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Abstract
A series of SiGeSn alloy samples with various Si and Sn compositions and thicknesses were grown on Ge-buffered Si substrates. The growth was conducted by using low-cost commercially available silane and germane precursors in a standard industrial reduced pressure chemical vapor deposition reactor. Si and Sn compositional- and film thickness-dependent material and optical properties have been characterized using X-ray diffraction (XRD), Raman, photoluminescence (PL), and ellipsometry spectroscopies. Moreover, thermal stability in harsh growth environment, such as in subsequent III-V growth, was studied for future multi-junction solar cell applications. In situ rapid thermal annealing at 650°C was conducted to investigate the enhanced material quality and direct bandgap emission, which were confirmed by XRD, transmission electron microscopy, Raman, and PL measurements.
© 2017 Optical Society of America
1. Introduction
Group-IV SiGeSn alloys have attracted research interest worldwide. This is because of their complementary metal oxide semiconductors (CMOS) compatibility [1,2] and unique optical properties that could be adopted in many device applications [3,4]. From growth perspective, SiGeSn alloys feature the capability of separately engineering the bandgap energy and lattice constant by tuning the Si and Sn compositions. For the Si-based light source, which has long been desired, it is noted that the direct bandgap emission can be significantly enhanced with higher Sn-content [5,6]. To date, GeSn alloys (i.e., Si composition of 0 at.%) show high-quality growth and direct bandgap material [7], leading to the successful demonstration of optically pumped lasers [8,9], light emitting diodes [10], and photodetectors [11,12].
The unique optical properties of SiGeSn make it of great potential in Si photonics applications. For example, the compositions of Si and Sn can be chosen in a way that the SiGeSn/GeSn/SiGeSn double heterostructures or quantum wells not only form type-I alignment but also are lattice matched. The SiGeSn alloy with 1.0 eV direct bandgap energy can also be used in multi-junction (MJ) solar cell systems comprising of III-V junctions, thereby leading to a much cheaper option of using Si substrate [13]. For light emission and detection applications, the wavelength coverage could extend well beyond that of Ge into broad infrared (IR) range by tuning the Si and Sn compositions. Currently, the growth of high-quality SiGeSn alloys satisfying the device-level criteria is a formidable challenge, which is mainly due to the following two factors: the low thermodynamic solubility of α-Sn in both Si and Ge [14]; and the large lattice mismatch between Ge, Sn, and Si [15,16]. Therefore, the low-temperature growth techniques have been developed to alleviate the growth difficulties. The SiGeSn material growth by molecular beam epitaxy, by magnetron sputtering [17–19], and by chemical vapor deposition (CVD) [14,20,21] have been reported. However, most of the CVD reported works were done using high-cost higher order GexH2x + 2 and SixH2x + 2 hydrides which are not favorable in industry. Thus, the CVD growth using cost effective precursors, such as germane (GeH4) and silane (SiH4), is desired and has not been fully developed. Moreover, the systematic compositional study of SiGeSn material and optical characteristics have not been reported yet.
In this work, we present a systematic study on material and optical properties of SiGeSn alloys that were grown on Ge buffered Si substrate via reduced pressure CVD (RP-CVD) using SiH4, GeH4, and SnCl4 as Si, Ge, and Sn precursors, respectively. The characterizations revealed the Si and Sn compositional- and film thickness-dependent material and optical properties. Moreover, to investigate the thermal stability, the annealing study of a selected SiGeSn alloy was performed.
2. Experimental methods
Epitaxial SixGe1-x-ySny layers were grown Ge-buffered Si(100) substrate using a RPCVD ASM Epsilon 2000 reactor. Germanium buffer layers (approximately 700 nm thick) were grown in situ prior to SiGeSn growth using 10% GeH4 in purified H2 carrier gas. Prior to growth of the Ge buffer layer, the silicon substrate is heated to 1060 ºC at 20 torr to remove the native oxide. The buffer layers were grown by a two-step growth method. First, a 150-nm seed layer was grown at <400°C in H2 carrier at a GeH4 partial pressure of 0.2 Torr, then the temperature was increased to 600°C. Once the temperature has stabilized at 600°C, the remaining ~500 nm was grown and a post-growth in situ anneal was done at >800°C. The chamber was then cooled down to <400°C in H2 and then the GeH4, SiH4, and SnCl4 precursors were introduced into the chamber to initiate the SiGeSn film growth. SnCl4 is a liquid at room temperature and must be delivered using a bubbler held at room temperature in which H2 gas is metered to it to increase or decrease to the desired SnCl4 mass flow rate. The SnCl4 mass flow rate was measured up-stream of the bubbler by a piezoelectric acoustic sensor. This signal was fed back to H2 bubbler mass flow controller in a dynamic control loop to continually monitor and adjust the SnCl4 relative to the desired set-point. The SiH4/GeH4 flow ratios were varied between 0.15 and 0.6 at a fixed SnCl4/GeH4 mass flow ratio of 0.01.
To determine the compositions of Si and Sn in the alloy, the high-resolution x-ray diffraction (HRXRD) and Rutherford backscattering spectroscopy (RBS) techniques were employed. The Sn composition was first obtained from RBS and then the Si composition was extracted via XRD based on the Sn composition. Figure 1 shows typical Si (001) channeling and random RBS spectra of a SiGeSn sample with a 50-nm film thickness and 7.3 at.% and 5.5 at.% incorporation of Si and Sn, respectively. RUMP software was used to fit the peaks.
Three sets of SiGeSn samples were investigated in this study: i) three samples (A, B, and C) with the same Si and Sn compositions of 7.3 and 5.5 at.% but different film thicknesses ranging from 50 to 200 nm; ii) three samples (A, D, and E) with the same Sn composition of 5.5 at.% and film thickness of 50 nm but different Si compositions ranging from 7.3 to 10.0 at.%; and iii) two samples (F and G) with relatively high Si incorporations of 12.0 and 19.0 at.%. The detailed sample information is listed in Table 1.
Table 1. Summary of SiGeSn samples
Study of the structural properties was conducted using a Phillips X'pert PRO diffractometer equipped with a Ge (220) monochromator HRXRD. A TITAN transmission electron microscope (TEM) has been used to study crystal orientation and defects. The crystallinity and Sn incorporation were studied with Raman spectroscopy equipped with a 5 mW 632-nm laser. Temperature-dependent photoluminescence (PL) measurement was performed by using a 532-nm continuous wave laser with the power of 500 mW as pumping source. The PL emissions were sent to an iHR spectrometer and then were collected by a liquid-nitrogen-cooled InGaAs detector. A Variable-Angle Spectroscopic Ellipsometry (Woolam Model VASE32) was used to investigate the spectral absorption coefficient in the range of 0.496-4.768 eV (260-2500 nm) with a resolution of 10 nm at three angles of incidence (65°, 70°, and 75°). The data fitting process was performed using the built-in WVASE32 software. The detailed data fitting procedure can be found elsewhere [22].
To investigate the thermal stability of the SiGeSn alloy for potential MJ solar cells applications, in situ annealing was performed with sample F at 650°C in H2 ambient for 2 minutes at atmospheric pressure. The material and optical properties were characterized for as-grown and after annealing samples, and the results were compared to illustrate the effects of thermal treatment.
3. Results and discussion
3.1 Material characterization
Figure 2(a) shows the 2θ-ω scans from symmetric (004) planes of samples A, B, and C that have the same Si and Sn composition (7.3 at.% and 5.5 at.%) but with different film thickness. The curves were aligned with Si substrate peaks at 69°. The peak from the Ge buffer is at 66°, while the peaks at angles below 66° are assigned to the SiGeSn. The SiGeSn peak intensity increases by one order of magnitude as the film thickness increases from 50 to 200 nm. Moreover, Ge and SiGeSn peaks are slightly shifted towards higher angles as the thickness increases due to strain relaxation of the material. Figure 2(a) inset shows the typical reciprocal space map (RSM) contour plot of sample C, which indicates that the SiGeSn layer was grown almost fully pseudomorphic to Ge buffer as they feature almost the same in-plane lattice constant (a||). Being below the critical thickness for the Si0.073Ge0.872Sn0.055 layers, these films predominantly show pseudomorphicity even for the 200-nm thick film in agreement with our theoretical calculations although a slight relaxation in the compressive strain was observed as film thickness increased. The in-plane and out-of-plane (a⊥) lattice constants for each sample were extracted from RSM contour plots, as summarized in Table 1.
Fig. 2 2θ-ω scan from (004) plane for three sets of samples. (a) Samples A, B, and C feature the same Si and Sn compositions but different film thicknesses. Inset: typical RSM of sample C. (b) Samples A, D, and E feature the same Sn composition and film thickness but different Si compositions. Inset: typical RSM of sample A. (c) Samples F and G feature relatively high Si incorporation. (d) RSMs of samples F and G.
The 2θ-ω scans of SixGe0.945-xSn0.055 films with the same Sn composition and film thickness but different Si compositions (samples A, D, and E) are presented in Fig. 2(b). Increasing Si incorporation from 7.3 at.% to 10 at.% in SixGe0.945-xSn0.055 shifts the SiGeSn peaks toward the Ge peaks as expected. The thickness fringes that are clearly discernible below the SiGeSn peaks due to interference between the reflected X-ray beams from the film interfaces, demonstrate smooth morphology in these high-quality material films [23]. Figure 2(b) inset shows the RSM contour plot of sample A. The pseudomorphic growth is clearly indicated.
Figure 2(c) shows the 2θ-ω scans of SiGeSn samples with relatively high Si incorporation of 12 at.% (sample F) and 19 at.% (sample G), the corresponding Sn compositions are 9.0 at.% and 2.7 at.%, respectively. Due to the higher incorporated Si, the SiGeSn peak shifts considerably towards larger angle, resulting in partial overlap of the SiGeSn and Ge peaks. The broadened peak at 66° in Fig. 2(c) indicates the existence of two overlapped peaks. For sample F, the lower angle shoulder was attributed to the SiGeSn peak. While for sample G, since more Si and less Sn are incorporated compared to sample F, the SiGeSn peak features more overlap with the Ge peak, indicating it is nearly lattice-matched to the Ge buffer. The RSM contour plots of samples F and G further confirmed the existence of SiGeSn. As shown in Fig. 2(d), the contours of SiGeSn and Ge are mostly overlapped, leading to the broadened contour plots.
Raman spectroscopy was used to investigate the SiGeSn film crystallinity. Figure 3 shows normalized Raman spectra of the samples A, D, and E at room temperature. The spectra were stacked for clarity. For each curve, the Ge-Ge longitudinal optical (LO) peak was observed at slightly less than 300 cm−1 close to the standard Raman shift of the reference Ge. The shift of Ge-Ge LO peak in SiGeSn samples is mainly due to the Sn and Si incorporation in Ge lattice which changes the average bond size and strength of Ge-Ge lattice. Another peak with relatively lower intensity was observed at ~385 cm−1, which corresponds to the Si-Ge bond. As Si incorporation increases, the peak shifts towards larger Raman shift, i.e., towards Si-Si bond at 520 cm−1 (not shown here). The inset of Fig. 3 shows a magnified illustration of the expected Ge-Sn mode positions. For sample A, the hump at ∼262 cm−1 is associated with the Ge-Sn bond [24].
Fig. 3 Raman spectra of samples A, D, and E. The inset shows a zoomed-in plot of the Ge-Sn modes region.
3.2 Optical characterization
Photoluminescence measurement was conducted to investigate the optical properties of the samples. Figure 4(a) shows the room temperature PL spectra of samples A, B, and C. Two peaks and one shoulder can be clearly observed, and interpreted based on bandgap energy calculation [25] as following: i) the peak at ~0.79 eV was assigned to the overlap of direct bandgap emissions from Ge and SiGeSn; ii) the peak at ~0.67 eV was attributed to the indirect bandgap emission from Ge buffer; iii) the shoulder at ~0.63 eV was associated with the indirect bandgap emission from SiGeSn. For sample A, since the SiGeSn layer features wider bandgap compared to the Ge buffer, the SiGeSn could act as barrier and therefore most photo-generated carriers remain confined in the Ge buffer resulting in the Ge emission dominating the PL. For samples B and C, due to the penetration depth of 532-nm pump laser (< 100 nm) being less than the film thickness, the photon absorption and recombination occur mainly in the SiGeSn layer, leading to the emission from SiGeSn contributing equally to the PL (considering that the carrier confinement still affects the carrier re-distribution). The PL spectra of samples A, D, and E are plotted in Fig. 4(b). For each spectrum, the Ge emission dominates the PL. As Si content increases, the ratio of emissions from SiGeSn over Ge decreases, which is due to the fact that indirect bandgap dominates more in the material with higher Si incorporation. Moreover, the shoulder at ~0.62 eV, associated with the indirect bandgap emission from SiGeSn films, shifts towards higher energy as Si composition increases. This is expected because of the increased indirect bandgap energy with higher Si incorporation in the alloy.
Fig. 4 (a) PL spectra as the thickness varies in samples A, B, and C. (b) PL measurements as the Si content increases in samples A, D, and E. (c) PL spectra of samples with high Si content. (d) Temperature dependent PL spectra of sample G.
Figure 4(c) shows room temperature PL spectra of samples F and G. For sample F, since the incorporation of 9 at.% Sn overcompensates the 12 at.% Si incorporation-induced bandgap indirectness, a narrower direct bandgap of SiGeSn compared to that of Ge was obtained, leading to the direct peak red-shift of SiGeSn. For sample G, two main peaks with a strong higher energy peak at ~0.79 eV and a lower energy peak at ~0.69 eV were observed, which were assigned to the overlapped direct bandgap emissions from Ge and SiGeSn, and overlapped indirect bandgap emissions, respectively. The direct bandgap energy separation between Ge and SiGeSn is too small hence their direct transitions cannot be identified from PL spectra. While for the indirect bandgap transition, the broadened peak indicates that both SiGeSn and Ge emissions contribute to PL.
The temperature-dependent PL spectra of sample G are shown in Fig. 4(d). As temperature decreases, the PL peaks shift towards higher energies due to bandgap increase as expected. Moreover, the overall PL intensity decreases as the temperature decreases with the direct peaks dropping more rapidly than the indirect peaks, resulting in the direct peaks dominating the PL at 300 K whereas the indirect peaks dominate the PL at the temperatures below 100 K. This PL spectra behavior indicates that sample G is a typical indirect bandgap material. In fact, adding Sn into Ge would engineer the bandgap towards direct bandgap while incorporation of Si shifts the bandgap to the opposite direction. Therefore, the bandgap property can be tuned by independently controlling the Si and Sn compositions. The broad peaks that appear at ∼0.6 eV are attributed to radiative recombination from states introduced by the impurities and other defects in the Si0.19Ge0.783Sn0.027 film [26]. The line-widths of the direct and indirect transitions in the temperature-dependent PL spectra of sample G were extracted after Gaussian fitting was performed. The results are plotted in Fig. 5(a). As expected, for the indirect peaks the line-widths decrease dramatically as the temperature decreases. However, the line-widths of the direct peaks become slightly wider with decreasing temperatures. This behavior confirms that sample G is indirect bandgap material. The peak positions of both the direct and indirect peaks were extracted and fitted using the Varshni equation:
where ESiGeSn(0) is the bandgap energy at 0K, α and β are fitting parameters. The results are presented in Fig. 5(b) and the parameters are given in Table 2.Fig. 5 Temperature-dependent of the PL spectra of sample G as a function of (a) Line-widths and (b) PL peak positions. The solid lines in (b) represent Varshni fittings.
Table 2. Varshni’s equation fitting parameters of the temperature-dependent PL spectra of sample G
Spectroscopic ellipsometry was used to investigate the absorption properties of the samples. Figure 6(a) shows the spectral absorption coefficient curves of samples A, B and C. Since the increase in film thickness relaxes the material, and consequently reduces the bandgap energy, the absorption cut-off edge shifts from 0.78 to 0.75 eV as the film thickness increases, which was assigned to direct bandgap absorption. The indirect bandgap absorption features small value and hence cannot be extracted accurately due to the equipment limitations. Figure 6(b) shows the absorption coefficient data of samples A, D, E, F, and G. The absorption coefficient curves shift towards longer wavelengths when increasing the Sn incorporation, while the curves shift towards shorter wavelengths as the Si composition increases. As a result, the absorption cut-off edges are determined by the overall Si and Sn incorporations.
Fig. 6 Spectral absorption coefficient curves of (a) samples A, B, and C. (b) Samples A, D, E, F, and G.
3.3 Annealing study
Since the SiGeSn alloy is a viable candidate for a solar cell junction within a III-V based high efficiency MJ junction solar PV system, annealing study was conducted to investigate the material property under thermal treatment. Sample F was selected as it features relatively high Si and Sn compositions, and the annealing temperature of 650°C was set to investigate the thermal stability under a typical MOCVD deposition temperature for III-V materials. Figure 7(a) shows 2θ-ω scans of sample F as-grown and after annealing. It can be seen that the SiGeSn peak features reduced peak line-width and increased peak intensity after annealing, indicating the improved material quality. The RSM contour plots of sample F as-grown and after annealing are shown in Fig. 7(b). Unlike RSM for the as-grown sample, the SiGeSn contour after annealing can be clearly resolved from broadened Ge plot, confirming the improved material quality. Moreover, the out-of-plane lattice constant of the annealed sample increases while the in-plane lattice constant remains almost the same, suggesting the almost unchanged pseudomorphicity of the material.
The material quality was further investigated by the cross-sectional TEM images of sample F as-grown and after annealing. Figure 8(a) shows TEM image of sample F as-grown, which clearly illustrates black and white thickness fringes. The arrows at the Ge buffer/Si interface indicate the stacking fault defects due to dissociation of dislocations at the [111] plane. Threading dislocations that propagate through the Ge buffer are resulting from the large lattice mismatch in the Ge/Si interface. The high-resolution image of SiGeSn/Ge interface as-grown is shown in Fig. 8(b), which illustrates a few misfit dislocations and provides evidence of lattice matched growth. The TEM image of sample F after annealing is shown in Fig. 8(c). The smooth cross-section with no observed threading dislocations in SiGeSn/Ge/Si interfaces implies improved material quality as a result of the annealing process. The high-resolution TEM image of SiGeSn/Ge interface shown in Fig. 8(d) illustrates periodic defects with ~250 nm spacing on average that are interpreted as Lomer dislocations. Such dislocations appear in the SiGeSn/Ge interface to accommodate the increased vertical lattice constant as the annealed SiGeSn film tends to approach relaxation.
Fig. 8 Cross sectional TEM of sample F. As-grown in (a) and (b). After rapid thermal annealing in (c) and (d).
Optical characterization including Raman and PL spectroscopy were also performed to study the optical properties of the material under thermal treatment. Raman spectroscopy shown in Fig. 9(a) indicates that after annealing the Ge-Ge LO mode is slightly lower than that of the as-grown sample, while the peak intensity and line-width almost remain the same. However, the Ge-Sn and Si-Ge modes show improved crystallinity as a result of the rapid thermal annealing. Figure 9(b) shows the PL spectra of as-grown and annealed samples. Compared to the PL spectrum of the as-grown sample, the PL peaks of the annealed sample shows a blue-shift, which is mainly due to the slightly relaxed material in high temperature environment. The improved material quality under thermal treatment suggests that the growth of SiGeSn alloy can be subjected to higher temperatures during the MJ junction solar cell growth without much detrimental effects, since the material quality could remain the same or even be improved slightly.
Fig. 9 Sample F as-grown and after annealing. (a) Stacked Raman spectra. The inset shows the Ge-Sn modes. (b) PL spectra.
4. Conclusion
5. In summary, SiGeSn alloys with various film thicknesses and varied Si and Sn compositions were grown on Ge buffered Si substrates via a standard industrial RP-CVD reactor. The XRD 2θ-ω scans showed the Si and Sn compositional- and film thickness-dependent material properties. The RSM contour plots indicated pseudomorphic growth of SiGeSn films. The PL spectra exhibited direct and indirect emissions from SiGeSn films, which were overlapped with emission peaks from Ge buffer. The spectral absorption coefficient of the alloys was studied by spectroscopic ellipsometry, which showed sharp absorption cut-off edge for each sample. Moreover, the annealing study was performed to investigate the material property under thermal treatment, which was aimed to study the thermal stability for MJ solar cell applications. The results showed an improved material quality that suggests SiGeSn is a versatile candidate for being part of a well-designed MJ of solar cell stack.
Funding
Air Force Office of Scientific Research (AFOSR) (FA9550-14-1-0205); National Aeronautics and Space Administration Established Program to Stimulate Competitive Research (NASA EPSCoR) (NNX15AN18A).
Acknowledgments
The authors would like to thank Dr. Mourad Benamara and Dr. Andrian Kuchuk from the Institute of Nano Science and Engineering at the University of Arkansas for TEM imaging and XRD measurements.
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