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Abstract
Silicon-tin (Si1-xSnx) films have been grown using plasma-enhanced chemical vapor deposition on Si (001) substrate. X-ray photoelectron spectroscopy characterization of the thin films show successful substitutional incorporation of Sn in Si lattice up to 3.2%. The X-ray diffraction characterizations show epitaxial growth of Si1-xSnx films (001) direction. The Sn incorporation has been measured using X-ray photoelectron spectrometry and the film uniformity was confirmed using energy dispersive X-ray spectrometry.
© 2020 Optical Society of America under the terms of the OSA Open Access Publishing Agreement
1. Introduction
In the past decades, Si Photonics has been pursuing integration of optoelectronics on the electronics chips to enable high-speed data transfers [1,2]. Finding a direct bandgap material that is compatible with Si has been pursued to fabricate highly efficient optoelectronic devices. Development of GeSn as a viable alloy that can achieve this goal has been successfully demonstrated by different groups [3–5]. A variety of optoelectronic devices including photodiodes [6,7], detectors [8–10], and lasers [11–14] have been fabricated using GeSn alloys. In order to expand the success in the achievements of GeSn, incorporation of Si in the GeSn matrix has been investigated to provide the wider bandgap SiGeSn ternary alloy in order to simultaneously engineer the bandgap and lattice constant. The SiGeSn alloy has been used in various device designs as barriers in the GeSn quantum well devices for better optical confinement or as a wider bandgap alloy for solar cells [15–19]. However, the development of SiGeSn has been always difficult due to the technical challenges such as low mutual solubility of Sn in Si and Ge (<0.5%) [20], incompatibility of growth temperatures, and lattice mismatch (up to 20%). Although different groups have reported the growth of SiGeSn material [21], however, the reports have been limited and the material has not been studied as widely as GeSn [22,23]. The SiGe and GeSn alloys have been studied in details and the missing link in achieving the knowledge to grow SiGeSn ternary alloy is Si1-xSnx. The Si1-xSnx alloy can be direct bandgap material like GeSn with the difference of having a larger bandgap due to the high direct band edge of Si at 3.2 eV (Γ-valley) [24,25]. However, achieving direct bandgap in Si1-xSnx requires high Sn incorporations (>25% Sn) in contrast to GeSn (∼8%Sn) [24]. Sn solubility is very low for both cases of Ge and Si in thermal equilibrium conditions, however, Sn is successfully shown to be incorporated in Ge as high as 22.3% with high material quality and optical efficiency [13]. Therefore, it is expected that high Sn content Si1-xSnx alloys with Sn composition significantly higher than the reported solid solubility to be grown as well. The desired Sn incorporation for pure SiSn is between 25% to 48% for a fully direct bandgap alloy. However, lower Sn contents (5-15%) are desirable for SiGeSn as the growth method could be implemented to achieve SiGeSn with direct or pseudo-direct energy bandgap.
Since 1995, only a few research groups have been able to show the successful growth of Si1-xSnx. The initial studies were using molecular beam epitaxy (MBE) which resulted in 5-16% Sn incorporation for single layer and multiple quantum well structures with very small thicknesses of 1-2 nm [26]. Reports on the synthesis of nanoscale Si1-xSnx were limited to laser ablation and crystallization methods which result in nanocrystal Si1-xSnx with metallic behavior or quantum-confined nanocrystals with compositions up to 5% [27]. However, a more recent study shows that SiSn nanostructures with 91% Sn can form [28]. The best attempts in achieving high Sn-content Si1-xSnx films using crystallization methods are reported by the Zaima et al., at the University of Nagoya by achieving up to 40% Sn on SOI, Si, Ge and InP substrates [29–31]. As of now, the only report of Si1-xSnx growth using CVD is by the Arizona State University using high order Silanes (Si3H8) and an unstable Sn precursor of SnD4 [24]. The growth of Si1-xSnx using conventional SiH4 and even higher-order Silane (Si2H6) is not possible below 300°C due to not being sufficiently reactive [24]. Although increasing the growth temperature breaks surface terminated Si-H bonds on the substrate, it concurrently disfavors Sn incorporation due to possible Sn droplet formation and phase segregation. Growth of Si1-xSnx using these precursors has not been successful at a broad range of growth conditions which resulted in either amorphous and/or phase-segregated films. Therefore, Tolle, et al. have used Si3H8 and SnD4/H2 mixtures at 275 °C to grow Si1-xSnx on a GeSn buffered Si [24]. It is worth mentioning that SiSn properties have been studied for a variety of applications such as photoelectricity [32], near-infrared brain imaging technology [33], or as donors in Si [34]. Table 1 shows a summary of the best growth achievements on Si1-xSnx.
Table 1. Comparison of different methods achieving Si1-xSnx in Caltech, AIST, Nagoya University, Arizona State University (ASU), and the University of Arkansas (UA)
In this work, we have adopted an ultra-high vacuum plasma-enhanced chemical vapor deposition (UHV-PECVD) technique by using commercially available precursors of Silane (SiH4) and tin tetrachloride (SnCl4) in high-hydrogen dilution environment. The deposited Si1-xSnx films are characterized by different methods to investigate their material and optical properties. The optical characterization of the films has been performed using Raman and Spectroscopic Ellipsometry. The X-ray diffraction (XRD) characterization is performed to study the crystallinity of the alloys. X-ray photoelectron spectrometry (XPS) is performed to verify the elemental incorporation. Energy-dispersive X-ray spectrometry (EDX) performed on the film surface through scanning electron microscopy (SEM) verifies uniform incorporation and smooth surface morphology.
2. Experimental
2.1 Growth method
The Si1-xSnx films were grown in a custom-built cold wall UHV-PECVD system. The CVD system consists of a load-lock chamber that is pumped down to 10−8 Torr and a growth chamber with the base pressure of 10−10 Torr described by Mosleh, et al. [35]. The growth chamber is equipped with a cryogenic sublimation pump to pump down oxygen and water vapor prior to the growth. The Si1-xSnx films were grown directly on a 4-inch Si (001) substrate without any buffer layer. The Si substrates were cleaned using Piranha etch solution for 15 min followed by HF dipping in 5% diluted HF for 30 seconds to remove the thin oxide. The wafers were transferred to the vacuum chamber within 15 minutes of preparation. The films were grown at a substrate temperature of 300°C. In order to make sure that the substrate has reached the growth temperature, a 30 min heat soak was performed prior to the flow of the gases. The growth matrix was designed based on successful low-temperature Si epitaxy in high hydrogen dilution and low plasma that was previously reported by Mosleh et al. [36]. The growth was accomplished using commercially available Si precursor (SiH4), Tin tetrachloride (SnCl4) as precursors, and Hydrogen (H2) as the carrier and diluent gas. The hydrogen flow rate was set to 100 standard cubic centimeters per minute (sccm) for all samples. The Silane (SiH4) was fixed at 5 sccm which made the dilution ratio at 20. The flow rate of SnCl4 was varied from 0.03-0.2 sccm. The capacitively coupled plasma (CCP) with 13.56 MHz radio-frequency (RF) power supply and an L-shaped automatic impedance matching network was connected to the counter electrode and the samples were grounded. The 4-inch Si wafer holder was grounded, as the cathode of CCP, while conductor plate in parallel with the wafer holder was powered, as the anode. The counter electrode was 20 mm away from the substrate as shown in Fig. 1(a). The equivalent capacitance was formed between Si wafer and conductor plate. The plasma discharge between the wafer and the metallic plate is shown in Fig. 1(b) as the glowing region. The color change in the plasma glow from purple (b) to navy blue (c) could be observed as a result of the increase in the SnCl4 flow rates. Although the gas entry is from the right, no difference is observed in the color of the plasma glow towards gas entry and the plasma was uniformly formed between the electrodes as it can be seen in Fig. 1(b) and (c). This is in contrast with the last report of GeSn plasma growth where the color was different in front of the gas entry and the Sn was shown to be depleted at the edge and resulted in non-uniform growth [37].
Fig. 1. (a) A schematic diagram of the PECVD system. The generation of plasma between the anode plate and Si wafer in the CVD reactor at a low SnCl4 flow rate (b) and high flow rate (c).
The Plasma power was varied between 5 to 20 W which makes the plasma power density to be 0.05-0.2 W/cm2. The substrate was rotated to ensure uniform growth. The growth temperature was set at 300°C and the chamber pressure was fixed at 0.5 Torr for all the growths. The growth conditions are summarized in Table 2.
Table 2. Growth matrix of Si1-xSnx on Si samples.
2.2 Characterization method
The Si1-xSnx samples were characterized using different methods to investigate the optical and material properties. The film thickness and absorption coefficient spectra were measured using a J.A. Woollam V-Vase ellipsometer. The ellipsometry data was measured at a fixed angle of 70 degrees. Raman spectroscopy of the samples was performed using a Horiba LabRAM HR Evolution Raman Spectrometer with a 532 nm laser as an excitation source and 100X objective to better focus on the top film. The crystallinity and lattice constant of the films were measured using a Panalytical Phillips X'pert PRO XRD system that was equipped with a parabolic mirror and a 4-bounce Ge (220) monochromator. The XRD 2Theta-Omega scans were measured along the Si (004) plane. The XRD reciprocal space maps (RSM) were measured from $({\bar{2}\bar{2}4} )$ plane to investigate the relaxation and Sn content of the Si1-xSnx lattice. The XPS was also used to measure the composition of the films. The XPS system uses a thermo K-α X-ray (1486.7 eV) with Ag 3d5/2 peak monochromator with full width at half-maximum (FWHM) energy resolution of 0.50 eV. The SEM/EDX was used to study the surface morphology and uniformity of Sn incorporation in the films.
3. Optical and material characterization results
3.1 Raman spectroscopy
The Si1-xSnx samples were characterized with Raman spectroscopy to explore the effect of flow ratio and plasma power on the material quality as well as the Sn incorporation. Si1-xSnx films have not been comprehensively studied for all the peaks, however, the literature suggests that the Si1-xSnx peak from transverse optical (TO) mode should be observed between 300-400 cm-1 [24,38]. The shifted Si-Si TO mode peak is a good indication of the formation of Si1-xSnx layer where the shift could be used to estimate the incorporation of Sn in the Si lattice [38]. In our Si1-xSnx films, four separate peaks were observed in the Raman spectra of the Si1-xSnx films (Fig. 2). Two peaks were from the substrate and two peaks from the Si1-xSnx film.
Fig. 2. The Raman spectra of the Si1-xSnx samples on Si. Due to the low thickness of the Si1-xSnx films, and high resolution of the Raman system, high count peaks from Si substrate is observed along with the Si1-xSnx films. The Si-Si 2TA peak at 300 cm-1, Si-Si TO peak at 520 cm-1 is arising from the Si substrate. The peak at 330 cm-1 is attributed to Si-Sn TO peak and the peak observed as a shoulder at 480-490 cm-1 is due to the strained Si-Si peak in the Si1-xSnx film.
The Raman spectroscopy was performed with two different lasers of 532 nm and 785 nm. The 785 nm results are not presented here because the penetration depth of 532 nm laser is ∼700 nm in Si and spatial resolution of the Raman system for 100X objective is ∼300 nm, while these specifications are much higher for the 785 nm laser wavelength (penetration depth 12000 nm; spatial resolution: 500 nm). Therefore, the 532 nm laser could give better results from the SiSn film. Due to the low thickness of the Si1-xSnx films (92-185 nm), and high resolution of the Raman system, high-count peaks from Si substrate is observed along with the Si1-xSnx films. The two substrate peaks are the second-order transverse acoustic mode (2TA) at 300 cm-1 and the TO mode peak at 520 cm-1. The peak at 330 cm-1 is attributed to Si-Sn TO mode and the peak observed as a shoulder at 480-490 cm-1 is due to the strained Si-Si TO mode peak in the Si1-xSnx film. It can be observed that the Si-Si TO peak is more pronounced as the plasma power is increased which could be interpreted as thicker film formation.
3.2 X-ray diffraction
In order to study the crystallinity and lattice constant of the Si1-xSnx layers, the samples were characterized using XRD technique. Figure 3 shows the 2θ- ω and RSM of the samples. Figure 3(a) delineates the change in the diffraction angle as a result of the increase in the SnCl4 flow rate. The shift in Si1-xSnx peaks from 69.12° of Si to lower angles which indicates incorporation of Sn in Si lattice. A similar pattern is observed in Fig. 3(b) where the effect of plasma power is studied in the incorporation of Sn in the Si lattice. As the 2θ-ω scans only show the out-of-plane lattice constant (${a_{\bot}}$) of the films, a reciprocal space map is performed on the sample to find the in-plane lattice constant (${a_\textrm{||}}$), strain and total lattice size.
Fig. 3. XRD results for Si1-xSnx on Si substrates (001). 2θ- ω scan of Si1-xSnx samples from (004) plane shows the increase in the Sn composition as a result of the increase in the SnCl4 flow rate (a) and increase in the plasma power at fixed flow rates show a reduction in the Sn content which resulted in lower out of plane lattice constant.
Figure 4 shows the RSM of sample D indicates that the Si1-xSnx layer is fully pseudomorphic. All other RSMs show the same result with fully pseudomorphic films which is consistent with the thin film thicknesses calculated using ellipsometry fittings. Using the measured in-plane (${\textrm{a}_\parallel }$) and out-of-plane (${\textrm{a}_ \bot }$) lattice constants from the XRD RSM scans, the total lattice constant of SixSn1-x layer is calculated through $\textrm{a}_\textrm{0}^{\textrm{SiSn}}\textrm{ = }\frac{{\textrm{(}{\textrm{a}_ \bot }\textrm{ + 2}{\textrm{a}_\parallel }\textrm{)}}}{{\textrm{(1}\,\textrm{ + }\,\textrm{2}\mathrm{\nu}\textrm{)}}}$ where ν is the elastic (Poisson) ratio [39]. Different bowing parameters have been proposed for Si1-xSnx, however, none has been verified with experimental data due to the limited available compositions of Si1-xSnx alloys. In many lattice calculations such as SiGeSn alloys, the lattice bowing parameter was assumed to be zero [25], however, recent modeling results have reported 0.084 Å or 0.2 Å which matches more with the experimental data [40]. In our calculations, shown in Table 3, the latter is taken as the bowing parameter (${b^{SiSn}}$ = 0.2).
Table 3. Calculated values of d-spacing, Lattice constant, and incorporation percentages from XRD 2θ- ω peak positions with the consideration that the films are all fully pseudomorphic found from RSM scans.
Figure 5(a) shows the effect of the SnCl4 flow rate in the incorporation of Sn in the Si lattice. It can be seen that increasing the flow rate results in higher Sn incorporation, however, this slightly reduces the material quality as evidenced by the increase in the FWHM of the peaks. This suggests a relaxation initiation process, as evidenced by the possibility of the film approaching critical thickness. However, all films were observed to be fully pseudomorphic. It has been shown that tensile strain in the film is a hindrance to Sn incorporation in GeSn on Ge films [41]. Therefore, it is expected that continued growth under the same condition would result in the thicker film as well as higher Sn incorporation. Besides, higher flow rates could result in higher Sn incorporation as the trend is indicating. Figure 5(b) shows the effect of the plasma power on the Sn incorporation in Si lattice. We can see that, although increasing the power from 5 W to 10 W has slightly increased the incorporation, however, as it can be seen from the 2Theta-Omega scans in Fig. 3(b) that the peak intensity has dropped one order of magnitude, which indicates a lowering of the material quality. This can be better observed as the power increases to 20 W where both incorporation and peak intensity decreases. This is attributed to the deterioration of the crystal structure by bombarding the surface with higher energy ad-species.
Fig. 5. (a) Effect of SnCl4 flow rate and (b) plasma power in the incorporation of Sn in the Si lattice.
3.3 XPS, SEM, and EDX analysis
The Si and Sn concentration profiles in the Si1-xSnx layers were determined by XPS depth profile analysis. The depth profile of the samples was achieved by Ar+ sputtering during which the charge compensation was performed by electrons and low energy ions of Ar+. The Ion Gun has a spot size of 400 µm spot size ranging from 100 eV to 4keV. The portions of the XPS spectrum, showing the Si and Sn binding energies are shown in Fig. 6. The Sn 3d peaks are 493 eV (3d3/2), 485 eV (3d5/2) binding energies and Si 2p (2p1/2, 2p3/2) is at 99.4 eV. The Si has an asymmetrical peak due to the overlapping of (2p1/2, 2p3/2) peaks which are very close to each other.
Fig. 6. The XPS results from depth profile of sample D. (a) Si 2p (2p1/2, 2p3/2) peaks at 99.4 eV and (b) Sn 3d peaks at 493 (3d3/2) and 485 (3d5/2) at 200 s of the etch time. (c) and (d) are the Si 2p and Sn 3d peaks for the depth profile of the Si1-xSnx film with respect to etching time.
In order to estimate the alloy composition, the ratios between the integrated intensities of the Sn 3d and Si 2p peaks were plotted as a function of the sputtering time and simulation (Fig. 7(a)). An O peak is observed in the depth profile, however, as a Si-O peak at 104 eV is not observed, it is attributed to a background in the sputtering chamber. The depth profile for Si (2p) and Sn (3d) are plotted in Fig. 6(c) and (d) which indicates the uniform incorporation of the alloy across the film.
Fig. 7. (a) The ratio of Si and Sn from the XPS depth profile of sample D for different etching times. The surface EDX map of sample D shows uniform Sn and Si. (b) shows the map for the Kα peak of Si at 1.739 eV and (c) shows Lα peak of Sn at 3.443 eV.
The Si1-xSnx samples were characterized using SEM/EDX in order to investigate the uniformity of the Si1-xSnx film on the surface as well as the distribution of Si and Sn on the surface. Figure 7(b, c) shows the SEM/EDX map of sample D. The SEM image did not show any granularity for the Si1-xSnx film. The EDX map is taken from Si Kα peak at 1.739 eV and Sn Lα peak at 3.443 eV. The EDX maps were taken randomly on the samples and as it can be seen in Fig. 7 they show that both Si (b) and Sn (c) are uniformly distributed, and no sign of Sn precipitation is observed.
3.4 Spectroscopic ellipsometry
The room temperature spectroscopic ellipsometry was performed to study the optical absorption behavior and film thickness of the samples. Figure 8(a) represents the absorption coefficient of the samples. The absorption cutoff wavelength for the Si reference sample is observed at 1100 nm which is corresponding to the indirect band edge (X). The absorption edge for the Si1-xSnx sample does not undergo a big shift as the difference in the value for Sn and Si is only 0.2 eV for the full range [25]. The incorporation of Sn mainly affects the Γ valley where the difference is 3.9 eV, however, as the wavelength scan range does not pass lower than 900 nm, it cannot reach the direct gap of Si at 3.4 eV (364 nm). Therefore, there could be a slight overestimation of the thicknesses. The thicknesses are estimated to be 92-185 nm for different samples. The refractive index (n) and extinction coefficient (k) of the samples are very similar. Figure 8(b) shows the n and k of sample C where the refractive index is slightly lower than Si reference and the absorption coefficient is slightly higher than Si. They both indicate that the bandgap has shrunk.
Fig. 8. (a) Typical absorption cutoff wavelengths for the Si reference and Si1-xSnx samples (Sample C) measured by spectroscopic ellipsometry. (b) Extinction coefficient and refractive index of sample C.
4. Discussion
The growth of SiSn using commercially available gases such as SiH4 has been a challenge due to the high-temperature requirement to break the Si-Si bond. On the other hand, due to the large lattice mismatch (>20%), unstable α-Sn diamond lattice above 13°C, and low solid solubility of Sn in Si (<0.5%), Sn tends to precipitate out of Si lattice, especially at temperatures above 350°C. This incompatibility of growth temperature and incorporation was the main hindrance in the alloy formation using CVD reactors. However, hydrogen has been shown to have a critical role in bringing down the Si epitaxy temperature to as low as 250 °C [42]. In our low-temperature Si1-xSnx growth, the same method has been adopted in order to use the benefit of growing Si as the host crystal matrix as well as keeping the surface at low enough temperature so that Sn would incorporate without precipitation out of the lattice. It is noteworthy that all of these are possible if the surface is hydrogen-terminated, therefore, sample preparation and hydrogen terminating the sample is of great importance [42]. In addition, keeping the hydrogen on the surface during growth plays an important role. Hydrogen has a 2×1 surface reconstruction at temperatures below 370°C which enables the topmost Si atom to have two hydrogen bonds [43,44]. Increasing the temperature above 407°C would result in almost all 1×1 surface reconstruction and would desorb hydrogen from the surface and result in a monohydride surface [45–47] This would reduce the role of hydrogen on the surface, and for this reason, the growth matrix was designed well below 370°C.
In addition, hydrogen dilution of the precursors (SiH4, SnCl4) would change the plasma environment in comparison to regular PECVD growth [37]. Hydrogen is not just a diluent gas and is an active gas in the growth. Atomic hydrogen produced in plasma enhancement has three major roles in this growth: increasing the surface mobility of ad species, etching the amorphous grown films and promoting layer by layer growth rather than 3D island growth. Si adatom diffusion length decreases significantly at low temperatures, therefore, the adatoms would not be able to find the crystalline sites and would result in a very slow amorphous growth. Atomic hydrogen provides the surface hydrogen coverage and enhances the surface diffusion of SiH3 by generating local heating through hydrogen exchange reactions [48]. In addition, plasma produces other SiHx species [49] through abstracting a hydrogen from the precursor gasses such as generating SiH3 from SiH4 (dissociation energy 90 KCal/mol) which is the main depositing radical with high lifetime and surface mobility and preparing [50]. Such enhancement in surface mobility is through constant abstraction of H from the surface by atomic hydrogen and making crystal site available for SiHx clusters on the surface. In addition to these roles, hydrogen tends to etch amorphous Si growth as the bond energy becomes weaker than the Si in the crystalline site, therefore all the amorphous growth happening as a result of low-temperature growth would be preferentially etched and would not result in low-quality film growth. However, this role results in slower growth rates as well.
In the case of Sn incorporation, SnCl4 dissociates easily at low temperatures [35] and has been incorporated into materials such as Ge below 300 °C, however, plasma environment makes more Sn species available for growth. Dissociation of Cl from SnCl4 requires 55Kcal/mole which is close to Si-Si bond energy (54 Kcal/mole) [51]. Therfore, at a similar temperture, the bonding and dissociation can happen. Moreover, plasma enhancement enables more effective dissociation while also producing SnClx clusters. Chlorine produces HCl in the high-hydrogen dilution environment and would effectively etch Si, however, it plays a positive role in etching away amorphous growth similar to the role of atomic hydrogen. Therefore, plasma enhancement of high hydrogen diluted SiH4 and SnCl4 would be the most favorable condition for low temperature Si1-xSnx growth.
5. Conclusions
In summary, the growth of Si1-xSnx films was achieved using a UHV-PECVD system on Si (001) substrate. The optical and material characterization of the samples shows successful crystalline growth of the Si1-xSnx films. Si1-xSnx alloys with 0.65-3.2% Sn incorporation were grown under different SnCl4 flow rates and plasma powers, however, the analysis shows the possibility of growing Si1-xSnx alloys with higher Sn incorporation using higher SnCl4 flow rates. The structural analysis of the samples shows that the grown films are single crystal and fully pseudomorphic with uniform Sn incorporation in Si lattice. The current results show that achieving device quality higher-Sn contents SiSn (and SiGeSn) films are feasible using PECVD technique. High quality SiSn films are essential in fabricating photonic devices such as photodetectors, modulators and lasers that are fabricated with single or multilayers of SiSn, GeSn, and SiGeSn.
Funding
Air Force Office of Scientific Research (FA9550-19-1-0341); Office of the Provost, University of Arkansas.
Acknowledgments
The authors would also like to thank Zheng Ju and Prof. Yong-Hang Zhang at Arizona State University as well as Institute for Nano Science and Engineering at the University of Arkansas for assistance in XRD measurements and Dr. Fumiya Watanabe at the Center for Integrative Nanotechnology Sciences at the University of Arkansas-Little Rock for XPS/EDX characterizations.
Disclosures
The authors declare that there are no conflicts of interest related to this article.
References
1. A. H. Atabaki, S. Moazeni, F. Pavanello, H. Gevorgyan, J. Notaros, L. Alloatti, M. T. Wade, C. Sun, S. A. Kruger, and H. Meng, “Integrating photonics with silicon nanoelectronics for the next generation of systems on a chip,” Nature 556(7701), 349–354 (2018). [CrossRef]
2. K. Giewont, K. Nummy, F. A. Anderson, J. Ayala, T. Barwicz, Y. Bian, K. K. Dezfulian, D. M. Gill, T. Houghton, and S. Hu, “300-mm monolithic silicon photonics foundry technology,” IEEE J. Sel. Top. Quantum Electron. 25(5), 1–11 (2019). [CrossRef]
3. A. Mosleh, S. A. Ghetmiri, B. R. Conley, M. Hawkridge, M. Benamara, A. Nazzal, J. Tolle, S. Yu, and H. A. Naseem, “Material Characterization of Ge1−xSnx Alloys Grown by a Commercial CVD System for Optoelectronic Device Applications,” J. Electron. Mater. 43(4), 938–946 (2014). [CrossRef]
4. C. Xu, R. T. Beeler, L. Jiang, J. D. Gallagher, R. Favaro, J. Menéndez, and J. Kouvetakis, “Synthesis and optical properties of Sn-rich Ge1−x−ySixSny materials and devices,” Thin Solid Films 557, 177–182 (2014). [CrossRef]
5. S. Wirths, D. Buca, G. Mussler, A. Tiedemann, B. Holländer, P. Bernardy, T. Stoica, D. Grützmacher, and S. Mantl, “Reduced Pressure CVD Growth of Ge and Ge1−xSnxAlloys,” ECS J. Solid State Sci. Technol. 2(5), N99–N102 (2013). [CrossRef]
6. W. Du, Y. Zhou, S. A. Ghetmiri, A. Mosleh, B. R. Conley, A. Nazzal, R. A. Soref, G. Sun, J. Tolle, and J. Margetis, “Room-temperature electroluminescence from Ge/Ge1−xSnx/Ge diodes on Si substrates,” Appl. Phys. Lett. 104(24), 241110 (2014). [CrossRef]
7. Y. Zhou, W. Dou, W. Du, T. Pham, S. A. Ghetmiri, S. Al-Kabi, A. Mosleh, M. Alher, J. Margetis, and J. Tolle, “Systematic study of GeSn heterostructure-based light-emitting diodes towards mid-infrared applications,” J. Appl. Phys. 120(2), 023102 (2016). [CrossRef]
8. T. Pham, W. Du, H. Tran, J. Margetis, J. Tolle, G. Sun, R. A. Soref, H. A. Naseem, B. Li, and S. Yu, “Systematic study of Si-based GeSn photodiodes with 2.6 µm detector cutoff for short-wave infrared detection,” Opt. Express 24(5), 4519–4531 (2016). [CrossRef]
9. T. Pham, W. Du, B. Conley, J. Margetis, G. Sun, R. Soref, J. Tolle, B. Li, and S. Yu, “Si-based Ge0. 9Sn0. 1 photodetector with peak responsivity of 2.85 A/W and longwave cutoff at 2.4 μm,” Electron. Lett. 51(11), 854–856 (2015). [CrossRef]
10. J. Werner, M. Oehme, M. Schmid, M. Kaschel, A. Schirmer, E. Kasper, and J. Schulze, “Germanium-tin pin photodetectors integrated on silicon grown by molecular beam epitaxy,” Appl. Phys. Lett. 98(6), 061108 (2011). [CrossRef]
11. J. Margetis, S. Al-Kabi, W. Du, W. Dou, Y. Zhou, T. Pham, P. Grant, S. Ghetmiri, A. Mosleh, and B. Li, “Si-based GeSn lasers with wavelength coverage of 2 to 3 μm and operating temperatures up to 180 K,” ACS Photonics 5(3), 827–833 (2018). [CrossRef]
12. S. Al-Kabi, S. A. Ghetmiri, J. Margetis, T. Pham, Y. Zhou, B. Collier, R. Quinde, W. Du, A. Mosleh, J. Liu, R. A. Soref, J. Tolle, B. Li, M. Mortazavi, H. A. Naseem, and S. Yu, “An optically pumped 2.5 µm GeSn laser on Si operating at 110 K,” Appl. Phys. Lett. 109(17), 171105 (2016). [CrossRef]
13. W. Dou, Y. Zhou, J. Margetis, S. A. Ghetmiri, S. Al-Kabi, W. Du, J. Liu, G. Sun, R. A. Soref, and J. Tolle, “Optically pumped lasing at 3 μm from compositionally graded GeSn with tin up to 22.3%,” Opt. Lett. 43(19), 4558–4561 (2018). [CrossRef]
14. S. Wirths, R. Geiger, N. von den Driesch, G. Mussler, T. Stoica, S. Mantl, Z. Ikonic, M. Luysberg, S. Chiussi, and J. Hartmann, “Lasing in direct-bandgap GeSn alloy grown on Si,” Nat. Photonics 9(2), 88–92 (2015). [CrossRef]
15. M. Akie, T. Fujisawa, T. Sato, M. Arai, and K. Saitoh, “GeSn/SiGeSn multiple-quantum-well electroabsorption modulator with taper coupler for mid-infrared Ge-on-Si platform,” IEEE J. Sel. Top. Quantum Electron. 24(6), 1–8 (2018). [CrossRef]
16. S. A. Ghetmiri, Y. Zhou, J. Margetis, S. Al-Kabi, W. Dou, A. Mosleh, W. Du, A. Kuchuk, J. Liu, and G. Sun, “Study of a SiGeSn/GeSn/SiGeSn structure toward direct bandgap type-I quantum well for all group-IV optoelectronics,” Opt. Lett. 42(3), 387–390 (2017). [CrossRef]
17. H. Wang, G. Han, Y. Wang, Y. Peng, Y. Liu, C. Zhang, J. Zhang, S. Hu, and Y. Hao, “Theoretical calculation of performance enhancement in lattice-matched SiGeSn/GeSn p-channel tunneling field-effect transistor with type-II staggered tunneling junction,” Jpn. J. Appl. Phys. 55(4S), 04ED13 (2016). [CrossRef]
18. G. Sun and S. Yu, “The SiGeSn approach towards Si-based lasers,” Solid-State Electron. 83, 76–81 (2013). [CrossRef]
19. B. R. Conley, H. Naseem, G. Sun, P. Sharps, and S. Yu, "High efficiency MJ solar cells and TPV using SiGeSn materials,” in Photovoltaic Specialists Conference (PVSC), 2012 38th IEEE (2012).
20. H. Baker and H. Okamoto, “ASM handbook,” Alloy phase diagrams 3, 2 (1992).
21. J. Vanjaria, A. C. Arjunan, T. Salagaj, G. S. Tompa, and H. Yu, “PECVD Growth of Composition Graded SiGeSn Thin Films as Novel Approach to Limit Tin Segregation,” ECS J. Solid State Sci. Technol. 9(3), 034009 (2020). [CrossRef]
22. M. A. Alher, A. Mosleh, L. Cousar, W. Dou, P. Grant, S. A. Ghetmiri, S. Al-Kabi, W. Du, M. Benamara, and B. Li, “CMOS Compatible Growth of High Quality Ge, SiGe and SiGeSn for Photonic Device Applications,” ECS Trans. 69(5), 269–278 (2015). [CrossRef]
23. S. Wirths, D. Buca, Z. Ikonic, P. Harrison, A. Tiedemann, B. Holländer, T. Stoica, G. Mussler, U. Breuer, and J. Hartmann, “SiGeSn growth studies using reduced pressure chemical vapor deposition towards optoelectronic applications,” Thin Solid Films 557, 183–187 (2014). [CrossRef]
24. J. Tolle, A. Chizmeshya, Y. Fang, J. Kouvetakis, V. D’Costa, C. Hu, J. Menendez, and I. Tsong, “Low temperature chemical vapor deposition of Si-based compounds via SiHSiHSiH: Metastable SiSn∕ GeSn∕ Si (100) heteroepitaxial structures,” Appl. Phys. Lett. 89(23), 231924 (2006). [CrossRef]
25. P. Moontragoon, R. Soref, and Z. Ikonic, “The direct and indirect bandgaps of unstrained SixG1−xySy and their photonic device applications,” J. Appl. Phys. 112(7), 073106 (2012). [CrossRef]
26. K. S. Min and H. A. Atwater, “Ultrathin pseudomorphic Sn/Si and SnxSi1−x/Si heterostructures,” Appl. Phys. Lett. 72(15), 1884–1886 (1998). [CrossRef]
27. V. Švrček, D. Mariotti, R. Blackley, W. Zhou, T. Nagai, K. Matsubara, and M. Kondo, “Semiconducting quantum confined silicon–tin alloyed nanocrystals prepared by ns pulsed laser ablation in water,” Nanoscale 5(15), 6725–6730 (2013). [CrossRef]
28. V. Timofeev, V. Mashanov, A. Nikiforov, I. Skvortsov, T. Gavrilova, D. Gulyaev, A. Gutakovskii, and I. Chetyrin, “Effect of Sn for the dislocation-free SiSn nanostructure formation on the vapor-liquid-crystal mechanism,” AIP Adv. 10(1), 015309 (2020). [CrossRef]
29. M. Kurosawa, M. Kato, K. Takahashi, O. Nakatsuka, and S. Zaima, “Self-organized lattice-matched epitaxy of Si1−xSnx alloys on (001)-oriented Si, Ge, and InP substrates,” Appl. Phys. Lett. 111(19), 192106 (2017). [CrossRef]
30. M. Kurosawa, Y. Inaishi, R. Tange, M. Sakashita, O. Nakatsuka, and S. Zaima, “Synthesis of heavily Ga-doped Si1−xSnx/Si heterostructures and their valence-band-offset determination,” Jpn. J. Appl. Phys. 58(SA), SAAD02 (2019). [CrossRef]
31. M. Kurosawa, O. Nakatsuka, and S. Zaima, “Growth and Applications of Si1−xSnx Thin Films,” ECS Trans. 80(4), 253–258 (2017). [CrossRef]
32. J. Tan, K. Chen, and L. Tang, “Out-of-plane spontaneous polarization and superior photoelectricity in two-dimensional SiSn,” J. Phys.: Condens. Matter 32(6), 065003 (2020). [CrossRef]
33. X. Jiang, M. Li, Y. Ming, J. Wang, and W. Li, "The Structure and Properties of Sn/SiSn-nanodisperse Alloy Thin Films,” in IOP Conference Series: Materials Science and EngineeringAnonymous (IOP Publishing, 2019).
34. L. Scheffler, S. Roesgaard, J. L. Hansen, A. Nylandsted Larsen, and B. Julsgaard, “Tin-based donors in SiSn alloys,” J. Appl. Phys. 126(3), 035702 (2019). [CrossRef]
35. A. Mosleh, M. A. Alher, L. C. Cousar, W. Du, S. A. Ghetmiri, T. Pham, J. M. Grant, G. Sun, R. A. Soref, and B. Li, “Direct Growth of Ge1-xSnx Films on Si Using a Cold-Wall Ultra-High Vacuum Chemical-Vapor-Deposition System,” Front. Mater. 2, 30 (2015). [CrossRef]
36. A. Mosleh, "Investigation of Growth Mechanism and Role of H2 in Very Low Temperature Si Epitaxy,” in 2014 ECS and SMEQ Joint International Meeting (October 5-10, 2014)Anonymous (Ecs, 2014).
37. W. Dou, B. Alharthi, P. C. Grant, J. M. Grant, A. Mosleh, H. Tran, W. Du, M. Mortazavi, B. Li, and H. Naseem, “Crystalline GeSn growth by plasma enhanced chemical vapor deposition,” Opt. Mater. Express 8(10), 3220–3229 (2018). [CrossRef]
38. M. Kurosawa, M. Kato, T. Yamaha, N. Taoka, O. Nakatsuka, and S. Zaima, “Near-infrared light absorption by polycrystalline SiSn alloys grown on insulating layers,” Appl. Phys. Lett. 106(17), 171908 (2015). [CrossRef]
39. R. Beeler, R. Roucka, A. Chizmeshya, J. Kouvetakis, and J. Menéndez, “Nonlinear structure-composition relationships in the Ge1−ySny/Si (100)(y< 0.15) system,” Phys. Rev. B 84(3), 035204 (2011). [CrossRef]
40. P. M. Yakibchuk, O. V. Bovgyra, M. V. Kovalenko, and I. V. Kutsa, “Electronic Structure Of Si1−xSnx Disordered Solid Solutions,” Zh. Fiz. Dosl. 23(2), 2703 (2019). [CrossRef]
41. W. Dou, M. Benamara, A. Mosleh, J. Margetis, P. Grant, Y. Zhou, S. Al-Kabi, W. Du, J. Tolle, and B. Li, “Investigation of GeSn strain relaxation and spontaneous composition gradient for low-defect and high-Sn alloy growth,” Sci. Rep. 8(1), 1–11 (2018). [CrossRef]
42. A. Mosleh, S. A. Ghetmiri, B. R. Conley, H. H. Abu-Safe, M. Benamara, Z. Waqar, S. El-Ghazaly, S. Yu, and H. A. Naseem, “Investigation of Growth Mechanism and Role of H2 in Very Low Temperature Si Epitaxy,” ECS Trans. 64(6), 967–975 (2014). [CrossRef]
43. D. Eaglesham, G. Higashi, and M. Cerullo, “370° C clean for Si molecular beam epitaxy using a HF dip,” Appl. Phys. Lett. 59(6), 685–687 (1991). [CrossRef]
44. M. V. R. Murty and H. A. Atwater, “Empirical interatomic potential for Si-H interactions,” Phys. Rev. B 51(8), 4889–4893 (1995). [CrossRef]
45. P. Gupta, V. Colvin, and S. George, “Hydrogen desorption kinetics from monohydride and dihydride species on silicon surfaces,” Phys. Rev. B 37(14), 8234–8243 (1988). [CrossRef]
46. M. C. Flowers, N. B. H. Jonathan, Y. Liu, and A. Morris, “Temperature programmed desorption of molecular hydrogen from a Si (100)-2×1 surface: Theory and experiment,” J. Chem. Phys. 99(9), 7038–7048 (1993). [CrossRef]
47. Y. Chabal and K. Raghavachari, “New ordered structure for the H-saturated Si (100) surface: the (3× 1) phase,” Phys. Rev. Lett. 54(10), 1055–1058 (1985). [CrossRef]
48. A. Matsuda, “Formation kinetics and control of microcrystallite in μc-Si: H from glow discharge plasma,” J. Non-Cryst. Solids 59-60, 767–774 (1983). [CrossRef]
49. A. Matsuda, “Plasma CVD Processes for Thin Film Silicon Solar Cells,” Advanced Plasma Technology 197-210.
50. J. H. Purnell and R. Walsh, “Some comments on kinetics and mechanism in the pyrolysis of monosilane,” Chem. Phys. Lett. 110(3), 330–334 (1984). [CrossRef]
51. G. Oldershaw and K. Robinson, “Flash photolysis of SnCl2 and SnCl4. Dissociation energies of SnCl2 and SnCl,” Journal of the Chemical Society A: Inorganic, Physical, Theoretical2963–2967 (1971). [CrossRef]
