1. Introduction

GaSb, an important member of the 6.1 Å family materials [1], has attracted wide attention from researchers due to its application in high electron mobility transistors [2,3], photovoltaic devices [4–6], and optoelectronic devices such as infrared lasers [7–10], infrared light-emitting diodes [8,11], and infrared photodetectors with short, medium and long wavelengths [6,12]. However, the research is limited by the high cost of native substrates and the lack of large size and semi-insulating GaSb substrates. In recent years, researchers have realized high-quality GaSb with a smooth surface on other conventional substrates (GaAs or Si) through heterogeneous epitaxial growth technology [13–16]. Among them, silicon substrate has several advantages including low cost, large size, high hardness, and high thermal conductivity compared with compound semiconductors [17]. Therefore, the realization of high-quality GaSb epitaxial layers on silicon substrates has attracted much attention since it allows the integration of optoelectronic and microelectronic devices based on 6.1 Å family materials with mature silicon technology [18,19]. Despite lots of research, it is very challenging to grow GaSb directly on silicon substrates especially complementary metal-oxide-semiconductor (CMOS)-compatible ones because of lattice mismatch, thermal expansion coefficient mismatch, and polarity difference between GaSb and Si, which lead to a large number of crystallographic defects include dislocations, stacking faults, cracks, twins, and antiphase domain boundaries (APBs). The defects drastically degrade the performance of the devices grown on silicon substrates [20–26]. These defects are related to the atomic arrangement at the heterointerface, which is extremely important for the quality of III-V epilayer [27]. The initial growth of III-V materials on Si always proceeded in a Volmer-Weber growth mode due to the interface energy of III-V/Si [28]. Therefore, significant differences in island size could be observed under different growth conditions and the early reports showed that the high defects density of III-V epilayer grown on Si substrates may be caused by the coalescence of these inhomogeneous large nucleation islands at the interface [29,30].

There is plenty of research showing that the initial process of the nucleation layer at the interface and its growth conditions are critical to the surface roughness and crystal quality of the GaSb epilayer when growing on a silicon substrate [14,31]. S.H. Huang et al. grew GaSb with an AlSb nucleation layer on the 5° miscut silicon substrate and found that 90° interfacial misfit arrays (IMF) formed at the interface between AlSb and Si, which led to a GaSb epilayer on Si with a low dislocation density [32]. Based on these results, they realized the electrically pumped GaSb quantum well edge-emitting lasers on the 5° miscut Si substrates [33]. Kyu Hyoek Yoen et al. obtained the III-Sb epilayer with surface RMS roughness of 1 nm on the 5° miscut silicon substrate in molecular beam epitaxy (MBE) system by utilizing thick AlSb buffer and AlSb/AlGaSb short-period superlattice as the dislocation-filter [34]. The results from E. Tournié et al. also confirmed the importance of the AlSb nucleation layer and dislocation arrays composed of Lomer dislocations were formed at the interface between the Si and the AlSb [35]. Besides, E. Tournié et al. also tried GaAs as a nucleation layer when growing GaSb on the 6° miscut silicon substrates and obtained a 500 nm GaSb epilayer with low defects density since the strain relaxed at the GaAs/Si and GaSb/GaAs interface through the formation of misfit dislocation grids [23]. Evangelia Delli et al. also obtained a smooth GaSb layer by inserting the AlSb nucleation layer on the 4° miscut silicon substrates. The RMS roughness and threading dislocation pit density of the 2 µm GaSb epilayer was 2 nm and 2 × 10⁸cm^-2, respectively. The values further decreased to 1.18 nm and 3 × 10⁷cm^-2 by implementing the stained AlSb/GaSb superlattices as the dislocation-filter [6]. Until now, various III-Sb optoelectronic devices have also been realized on the miscut silicon substrates such as lasers [36–38], photodetectors [6,18], and light-emitting diodes [11,39].

However, most research uses silicon substrates with a large miscut angle to avoid the generation of APBs, which are incompatible with standard CMOS fabrication technology where wafers are exact or on-axis (001) silicon substrates with a miscut angle less than 0.5° [40]. To obtain GaSb on on-axis Si, T. Baron et al. obtained a 250 nm APBs-free GaSb epilayer with outstanding optoelectronic properties after annealing the surface of on-axis silicon substrates under H₂ atmosphere in metal-organic chemical vapor deposition (MOCVD) systems to promote the formation of bi-atomic steps [41]. Kouichi Akahane et al. grew 2.5 µm GaSb epilayer on on-axis silicon substrate by using MBE technology, and the FWHM of the X-ray diffraction (XRD) rocking curve of GaSb epilayer was 215 arcsec [14]. Uğur Serincan et al. also grew 1 µm GaSb on on-axis Si substrate by applying post-annealing in MBE, and the FWHM of XRD rocking curve of GaSb epilayer was 260 arcsec [42]. Rio Calvo et al. experimentally demonstrated that the growth rate difference of the two polar domains leads to the APD burying process [43]. Recently, E. Tournié et al. have also realized mid-infrared semiconductor lasers based on on-axis silicon substrates in MBE [44,45], showing the feasibility of all-MBE technology for direct epitaxy of III-Sb semiconductors and devices on on-axis silicon substrate.

Although high-quality GaSb materials have been realized by MOCVD technology [41], it is very challenging to obtain Sb-based optoelectronic or microelectronic devices in MOCVD compared to MBE. This is because of the low volatility of Sb, the requirement of low growth temperature, and the strong affinity of AlSb-based compounds with O and C impurity [46,47]. Most III-Sb optoelectronic or microelectronic devices with high performance grown on native substrates were obtained by MBE technology which is less sensitive to thermodynamics [9,48–50]. There is a need to study all-MBE grown GaSb materials and devices on on-axis silicon substrate, which not only avoids quality degradation caused by potential contamination during the wafer transfer process but also simplifies the fabrication process and decreases the costs [51]. Moreover, there is little research on Sb-based materials and devices grown on on-axis silicon, especially systematic research on the defect annihilation mechanism.

In this letter, we extensively explore the nucleation layer's influence on defect annihilation, determine the growth conditions of the GaSb epilayer, and obtain a smooth and APB-free GaSb layer on on-axis Si (001) substrates by MBE eventually. The Al_xGa_1-xSb nucleation islands under different growth conditions were studied in detail. The results show that the size and density of the initial nucleation island were vital to the self-annihilation of defects. The defects were self-annihilated quickly in the initial stage following the coalescence of nucleation islands by using the optimized growth conditions of the AlSb nucleation layer and GaSb epilayer. The atomic force microscopy (AFM) image and (1 × 3) reflective high-energy electron diffraction (RHEED) patterns of GaSb epilayer grown on on-axis Si substrate show an APBs-free surface. In a word, we have obtained 2 µm APBs-free GaSb layer grown on on-axis (001) Si substrate in MBE with a surface roughness of 0.69 nm and an FWHM of 251 arcsec. Moreover, the optical performance of the GaSb epilayer and GaSb/AlGaSb quantum well grown on on-axis (001) Si substrate was also evaluated.

2. Experimental

2.1 Growth of GaSb on Si

The epitaxy growth of III-Sb semiconductor materials was carried out in the Riber Compact 21 T solid-source MBE system, which is equipped with antimony (Sb) valved cracker cell, aluminum (Al), and gallium (Ga) dual filaments effusion cells. Sb₂ was used to grow III-Sb materials by keeping the cracker temperature at 900 °C. The beam equivalent pressure (BEP) was used to evaluate the flux of III and V and calibrate the V/III ratio. The substrate temperature was measured using a C-type thermocouple, calibrated by monitoring the GaSb RHEED transition from (1 × 3) to (2 × 5) under the same Sb flux. And the different growth rates were calibrated through the RHEED oscillations of different layers on GaSb substrates. The 3-inch on-axis (001) silicon substrates were dipped in hydrofluoric acid (HF) (5%), rinsed with deionized water, and dried with high-purity nitrogen. Before transferring into the MBE growth chamber, the substrate was outgassed at 400 °C in the buffer chamber. The substrate was then heated to 1000 °C and held for 30 minutes in the growth chamber to remove residual oxide and cooled down to grow the nucleation layer. RHEED in the MBE growth chamber enables us to analyze and monitor the surface of the epilayer during the growth. The RHEED showed a clear (2 × 2) surface reconstruction, indicating that the oxide layer has been removed. Before growth, the substrate was soaked under the Sb₂ beam for 7.5 min to obtain an Sb-stable surface, which is considered to reduce the generation of APBs [15,52]. Following the soak, the III-Sb nucleation layer growth was initiated without changing the substrate temperature.

For the growth of III-Sb materials, a 5 nm Al_xGa_1-xSb nucleation layer was first grown on the deoxidized substrate, followed by a 45 nm GaSb capping layer grown at the same temperature to avoid destroying the nucleation sites. The substrate was then heated up to high temperatures for the growth of the subsequent GaSb layer. To study the influence of growth conditions on defect annihilation, nucleation layers with different Al composition (x = 0, 0.3, 0.5, 0.7, 1) was first grown to determine the optimal composition of the nucleation layer, different growth conditions such as growth temperature and V/III flux ratio at the nucleation stage and the GaSb layer growth stage were then carried out, respectively. Note that the growth rate of the nucleation layer in our experiment did not show the effect on the quality of the epilayer, different from other reports [53,54]. In our experiment, the growth rates of the nucleation layer and GaSb epilayer were 0.36 ML/s and 0.7 ML/s, respectively. The growth process was monitored in situ by RHEED.

2.2 Characterization of GaSb on Si

AFM was done using an MFP 3D scanning probe microscope to determine the surface roughness and morphology of the samples. The material quality of the epilayer was studied by high-resolution XRD using a Bruker D8 Discover Plus diffractometer. To evaluate the interface between GaSb/AlSb/Si, we prepared a cross-sectional transmission electron microscopy (TEM) sample using a Thermo Fisher scientific Helios5 CX-FIB/SEM. The structural and chemical integrities of the epilayer were investigated by high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) and high-resolution cross-sectional transmission electron microscopy (HRTEM) using an FEI Talos 200S microscope with super EDS operating at 200 kV. Photoluminescence (PL) spectroscopy was done at room temperature (RT) using a 785 nm laser for excitation to characterize the optical performance of the GaSb epilayer grown on on-axis (001) Si substrates.

3. Results and discussion

We have monitored the growth process GaSb on on-axis Si substrate using 5 nm AlSb as the nucleation layer. The typical substrate temperature change and shutter switch sequences are shown in Fig. 1(a). The on-axis Si substrate was first deoxidized at 1000 °C and then soaked under Sb overpressure. After the Sb soak, the AlSb nucleation layer, the GaSb capping layer and the GaSb epilayer were grown on the substrate sequentially. Finally, the sample was cooled to 300°C to check the (1 × 3) to (2 × 5) transition and then to room temperature after the growth of the entire GaSb epilayer. And the growth process was monitored in situ by RHEED. The inset in Fig. 1(a) shows a typical RHEED pattern of the 5 nm AlSb nucleation layer grown on the Si substrates. The sharp interconnected chevron spots indicate that the initial AlSb nucleation layer grows in the form of 3D islands [55]. The RHEED patterns of the epilayer surface at each growth stage are also shown in Fig. 1(b) - Fig. 1(g). A 45 nm GaSb capping layer was grown after the growth of the AlSb nucleation layer. The (3 × 3) RHEED pattern along with the [110] and [1–10] directions in Fig. 1(b) and Fig. 1(e), not only shows a smooth GaSb surface with a 2D planar growth mode but also shows that mixed reconstruction stripes of multidomain from numerous APBs. After 550 nm GaSb epilayer grown at high temperature, the RHEED pattern varies to (1 × 3) along with the [110] and [1–10] incident directions (Fig. 1(c) and Fig. 1(f)), which suggested that the formation of single domain and the coalescence and self-annihilation of APBs. Besides, as shown in Fig. 1(d) and Fig. 1(g), the reconstruction transformation of (1 × 3) to (2 × 5) was also observed during the cooling process after the growth of the entire GaSb epilayer on Si substrate, which is similar to the GaSb layer homoepitaxially grown on GaSb substrates [56].

 

Fig. 1. (a) Typical temperature and shutters sequence during the growth of AlSb nucleation layer and GaSb epilayer on silicon substrates. Inset: RHEED patterns after the growth of 5 nm AlSb nucleation layer. RHEED patterns along [110] direction after the growth of (b) GaSb capping layer, (c) 550 nm GaSb epilayer, (d) the entire GaSb epilayer during the cooling process, respectively. RHEED patterns along [1–10] direction after the growth of (e) GaSb capping layer, (f) 550 nm GaSb epilayer, (g) the entire GaSb epilayer during the cooling process, respectively.

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To explore the role of the nucleation layer on the surface roughness and crystalline quality of GaSb grown on on-axis (001) silicon substrates, we first studied the influence of the Al composition in the Al_xGa_1-xSb (x = 0, 0.3, 0.5, 0.7, 1), other growth conditions remained the same to ensure consistency of the experiment. AFM and XRD rocking curve (RC) obtained around GaSb (004) were used to evaluate these samples’ surface condition and crystal quality, respectively. Figure 2(a) shows the typical AFM image of the 600 nm GaSb epilayer grown on Si substrate with the AlSb nucleation layer, which has a surface with APBs since the layer is not thick enough for their annihilation. The variation of the surface roughness of the GaSb epilayer with different Al compositions is shown in Fig. 2(c). With the increase of Al composition in the nucleation layer, the surface roughness of the GaSb epilayer has a noticeable reduction and the surface morphology has a significant improvement. The GaSb epilayer has a surface topography with a large number of APBs and micro twins (MTs) propagating to the surface. The XRD rocking curves around the GaSb (004) of these samples are shown in Fig. 2(b), and the extracted FWHM of the GaSb (004) peak of the GaSb epilayer with different Al_xGa_1-xSb nucleation layer is shown in Fig. 2(d). Compared with the pure GaSb nucleation layer, the crystalline quality of the samples with the Al composition nucleation layer has significantly improved, which may be attributed to different interface energy and the kinetically related processes such as the different migration length of Al atom and Ga atom [28]. Based on these results, we choose AlSb as the nucleation layer for the following growth.

 

Fig. 2. (a) Typical AFM image (5 × 5 µm²) of the GaSb epilayer grown on on-axis Si substrate with the AlSb nucleation layer. (b) XRD rocking curves around the GaSb (004) peak of the samples with different Al_xGa_1-xSb nucleation layers. (c) Variation of the RMS roughness (10 × 10 µm²) of the GaSb epilayer grown on Si substrate with different Al_xGa_1-xSb nucleation layers. (d) The FWHM values extracted from the GaSb (004) peak of the samples with different Al_xGa_1-xSb nucleation layers.

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We further studied the effect of the nucleation layer growth conditions such as V/III flux ratio and growth temperature on the quality of the subsequent GaSb layer by growing a 600 nm GaSb epilayer containing a 5 nm AlSb nucleation layer, in which the GaSb epilayer was grown at 470 °C. In order to study the influence of different V/III flux ratios, the AlSb nucleation layer growth temperature remained at 430 °C. Figure 3(a) and 3(b) show the RMS roughness and XRD rocking curves around the GaSb (004) peak of the GaSb epilayer based on the AlSb nucleation layer with different V/III flux ratio. With the increase of V/III flux ratio, the surface RMS roughness and FWHM of GaSb epilayer decrease first and then increase. Compared to the other V/III ratio, there were relatively few APBs on the surface when the V/III ratio was 6, which caused a low RMS of 3.81 nm, and the GaSb (004) FWHM also has a low value of 560.7 arcsec, showing good crystalline quality. For the AlSb nucleation layer grown at different temperatures with a V/III ratio of 6, the RMS roughness and GaSb (004) FWHM of the 600 nm GaSb epilayer are shown in Fig. 3(c) and 3(d), respectively. The GaSb epilayer had a low RMS of 3.81 nm with the nucleation layer grown at 430 °C. The FWHM of the GaSb epilayer was improved from 604.8 arcsec at the 310 °C nucleation layer growth temperature to 514.8 arcsec at the 430 °C nucleation layer growth temperature and then deteriorated to 550.8 arcsec at the 470 °C nucleation layer growth temperature. The growth temperature and V/III flux ratio of the AlSb nucleation layer have the same influence on the surface RMS roughness and crystalline quality of the GaSb epilayer. At the low nucleation layer growth temperature of 310 °C and high nucleation layer growth temperatures of 470 °C, a large number of defects such as threading dislocations and APBs existed in the GaSb epilayer, causing poor surface and crystal quality.

 

Fig. 3. (a) Dependence of the RMS roughness (10 × 10 µm²) of GaSb epilayer grown on Si substrate on different V/III flux ratio of the AlSb nucleation layer. Inset: AFM images (5 × 5 µm²) of 600 nm GaSb epilayer on Si grown on the optimized V/III ratio of the AlSb nucleation layer (b) The FWHM values of the GaSb epilayer grown on Si substrate based on the AlSb nucleation layer with different V/III flux ratio. Inset: XRD rocking curves around the GaSb (004) peak of the samples based on different V/III flux ratio. (c) Dependence of the RMS roughness (10 × 10 µm²) of GaSb epilayer grown on Si substrate on different growth temperatures of the AlSb nucleation layer. (d) The FWHM values of the GaSb epilayer grown on Si substrate based on the AlSb nucleation layer with different nucleation layer growth temperatures. Inset: XRD rocking curves around the GaSb (004) peak of the samples based on different nucleation layer growth temperatures.

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To further study the growth mechanism, samples with a 5 nm AlSb nucleation layer only were grown on on-axis (001) silicon substrates at different growth temperatures. As shown in Fig. 4(a) - (d), the surface morphology and RHEED patterns of the AlSb nucleation layer have considerable differences at different growth temperatures. When the 5 nm AlSb was grown on Si substrate at the low temperature of 310 °C, the AFM image (shown in Fig. 4(a)) of the surface indicated that a quasi-continuous film with an RMS roughness of 0.5 nm was formed. The short elongated lines observed in RHEED patterns (shown in the inset of Fig. 4(a)) after the growth of 5 nm AlSb at low temperatures also proved the coalescence of the AlSb islands. As the increase of growth temperature, the 5 nm AlSb grown on Si substrate formed 3D islands because of the long atom migration length at high temperatures. Compared to samples grown at a temperature of 430 °C, the islands grown at the high temperature of 470 °C have a lower density and larger size, as shown in Fig. 4(b) – (d). The density of AlSb island was 1 × 10¹¹cm^-2, 8.5 × 10¹⁰cm^-2, 5 × 10¹⁰cm^-2, corresponding to the nucleation growth temperature of 390 °C, 430 °C, and 470 °C. The average lateral size and height of the AlSb islands were 14 nm and 3 nm at 390 °C, 18 nm and 5 nm at 430 °C, and 22 nm and 8 nm at 470 °C. The quasi-continuous AlSb film grown at low temperature and the low-density AlSb islands grown at high temperature provided fewer nucleation sites for subsequent GaSb growth, resulting in a poor GaSb epilayer with a large number of defects such as dislocations, APBs, etc. While for the growth of the AlSb nucleation layer at a suitable temperature, the dense 3D AlSb islands provide sufficient nucleation sites for the subsequent growth of GaSb, dislocations and APBs with opposite vectors will also be coalesced and annihilate during the epitaxy.

 

Fig. 4. AFM images (1 × 1 µm²) of 5 nm AlSb nucleation layer on Si grown at (a) 310 °C, (b)390 °C, (c)430 °C, (d)470 °C, respectively. Inset: RHEED patterns of 5 nm AlSb nucleation layer at the corresponding growth temperature. (e) Variation of the density, lateral size, and height of the AlSb nucleation island at different growth temperatures. The shadow shows the quasi-continuous film region.

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The epilayer growth condition was also vital to the quality of GaSb grown on Si substrate. The optimal growth temperature of the GaSb epilayer was investigated based on the optimized nucleation layer growth conditions. Our results show that the crystalline quality of the GaSb epilayer has been improved significantly with increasing growth temperature. However, the surface of the GaSb epilayer will deteriorate due to the desorption of Sb atoms at overgrowth temperature. The V/III flux ratio of the GaSb epilayer has the same influence on the surface RMS roughness and crystalline quality of the GaSb epilayer. Using the optimized growth conditions, we have grown GaSb epilayers with different thicknesses. Figure 5(a) shows the AFM image of the sample with 1.5 µm GaSb epilayer. Due to the self-annihilation of defects during the growth process, the 600 nm GaSb epilayer shows a single-domain morphology with an RMS roughness of 3.48 nm. And there are some defect pits with an average depth of 10 nm caused by dislocations and APBs, which originated from the lattice mismatch and the polarity differences during the heteroepitaxy. As for the samples with 1.5 µm and thicker, most defect pits have disappeared due to the coalescence and annihilation of APBs. The high-resolution XRD was used to estimate the crystalline quality of GaSb layers. Figure 5(d) shows the XRD rocking curves around the GaSb (004) of four samples. X-ray diffraction peaks broaden when the crystal lattice becomes imperfect. The most common sources of peak broadening are dislocations, stacking faults, twinning, micro stresses, grain boundaries, sub-boundaries, coherency strains, chemical heterogeneities, and crystallite smallness [57]. The FWHM of the GaSb (004) peak decreases from 591 arcsec to 251 arcsec with the increase of the GaSb epilayer thickness from 600 nm to 2 µm, which is caused by the coalescence and annihilation of threading dislocations and APBs. The roughness of the GaSb layer grown on on-axis (001) Si substrate was 0.4 nm for 900 nm GaSb by MOCVD [41] and 1 nm for 1 µm GaSb in MBE [42]. The FWHM of GaSb (004) grown on on-axis (001) Si was ∼350 arcsec for 1 µm GaSb in MOCVD growth [41] and 215 arcsec for 2.5 µm GaSb in MBE growth [14]. Therefore, we have obtained 2 µm high-quality GaSb layer grown on on-axis (001) Si substrate in MBE with a surface roughness of 0.69 nm (3 × 3 µm²) and an FWHM of 251 arcsec.

 

Fig. 5. (a) AFM image (3 × 3 µm²) of GaSb epilayer grown on Si substrate with a thickness of 1.5 µm, (b) XRD rocking curves around GaSb (004) of GaSb epilayers grown on on-axis Si substrates. (c) Dependence of the RMS roughness (10 × 10 µm²) of GaSb epilayer grown on Si substrate on the layer thickness. (d) The FWHM values of the GaSb epilayer grown on Si substrate with different layer thicknesses.

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The XRD reciprocal space mapping (RSM) around (004) and the (224) reflection of the GaSb layer grown on on-axis silicon substrates are shown in Fig. 6(a) and Fig. 6(b) to further examine the crystalline quality and residual strain inside the GaSb epilayers. Figure 6(a) shows the symmetric (004) RSM with the GaSb peak and the Si substrate peak aligned along the vertical line, indicating that the GaSb epilayer was not tilted compared to the Si substrate. The broadening along the qx direction was related to the average crystallite size that exhibited coherent diffraction [58]. The residual defects in the GaSb epilayer broke the coherence of the single crystal layer and reduced the lateral correlation length, causing the broadening in the qx direction. As shown in Fig. 6(b), the full-relaxation line in the asymmetric (224) reflection almost passed directly through the center of the patterns representing GaSb and Si, implying that the GaSb epilayer was nearly fully relaxed with respect to the Si substrate and no residual strain was present. The compact pattern of GaSb indicated a high crystalline quality of GaSb epilayer with fewer defects or mosaic grown on Si substrate [59].

 

Fig. 6. (a) RSM of the (004) reflection of GaSb and Si. (b) RSM of the (224) reflection of GaSb and Si.

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We further studied the formation, propagation, and extinction of defects by cross-sectional HRTEM. Figure 7(a) - (c) show the cross-sectional TEM image of the 1.5 µm GaSb epilayer with 5 nm AlSb nucleation layer grown on on-axis (001) Si substrate before optimization, in which the AlSb nucleation layer and GaSb epilayer were grown at 390 °C and 410 °C, respectively. As shown in Fig. 7(a), there were lots of defects such as APBs, threading dislocations and MTs [35] existing in the GaSb epilayer due to the difference between III-V and Si. Some of these defects were annihilated during the growth process, but many still propagated to 1 µm or even further, which were the origin of the poor GaSb surface topography and crystalline quality. The zoom-in cross-sectional TEM image of the GaSb epilayer is shown in Fig. 7(b), where the defects can be observed. The HRTEM image of the GaSb/AlSb/Si interface in Fig. 7(c) shows mismatch dislocations and MTs origin from the interface. Figure 7(d) - (f) show the cross-sectional TEM image of the optimized 1.5 µm GaSb epilayer with 5 nm AlSb nucleation layer grown on on-axis (001) Si substrate under the optimized growth conditions, in which the AlSb nucleation layer and the GaSb epilayer were grown at 430 °C and 490 °C. As shown in Fig. 7(d), few defects were observed in the GaSb epilayer, and most defects have annihilated in the initial GaSb layer. In the zoom-in TEM image (Fig. 7(e)), the area at the AlSb/GaSb/Si interface with a lateral size of 140 nm and a height of 70 nm show that the AlSb nucleation islands grown at optimized conditions had provided appropriate nucleation sites for the subsequent GaSb epilayer. Most defects also self-annihilated quickly with the coalescence of the initial islands. In the HRTEM image of Fig. 7(f), a sequence of periodic lattice distortion along the entire interface can be observed clearly, which originated from the strain fields due to bond bending around the misfit dislocation arrays [60]. The strain caused by lattice mismatch during the heteroepitaxy process was released by forming the interfacial misfit dislocation arrays in the AlSb/Si interface, so fewer dislocations can be seen in the GaSb epilayer. Compared to the samples before optimization, MTs also disappeared.

 

Fig. 7. (a) - (b) Cross-sectional TEM image and zoomed-in cross-sectional TEM image and (c) HRTEM image of the GaSb/AlSb/Si interface of 1.5 µm GaSb epilayer with 5 nm AlSb nucleation layer grown on on-axis (001) Si substrate before optimization, respectively. (d) - (e) Cross-sectional TEM image and zoomed-in cross-sectional TEM image and (f) HRTEM image of the GaSb/AlSb/Si interface of 1.5 µm GaSb epilayer with 5 nm AlSb nucleation layer grown on on-axis (001) Si substrate under the optimized growth conditions, respectively.

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Figure 8(a) shows the optimized samples’ cross-sectional high-angle annular dark field (HAADF) image. Most defects have been annihilated in the initial epilayer, indicating the high crystalline quality of the GaSb epilayer with fewer defects. Figures 8(b) – (e) show the EDX elemental mapping with Al, Si, Ga and Sb, respectively. The uniform distribution of these elements indicated no element intermixing between the Si substrate, the AlSb nucleation layer and the GaSb epilayer. The AlSb islands distribution with a height of about 6 nm can be identified on the superimposed elementary EDX chemical mapping of these elements, which is close to the AFM results of AlSb islands grown at 430 °C. It is worth mentioning that the morphology of the AlSb nucleation did not be significantly changed with the sequent growth of the GaSb capping layer. The AlSb islands with uniform size are tightly connected, and the subsequent growth of the GaSb capping layer promotes the rapid annihilation of defects. The results further confirmed the importance of the nucleation layer for the self-annihilation of defects.

 

Fig. 8. (a) Cross-sectional TEM high-angle annular dark field image (HAADF) of GaSb grown on on-axis Si substrate under the optimized growth conditions. (b) – (e) EDX elemental mapping with the Al, Si, Ga, and Sb map, respectively. (f) The superimposed chemical mapping of the elements Al, Si, Ga, and Sb.

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Finally, PL was used to characterize the optical performance of the GaSb epilayer grown on on-axis (001) Si substrate. Figure 9(a) shows the PL spectra at room temperature (RT) of the GaSb epilayer with different thicknesses grown on on-axis Si substrates under optimized growth conditions. As shown in Fig. 9(a), compared to the 600 nm samples, the PL intensity of 2 µm sample has significantly improved due to the annihilation of most defects. It is worth noting that no barrier layers were introduced into the measured GaSb layer to enhance carrier confinement. The corresponding FWHM of the 1.5 µm GaSb epilayer grown on Si was 113 nm, which is close to the GaSb substrate of 108 nm and shows a high-quality GaSb epilayer grown on Si substrates. The measured PL peak energy of the 2 µm GaSb was 0.711 eV, which is close to that of the GaSb substrate with 0.713 eV. In a word, we have obtained GaSb epilayers on on-axis Si substrates with high optical quality, comparable to that of GaSb substrate. We also grew GaSb/AlGaSb quantum well based on the above GaSb epilayers. The PL spectrum in Fig. 9(b) shows a peak at 1637 nm with an FWHM of 45.72 eV, which shows the possibility of applying the material system in the near-infrared optoelectronic devices.

 

Fig. 9. PL spectra of (a) the GaSb epilayer with different thicknesses grown on the silicon substrate; (b) GaSb/AlGaSb quantum well.

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4. Conclusions

In conclusion, using all-MBE technology, we have realized a high-quality GaSb epilayer with low defect density on on-axis silicon (001) substrate. We also explored the influence of nucleation layer growth conditions on defect self-annihilation of the GaSb epilayer and found that the composition, density, and size of the nucleation played an essential role in the surface morphology and crystalline of the GaSb epilayer grown on Si substrate. The optimal growth conditions of the GaSb epilayer were also determined. The surface roughness of the obtained 2 µm GaSb epilayer grown on Si substrate is 0.69 nm, and the XRD FWHM of the (004) GaSb peak is 251 arcsec. The result shows a good crystal quality of the GaSb layer with less mosaic. We further confirmed by HRTEM that GaSb epilayers grown on Si substrate have a good crystalline quality because most defects annihilated in the initial GaSb epilayer. The PL of the GaSb epilayer and GaSb/AlGaSb quantum well grown on Si substrate were also demonstrated. The success of all-MBE growth of Sb-based material system for infrared optoelectronic and microelectronic devices on COMS-compatible Si substrate paves the way towards their practical applications. Improved performance and new applications of such lasers, photodetectors, and non-volatile memories can be envisioned.