability. 28,42 . Our observation shows that Ge doping enhances β→α-Sn transition at high temperatures while retards it at low temperatures. This result helps to reconcile some seemingly contrary results mentioned earlier.
The XRD patterns of as-deposited and annealed samples are shown in Figure 3(a) to further evaluate the phase transformation upon RTA. In the XRD pattern of the as-deposited sample, all the diffraction peaks are attributed to polycrystalline β-Sn, and no α-Sn peak is observed. Consistent with the Raman data discussed earlier, after RTA at 300-500°C the XRD patterns clearly show the peaks of α-Sn. Especially, samples annealed at 325-400°C show relatively strong α-Sn peaks. Theoretically, the peak intensity ratio of α-Sn and β-Sn I α /I β depends on their mass ratio w α /w β 43

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where RIR is the relative intensity ratio. I α and I β could be calculated by an integral of each peak with a range of ±0.5°.
Then the mass fraction of α-Sn w α /(w α + w β ) is derived from Equation (1) and shown in Figure 3(b). According to α-Sn (01-071-4637) and β-Sn (98-001-3770) powder diffraction files, the RIR of β-Sn is 10.78 while the RIR of α-Sn is 16.65 34 . The mass fraction of α-Sn versus annealing temperature shown in Figure 3(b) indicates that the optimal annealing temperature for β→α-Sn transition is around 325-400°C. The sample annealed at 375°C contains more than 50 wt.% α-Sn. The low amount of α-Sn formed at temperature lower than 325°C and higher than 400°C indicates low temperature cannot overcome the activation energy barrier for α-Sn nucleation on Ge-rich nanoclusters, while high temperature degrades its stability. A further measurement of residual stress in the sample annealed at 325°C (see Supporting Information) indicates that the in-plane stress is compressive. With 95% confidence, the average in-plane strain ε = (−0.63 ± 0.44)% across the Sn microdots, and the stress-free lattice constant of α-Sn phase is a 0 = 6.4540 ± 0.027 Å. Then using Vegard's law, we could estimate that the atomic percent of Ge in α-Sn phase is (4.2 ± 3.3) at.%, consistent with the EDS measurement within the error range. The level of Ge composition and compressive strain falls into the category of TDS as predicted theoretically 5,8 , offering a facile approach for 3D TDS material integration on Si. To further analyse the α-Sn vs. β-Sn regions in the microdots as well as their interfaces, cross sectional samples of the SnGe microdots annealed at 325°C were prepared and analyzed by HRTEM, EDS and scanning nanobeam diffraction acquired through four-dimensional scanning transmission electron microscopy (4D-STEM). Figure 4(a) and (b) show two big microdots, both being α-Sn capped with β-Sn. Figure 4(c) shows a small β-Sn dot. Our further analysis found that small dots with size up to 50 nm diameter mostly remain in β-Sn phase, suggesting that the critical size for nucleating β→α-Sn transformation might be larger than 50 nm at 325°C. The contrast between α-Sn and β-Sn regions in the two big microdots (in (a) and (b)) are also clearly shown in the high-angle annular dark-field (HAADF) STEM images in panels (d) and (g). These two examples indicate that the reversed transformation from β-Sn to α-Sn started at the Sn microdot/oxide interface and grew toward the top surface. The α-Sn phase in both microdots have similar size of ∼200 nm-diameter and ∼100 nm-thickness. These images also clearly show a native oxide layer between the SnGe microdots and the Si substrate. Therefore, the observed α-Sn formation is completed unrelated to hetero-epitaxy on Si. The EDS analysis in Figure 4(d) to (i) clearly show Ge-rich regions near the corner of the microdots, which serve as seeds for α-Sn nucleation and growth, as we have indicated earlier in the Raman analyses. The average Ge composition is 1.4-1.8 at.% inside the α-Sn regions of the microdots, while that in β-Sn regions is slightly lower. Figure 4 Figure 4(j) corresponds to the α-Sn region at the bottom of the microdot (near SiO 2 interface). The d-spacing of α-Sn(111) is 0.367 nm, so the strain is derived to be -1.1% given 3.9 at.% Ge composition near the SiO 2 interface from EDS analyses. Within the error range, this level of compressive strain is also consistent with average in-plane strain from XRD analyses (ε = (−0.63 ± 0.44)%). Figure 4(k) corresponds to the Ge-rich region at the corner of the microdot, which also echoes the Ge-rich region observed in the as-deposited sample (see Fig. 2(c) and (d)). The d-spacing of (111) planes in the boxed region is 0.323 nm, indicating that this region is almost pure Ge after RTA. This observation is fully consistent with the blue-shift of the Ge-Ge peak observed in Raman spectroscopy earlier, showing that the average Sn composition in the Ge nanoclusters is drastically reduced from ∼17 at.% to zero after RTA. . The larger d-spacing of α-Sn(111) near the phase boundary compared with that in 4(j) near the microdot/oxide interface is likely due to the difference in strain considering that the Ge compositions are similar. Furthermore, scanning nanobeam diffraction analyses from 4D-STEM data in Figure 5 clearly show α-Sn and β-Sn phase regions in a microdot. Figure 5(a) is the generated bright field image of this microdot, which is consistent with our earlier HRTEM and STEM images. Figure 5(b) is the reconstructed phase map, showing α-Sn capped with β-Sn. This map was generated by analyzing the spatial evolution of the Bragg disk positions; examples of indexing from different grains are shown in Figure 5(c)-(g) (regions shown are depicted by red rectangles in the corresponding generated dark field images). The corresponding conventional dark field TEM images of this microdot are also given in the Supporting Information. From this analyses, we can see that the α-Sn region is dominated by a large grain shown in Figure 5(c), ∼200 nm wide and 100 nm thick. It is indeed ∼10x larger than the largest critical size for α-Sn formation reported earlier [33][34][35] . The small α-Sn region at the top right corner in Figure 5(d) has a different orientation, overlapping with a small piece of β-Sn region (see Figure 5(g)) in the vertical direction. The β-Sn "cap" region is consisted of 3 small grains, as shown in Figure 5(e)-(g). Overall, the results further confirm the formation of large α-Sn microdots capped with β-Sn. α-Sn occupies ∼70% of the volume in this microdot, which equals to ∼65% weight percentage. This is largely consistent with the ∼50% mass ratio estimated from XRD results in Figure  3(b), by taking into account there are also small β-Sn dots. We also confirmed the presence of both α and β-Sn microdots in the sample annealed at 350°C via HRTEM analyses, as shown in Figure 6. Panel (a) shows a trapezoid-shaped microdot dominated by α-Sn while panel (b) shows a dome-shaped microdot dominated by β-Sn. The difference in shape could due to the different surface energy of α-Sn and β-Sn. Previous studies show that diamond cubic structures (corresponding to α-Sn) tend to form facets on top 45 . Panels (c), (f), and (i) show HRTEM images of different SnGe microdots after 350°C RTA. Panels (d), (g) and (j) zoom into certain regions of these HRTEM images and convert the grayscale into RGB color scale for better visibility. Panels (e), (h), and (k) show the corresponding Fourier transform of the images. First, panels (c) and (f) zoom into two different regions of Figure 6(a), which is dominated by a single-crystalline α-Sn grain extending throughout the microdot with ∼280 nm-diameter and ∼150 nm-thickness. Again, these dimensions are ∼10x greater than the upper size limit for stable α-Sn nanodots reported earlier. [33][34][35] Panels (d) and (e) show an α-Sn(111) inter-planar spacing of d=(0.367 ± 0.03) nm, in good agreement with the expected value of 0.370 nm for 4.2 at.% Ge and -0.63% in-plane compressive strain measured by XRD. Moiré patterns 46,47 are also observed in some regions (see Panel (f)-(h)) due to the vertical overlapping of the predominant α-Sn(111) planes with β-Sn(200) planes at an angle of ∼11.5°. In contrast, panels (i)-(k) shows a pure β-Sn microdot with its (101) planes almost parallel to the surface of the Si substrate. Overall, these data not only confirm the formation of α-Sn at hundreds of nm scale after RTA at 350°C, but also capture adjacent α and β-Sn grains during the reversed phase transformation process from β to α-Sn.
To investigate the optical transitions in these 3D TDS αmicrodots, the optical absorptance spectra of the as-deposited and annealed SnGe samples were measured by an ultraviolet-visible-infrared (UV-Vis-IR) spectrophotometer and plotted in Figure  7. The absorption edge at ∼1100 nm is due to the Si substrate. Samples with a significant amount of α-Sn (annealed at 350°C and 375°C) show an additional absorption edge at ∼1750 nm wavelength compared to those with no or little α-Sn, and the absorption further increases with the decrease of wavelength up to 1250 nm. This additional optical absorption feature at ∼0.7-1 eV caused by α-Sn is consistent with a recent work 48 , and it can be attributed to a transition corresponding to the spin-orbit splitting energy of ∆ 0 = Γ 8v − Γ 7v ≈ 0.7eV 1,48 . Moreover, it has been recently demonstrated that pulsed laser can be used to change the quantum states of 3D TDS 39 , so the measured absorption spectrum of α-SnGe suggests that commonly used lasers for telecommunications at 1310-1550 nm may be used for quantum state switching with potential for integration with silicon integrated photonics. For example, these α-Sn microdots can be easily deposited on Si or silicon nitride waveguides or microresonators, where the input optical power can be evanescently coupled to α-Sn 3D TDS to switch the quantum states. This feature may offer an approach to achieve optically tunable quantum materials and devices on silicon.
We also grew SnGe samples with higher nominal Ge composition and found a nominal Ge composition of 5-8 at.% in the as-deposited microdots is optimal for seeding β→α-Sn transformation. This is briefly summarized in the Supporting Information.

Discussion
In summary, we demonstrate that Ge-doped α-Sn can be grown on native oxide on Si substrate via a reversed phase transformation from β to α-Sn at 300-500°C, seeded by Ge-rich GeSn nanoclusters in the as-deposited materials. Such a reversed phase transformation has not been reported in previous literature. Referring to Figure 2(e), this reversed phase transformation is implemented via the following steps. (1) Ge-rich GeSn nanoclusters (with ∼17 at.%.) are spontaneously formed via Ge-Sn intermixing during the deposition process. From HRTEM analyses, they are mostly located near the microdot/oxide interface and the corners of the microdots. (2) Upon RTA at 300-500°C, the GeSn nanoclusters undergo phase separation into Ge and Sn, in which α-Sn starts to nucleate around the Ge-rich nanoclusters. This process initiate the reversed β to α-Sn transition. It also induces compressive strain in the α-Sn microdots (due to the lattice mismatch with Ge-rich seeds) towards 3D TDS materials integrated on Si. (3) In relatively large microdots with diameters greater than 50 nm, α-Sn can reach the critical size for nucleation and grow from the interface towards the surface, transforming β-Sn into α-Sn during the growth process. This process tends to result in α-Sn microdots capped with β-Sn (e.g. Figure 4(a),(b) and Figure 5), and some big microdots can be almost fully transformed to α-Sn (e.g. Figure 6(a)). (4) Small dots less than 50 nm in diameter cannot reach the critical size for nucleating α-Sn, thereby mostly remaining β-Sn. Such size-limited phase transformation has indeed been observed in other material systems such as silicides 49 .
An important implication of this work is that Ge-rich GeSn with ∼17 at.% Sn can effectively serve as a template for α-Sn quantum material growth at elevated temperatures. Since such GeSn virtual substrate have already been developed for mid-infrared optoelectronic device integration on 8-inch Si wafers 50 , they can be readily applied to implement monolithic Sn-rich SnGe quantum material and device integration on Si. Another implication, especially interesting for the GeSn infrared material and device community,is that that the actual room-temperature solubility of Sn in Ge nanostructures can reach as high as ∼17 at.% Sn since such a high Sn composition was spontaneously achieved during physical vapor deposition at room-temperature, indicating that it is thermodynamically favored. This is a favorable result for the stability of GeSn mid infrared materials and devices that require ∼10-20 at.% Sn composition.
Our work also clarifies a couple of questions related to previous literature: (1) We showed that the formation of α-Sn at elevated temperatures does not have to rely on epitaxy from single crystal substrates. Therefore, although lattice-matching epitaxy is preferred, actually it is not a necessary condition to stabilize α-Sn; (2) We found that the incorporating Ge facilitates α-Sn formation at elevated temperatures while retards its formation at low temperatures. This results reconciles some of the  seemingly contradictory reports in previous literature. [24][25][26] (3) We confirm that a moderate compressive strain on the order of -0.5 to -1 % does not necessarily retard the formation of α-Sn.
In conclusion, we show that Ge-doped α-Sn can be stabilized up to ∼200 nm in dimensions at room temperature without the help of coherent epitaxial interface as in the common case of α-Sn growth on InSb substrates. A compressive strain is introduced to Ge-doped α-Sn, which makes it one of the few candidates for 3D TDS according to theoretical predictions. This work presents the first instance of α-Sn grown at relatively high temperature on Si without epitaxy, breaking through the long-existing constraints in substrate selection, size/thickness limit, and thermal stability of α-Sn growth towards topological quantum material/device integration on Si. Furthermore, α-SnGe's optical absorption at ∼1250-1750 nm wavelength suggests a potential possibility to achieve optically tunable quantum materials and devices on silicon when integrated with silicon photonic circuits. Since we found that Ge-rich GeSn nanoclusters can effectively seed the growth of Sn-rich α-SnGe in this study, this process could potentially be scaled up using selective epitaxial growth on GeSn virtual substrates on Si 50 for high-quality quantum device integration.