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Abstract
Single crystalline GeSn growth on Si substrate was successfully demonstrated by using plasma enhanced chemical vapor deposition (PE-CVD) with commercially available GeH4 and SnCl4 precursors. Using the plasma enhancement technique, low temperature growth at 350°C for GeSn epitaxy on Si substrate was achieved with the growth rate of 51.4 nm/min and Sn content up to 6%. The relaxed GeSn films with 1 µm thickness were able to be grown despite of the huge lattice mismatch between GeSn and Si. Structural and optical characterizations were conducted to study the film properties. The demonstrated plasma enhancement growth showed its effectiveness to enhance the Sn incorporation of the crystalline GeSn at low temperature and to maintain the high growth rate.
© 2018 Optical Society of America under the terms of the OSA Open Access Publishing Agreement
1. Introduction
Scientific community has shown great interest in Group-IV GeSn alloys in the past few decades. GeSn alloys have demonstrated tunable energy bandgap covering broad near- and mid-infrared wavelength, as well as true direct bandgap with sufficient Sn content. The GeSn epitaxy process is compatible with current complementary metal–oxide–semiconductor (CMOS) process, suggesting that the devices based on GeSn could be monolithically integrated on Si substrate [1–4]. The GeSn growth is challenging due to the large lattice mismatch between GeSn and Si (>4.2%), and the low solid solubility (<0.5%) of Sn in Ge. Therefore, the growth techniques have used Ge buffer layer and non-equilibrium conditions growth to address these challenges. The GeSn growth using chemical vapor deposition (CVD) technique at low temperature has been investigated for more than a decade. Currently, higher order Ge hydrides, such as Ge2H6 and Ge3H8, show attractive low temperature growth results due to their higher decomposition rate at lower temperatures [5–10].Using such growth approaches the GeSn lasers with Sn compositions of 12.6% [5, 8] and 16% [7] were reported. However, GeH4 is preferred from an industrial manufacturing perspective due to its much lower cost [11–15] and the commercial availability. Our newly developed growth approaches based on spontaneous-relaxation-enhanced (SRE) Sn incorporation mechanism led to the demonstration of GeSn lasers with Sn content of 17.4% [16] and 22.3% [17, 18]. From device application perspective, the high-Sn-content alloys are desired since: i) for emitter, high Sn content increases the bandgap directness and thus enhances the light emission efficiency; ii) for detector, high Sn content extends the spectral response cutoff wavelength. In order to increase the Sn content, the lower temperature growth is required. However, low temperature growth results in Sn precipitation [19], breakdown of crystalline epitaxy due to the surface roughening, and lower growth rate, making the epitaxy process impractical. Therefore, increasing the Sn incorporation using GeH4 as the gas precursor would face a bottleneck with the conventional CVD techniques. In a recent publication it is shown that achieving high Sn content up to 22.3% requires multiple buffer layers of GeSn [17, 18]. This is the main motivation for seeking an alternative growth technique to facilitate the growth of higher Sn composition via CVD while the advantage of using the low cost and industry desired GeH4 precursor.
Plasma enhanced chemical vapor deposition (PE-CVD) was reported as a novel growth control mechanism for the growth of Ge on Si [20–22]. With the plasma enhancement, high-quality Ge could be grown at low growth temperature while maintaining a high growth rate. The success of plasma enhancement technique could potentially be applied to GeSn growth, which has not been studied yet. In this work we investigate the possibility to grow GeSn material system on Si substrate via PE-CVD growth technique using GeH4 and SnCl4 as precursors. The crystallographic properties of the deposited materials were studied by Raman spectroscopy, X-ray diffraction (XRD), and transmission electron microscopy (TEM). The Sn content was obtained by the data fitting of XRD reciprocal space mapping (RSM). The optical properties of GeSn on Si were studied by photoluminescence (PL) and Ellipsometry spectroscopies.
2. PE-CVD description and material growth
The configuration of plasma enhancement technique in cold wall ultra-high-vacuum CVD is schematically shown in Fig. 1(a). A Capacitively Coupled Plasma (CCP) was utilized to generate the plasma with 13.56 MHz radio frequency (RF) power supply and an L-shaped automatic impedance matching network. The 4-inch Si wafer holder was powered, as the anode of CCP, while conductor plate in parallel with the wafer holder was grounded, as cathode. The equivalent capacitance was formed between Si wafer and conductor plate. The GeH4 and SnCl4 precursors, and the Argon (Ar) carrier gas were delivered horizontally between the wafer and the plate.
Fig. 1 (a) A schematic diagram of PE-CVD system. (b) The generation of plasma between cathode plate and Si wafer in the CVD reactor. (c) The grown GeSn on Si wafer after PE-CVD process. The probing spots I, II, and III were marked on the wafer.
The plasma discharge between the wafer and the metallic plate is shown in Fig. 1(b) as the glowing region, which could be divided into two areas: i) the main area A between the substrate and cathode plate as a white glowing region and ii) the peripheral area B close to the gas inlet as a navy blue glowing region. To testify the uniformity of GeSn film across the wafer, Si substrates were intentionally not rotated during the growth. After the growth, the GeSn films were observed to be non-uniform with a cloudy region at the edge and a shiny area at the center of Si wafer, as shown in Fig. 1(c). To precisely investigate the GeSn film properties, three areas across the Si wafer, named as spots I (Center), II, and III (Edge), respectively, were investigated.
Prior to the material growth, the 4-inch (001) Si wafer was cleaned by the standard piranha solution followed by a hydrofluoric acid dipping process to remove the native oxide. For PE-CVD growth the plasma power supply was fixed at 50 W. In this study, three groups of Ge(Sn)/Si were grown:
- i) Samples A and B in group I are GeSn alloys that were grown with plasma enhancement at 350 and 400°C, respectively. To control the growth rate at low temperature, the Ar flow rate for sample A was 100 sccm in contrast with 200 sccm for sample B.
- ii) Samples C and D in group II are Ge that were grown with plasma enhancement aiming to be compared with samples A and B. During the Ge epitaxy, the wafers were rotating at 20 revolutions per minute (rpm).
- iii) Sample E in group III is GeSn that was grown via conventional CVD thermal method without plasma enhancement. During the GeSn epitaxy of sample E, the wafer was rotated at 20 rpm. The growth temperature and pressure of sample E were 350°C and 0.5 Torr, respectively. The flow fraction of GeH4 for sample E is 0.286, which is 5.7 times higher than 0.05 of sample A by PE-CVD. The high flow fraction of GeH4 ensures the GeSn deposition at growth temperature of 350°C. The flow rate of SnCl4 was 0.02 sccm in order to minimize the HCl etching effect that results from the byproduct reaction of gaseous during GeSn epitaxy [23]. The growth time was fixed at 120 minutes to obtain the comparable thickness with other samples.
The growth recipes of the five samples are summarized in Table 1.
Table 1. The summary of growth recipes of five Ge(Sn) samples by both PE-CVD and conventional CVD.
3. Results
3.1 Structural properties
The Raman spectra of samples A and B at spots I, II, and III were performed to investigate the crystallinity of GeSn, as shown in Fig. 2(a) and (b), respectively. The intensity of Raman spectra for sample A is low due to the high surface roughness and low material quality caused by low growth temperature (350°C). In order to obtain similar Raman intensity of sample A as sample B, the slit width of the light entrance of spectrometer for sample A was twice as big as sample B. The Ge-Ge longitudinal optical (LO) peaks of GeSn films were clearly observed, suggesting crystalline growth using PE-CVD technique was achieved. The Ge-Ge LO phonon of reference bulk Ge is located at 300 cm-1. Note that the full width at half maximum of Raman spectra for sample A is larger than that of sample B due to the wider slit width of the spectrometer entrance. The Ge-Ge LO peak of GeSn film shifts to the lower wavenumbers compared to bulk Ge due to the change of binding energy induced by Sn incorporation. From spot I to III of sample A (B), the Raman peak shifts from 297.8 (297.6) to 295.8 (296.7) cm−1, which is due to the supply change of Sn atoms across the wafer. At spots III of samples A and B, the minimum values of Raman peak shift corresponds to the maximum Sn incorporation.
Fig. 2 Raman spectroscopies of (a) sample A and (b) sample B at the spots of I, II, and III. A comparison of the Raman spectra for all samples is shown in (c).
A comparison of Raman shift for five samples (Spot III for samples A and B) is shown in Fig. 2(c). For samples C and D, the Ge-Ge LO peaks are at 302.9 and 301.5 cm−1, respectively. The longer wavenumber values of samples C and D, relative to bulk Ge reference, were attributed to the residual compressive strain at low growth temperature [24]. The Raman peak of sample E is located at 299.2 cm−1, corresponding to the Sn content of 0.75% [23]. Compared with bulk Ge reference, samples A and B show peak shifts of −4.5 and −3.3 cm−1, respectively. According to the relationship between Raman shift (∆ω) and Sn content ():for relaxed GeSn layer, Sn contents of samples A and B at spots III were estimated as 5.1 and 3.8%, respectively, where the coefficient of was adopted as −88 cm−1 for fully relaxed GeSn layer [25].
The XRD technique was performed to further investigate the Sn incorporation and compressive strain of GeSn. The spots III of sample A and B were chosen for the XRD measurements in order to evaluate the maximum Sn content. The 2Theta-Omega curves of samples A, B, and E along (004) direction are shown in Fig. 3(a). The peaks associated with Si substrate and GeSn can be clearly resolved. For sample E, the GeSn peak was located at the angle of 65.95°. The GeSn peaks of samples A and B shift to lower angles than that of sample E due to more Sn incorporation. Sample B exhibits a narrow GeSn peak at 65.64° while sample A presents a broadened GeSn peak between 65.33° and 65.57°. The broadening of XRD peak of sample A suggests the wide range of Sn incorporation. Both XRD peaks of sample A and B are wider than that of sample E, which is mainly due to the variance of Sn incorporation.
Fig. 3 (a) The 2Theta-Omega XRD curves of samples A, B, and E from (004) plane. The reciprocal space mapping of (b) sample A and (c) sample B from () plane.
To evaluate the compressive strain and Sn composition of GeSn layer, the XRD-RSMs of samples A and B were conducted from () plane, as shown in Fig. 3(b) and (c), respectively. The GeSn contour plots are close to the relaxation line indicating that GeSn layers are almost fully relaxed. The degree of relaxation of samples A and B were calculated as 92% and 97%, respectively. The Sn incorporation was obtained by the data fitting of rocking curves and RSMs. The bowing parameter of GeSn lattice constant was chosen as −0.066 Å. The Sn incorporation of sample A was calculated from 3.9% to 6% while Sn incorporation of sample B is 3%. By contrast, the Sn content of 0.75% for sample E was obtained [23], indicating that the plasma discharge significantly enhances Sn incorporation.
The typical bright field TEM images at spot III of sample B were shown in Fig. 4(a). The thickness of GeSn film was measured as 1027 nm. The threading dislocations were observed to propagate through the GeSn film in order to accommodate the large lattice mismatch between GeSn and Si. Based on TEM analysis results, the threading dislocation densities were estimated to be ~109 cm−2. The high-resolution TEM images at the surface of GeSn (marked as area A) and the interface between GeSn and Si (marked as area B) were shown in Fig. 4(b) and (c), respectively. The diamond cubic structure indicates the single crystal GeSn epitaxy on Si by PE-CVD with the ultra-high growth rate. At the interface between GeSn on Si (Fig. 4(c)), the stacking faults along (111) and () were observed. The inset of Fig. 4(c) shows the fast Fourier transform (FFT) pattern, which further confirms the single crystal formation while the streaks at <111> direction of FFT pattern is due to the introduction of stacking faults.
Fig. 4 (a) The bright field TEM images at the spot III of sample B. (b) The zoom-in TEM image at area A (Surface of GeSn). Inset: High resolution TEM image. (c) The zoom-in image at area B (Interface between GeSn and Si). Stacking faults (S.F.) were observed at the interface. Inset: Fast Fourier transform (FFT) pattern.
3.2 Optical properties
The room-temperature PL spectra at various spots of samples A and B were conducted, as shown in Fig. 5(a) and (b), respectively. A 1064 nm pulsed laser was used as pumping source with a 500 mW average pumping power, 45 kHz frequency and 6 ns pulse width. A liquid nitrogen-cooled InGaAs detector was utilized with the cut-off wavelength of 2300 nm. The different shifts of PL spectra at the probing spots I, II, and III is due to the non-uniformity of GeSn epitaxy under the current PE-CVD configuration. The PL emissions of both samples come from the direct bandgap transition of GeSn. The indirect bandgap transition was not observed, which is suppressed by the fast nonradiative recombination induced by defects. For sample A, the PL peak at spot III was clearly shown to be at ~2000 nm, while the PL spectra at the spots I and II were too weak to be identified. For sample B, the PL spectra from spot I to III were observed with different red shifts of PL spectra. Starting from the spot I, the peak is at 1486 nm, corresponding to the direct bandgap emission of GeSn with negligible Sn incorporation, i.e., Ge direct bandgap. At spot II, the main PL peak shifts to 1556 nm and peak intensity increases in comparison with spot I, suggesting that Sn atoms start to be efficiently incorporated into Ge. The spot III has the PL emission with the broadening peak wavelength between 1800 and 1950 nm. For both samples A and B, the maximum PL peak intensity was obtained at the spot III. Meanwhile, the longest wavelengths of PL peaks at spot III of both samples suggest the highest Sn incorporation. We do notice that at spot III the broaden PL peaks for both sample A and B were observed, which is mainly due to the spatial variance of Sn incorporation confirmed by XRD analysis. The defects might contribute to the broadening of PL spectra as well, which, however could not be distinguished in PL spectra.
Fig. 5 Room-temperature PLs of (a) sample A and (b) sample B at different probing spots on the Si wafer. (c) The comparison of PL spectra for five samples.
The PL spectra for all samples were normalized and plotted side-by-side in Fig. 5(c). The PL spectrum of bulk Ge is also included as the reference. Compared with direct (0.8 eV) and indirect (0.66 eV) bandgap energies of bulk Ge, the PL spectra at spots III of samples A and B have a clear red shift, suggesting that Sn atoms are efficiently incorporated into Ge lattice sites. Comparing with PL spectra of sample B, the PL peak of sample A exhibits the longer wavelength, indicating the higher Sn incorporation at lower temperature growth. No PL emission was observed for sample C. The non-radiative recombination introduced by defects dominates over both direct- and indirect-bandgap radiative recombination. Sample D shows the light emission of direct bandgap at the peak wavelength of 1500 nm, which has the blue shift compared with bulk Ge due to the residual compressive strain. The enhanced optical properties of sample D compared with sample C is due to the higher temperature growth. For sample E of GeSn by non-PE-CVD, the light emission from direct bandgap transition was observed at peak wavelength of 1650 nm, corresponding to the Sn content of 0.75% [23]. According to the PL spectra of sample A, B, and E, the plasma enhancement dramatically increases the Sn incorporation at low temperature, which is evidenced by the clear red shift of light emission.
The samples thicknesses were obtained by Ellipsometry measurements, as shown in Table 2. The growth time of GeSn samples (A, B) and Ge samples (C, D) by PE-CVD was 20 mins while it is 120 mins for GeSn sample (E) by non-PE-CVD. For GeSn on Si by PE-CVD, the growth rates at the spot III and I were measured as 57.7 and 35.2 nm/min for sample A, and 51.4 and 25.4 nm/min for sample B, respectively. For Ge on Si by PE-CVD, the growth rates are 21.1 and 21.9 nm/min for samples C and D, respectively. In comparison with Ge by PE-CVD, the growth rate of GeSn is significantly higher, suggesting that the introduction of Sn facilitates the film deposition on Si under the current PE-CVD configuration. For sample E using non-PE-CVD growth, the growth rate of GeSn is 5 nm/min, which is less than one tenth of the growth rate of spot III of sample A by PE-CVD, indicating that the growth rate of GeSn increases dramatically with plasma enhancement.
Table 2. The summary of film thicknesses and growth rates of five Ge(Sn) samples grown by PE-CVD and conventional CVD.
The absorption coefficients of the five samples were extracted by Ellipsometry, as shown in Fig. 6(a). For samples A and B, the measurements were performed at spots III which shows the maximum Sn content. For all the samples both direct and indirect bandgap absorptions were observed. When photon energy exceeds the direct bandgap energy , the direct bandgap transition is dominant over the indirect bandgap transition. The absorption coefficients of samples C and D were mostly overlapped with that of bulk Ge. Compared to the bulk Ge, both direct and indirect absorption edges of samples A, B, and E show pronounced red shift, suggesting the reduced bandgap energies due to the incorporation of Sn. The absorption edges of samples A and B shift towards a lower energy numbers than the absorption edges of sample E while sample A exhibits the lowest energy values, which is consistent with the PL results.
Fig. 6 (a) The absorption coefficients of five samples. (b) The data fitting of absorption edges for direct bandgap energy for samples A, B, and E.
Near the direct band edge, the interband absorption could be expressed as [26,27],
where α is the absorption coefficient of direct bandgap transition, is the photon energy, A is the constant and is the direct bandgap energy. By fitting the absorption coefficient of band edge, the direct bandgap energies of samples A, B, and E were extracted. Figure 6(b) shows the data fitting process of , where the direct bandgap energies of samples A, B, and E were extracted as 0.690, 0.735 and 0.772 eV, respectively.4. Discussion
The PE-CVD technique has been demonstrated to significantly enhance the Sn incorporation and the growth rate of GeSn epitaxy. The mechanism of Sn enhancement by plasma was first discussed in this section. The generation of energetic gas species with high reactivity in plasma discharge pushes the GeSn growth condition far beyond the normal equilibrium of conventional CVD. The surface reaction dynamics was strongly modified by plasma, which makes the growth rate less sensitive to the growth temperature compared to the traditional CVD [28]. With plasma assistance, the precursors of GeH4 and SnCl4 are expected to be easily decomposed at low growth temperatures, which therefore ensures the high growth rate, high Sn incorporation of GeSn epitaxy on Si.
A general view of the growth mechanism of PE-CVD is illustrated in Fig. 7. The gases are dissociated into free radicals and ions in plasma, and thus increase the growth rate at low temperature. The relative high growth rate could result in high Sn incorporation as more Sn atoms are buried inside the Ge lattice matrix before their segregation on the surface. Meanwhile, a high-voltage capacitive sheath is formed between the wafer and the bulk plasma. At the steady-state plasma discharge, the time-averaged electric potential in bulk plasma is presented as a positive value and the potential drops sharply across the sheath [28]. The ions of precursors in the bulk plasma could enter the sheath and be accelerated by the built-in electrical field in the sheath. As a result, the ions gain sufficient energy when arriving at the GeSn surface and the collision probability among reactive ions increases at the surface. Therefore, the surface mobility of adatoms enhances at low growth temperature and consequently increases the material crystallization.
Fig. 7 Left: The buildup of electric potential in both bulk plasma and sheath. Right: The ion transportation process from bulk plasma to Si wafer.
In order to explain the non-uniform plasma discharge across the wafer and its effects on the Sn incorporation, we divided plasma glowing region into area A and B. The area A (bright white glowing region) mainly corresponds to the discharge of Ar and GeH4 [29,30] while the area B (navy blue glowing region) corresponds to the discharge of GeH4 and SnCl4 [30,31]. The non-uniform plasma discharge could be explained as follows: The gaseous SnCl4 experiences a much higher decomposition rate than GeH4 [32,33]. As a result, the majority of SnCl4 is immediately dissociated and depleted at the edge of bulk plasma at area B. High Sn-content GeSn film was deposited at the edge of Si wafer. In contrast, since most of the SnCl4 was consumed at area B, the plasma discharge of SnCl4 at area A is negligible. Therefore, comparing with area B, much less Sn was incorporated into Ge at area A. The discrepancy of plasma discharge beneath the wafer results in the variance of spatial Sn incorporation across the Si wafer.
Although the PE-CVD technique has shown its promising prospects on the GeSn growth, the following issues must be addressed during the PE-CVD epitaxy: i) the plasma involvement in the CVD growth introduces the associated physical sputtering on the surface. The sputtering could physically etch the surface and make the GeSn growth rate difficult to be controlled. ii) The epitaxy uniformities regarding growth rate and Sn composition across the wafer are critical for the PE-CVD growth due to the complicated controllability of uniform plasma density, high reaction rates of ions, and high collision frequency between ions. The showerhead design beneath the wafer for gas feeding is one possible solution to ensure the uniform gas distribution and improve homogeneous GeSn deposition by PE-CVD. Other parameters such as plasma power, growth pressure, and gas flow ratio are critical as well to address the abovementioned issues. While more Sn could be incorporated into Ge lattice matrix at low growth temperature by PE-CVD technique, the GeSn crystal quality deteriorates as the growth temperature decreases. For the current growth condition there is a trade-off between Sn incorporation and crystal quality with respect to the growth temperature. In order to improve the GeSn material quality at low temperature, three solutions would be pursued: i) to grow relaxed Ge as the buffer; ii) to grow the Sn compositional gradient layer as the buffer; iii) to optimize the current growth condition of PE-CVD technique.
5. Conclusions
In summary, we present the significant enhancement of Sn incorporation of single crystalline GeSn on Si by using PE-CVD technique. A high growth rate of 51.4 nm/min was obtained by PE-CVD and the Sn composition was achieved up to 6%. Higher Sn content could be achieved if the growth condition of PE-CVD is optimized. The PE-CVD technique potentially offers an efficient solution to modify the surface reaction dynamics during GeSn epitaxy towards enhanced Sn-content GeSn film compared with the conventional CVD technique.
Funding
Air Force Office of Scientific Research (AFOSR) (FA9550-14-1-0205, FA9550-16-C-0016); National Science Foundation (NSF) (DMR-1149605); National Aeronautics and Space Administration Established Program to Stimulate Competitive Research (NASA EPSCoR) (NNX15AN18A).
Acknowledgments
We thank Dr. Mourad Benamara and Dr. Andrian Kuchuk for their assistance in TEM imaging and XRD measurements at Institute for Nanoscience & Engineering, University of Arkansas.
Disclosures
The authors declare that there are no conflicts of interest related to this article.
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