Growth and Characterization of GaAs 1-x-y Sb x N y /GaAs Heterostructures for. Multijunction Solar Cells Applications

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1 Growth and Characterization of GaAs 1-x-y Sb x N y /GaAs Heterostructures for Multijunction Solar Cell Applications Running title: Growth and Characterization of GaAsSbN/GaAs Heterostructures for Multijunction Solar Cells Applications Running Authors: Maros et al. Aymeric Maros a), Nikolai Faleev, Richard R. King and Christiana B. Honsberg School of Electrical Computer and Energy Engineering, Arizona State University, Tempe, AZ a) Electronic mail: amaros@asu.edu The GaAsSbN dilute-nitride alloy can be grown lattice-matched to GaAs with a bandgap of 1 ev, making it an ideal candidate for use in multijunction solar cells. In this work, using molecular beam epitaxy in conjunction with a radio-frequency nitrogen plasma source, we focus first on the growth optimization of the GaAsSb and GaAsN alloys in order to calibrate the Sb and N compositions independently of each other. After the optimum growth conditions to maintain 2D growth were identified, the growth of GaAsSbN films was demonstrated. Both a GaAsSb N /GaAs heterostructure (100 nm thick) and a GaAsSb N /GaAs quantum well (11 nm thick) were grown. XRD analysis reveals quite high crystal quality with a small lattice mismatch of %. SIMS profiling revealed that nitrogen was unintentionally incorporated in the GaAs buffer layer during the plasma ignition and stabilization. Nevertheless, a low temperature photoluminescence peak energy of 1.06 ev was measured for the GaAsSbN heterostructure sample while the quantum well emitted photoluminescence at 1.09 ev, which demonstrates promise for realizing 1-eV solar cells. 1

2 I. INTRODUCTION Conventional lattice-matched multijunction solar cells are composed of three subcells with Ge (0.67 ev) as the bottom cell, Ga(In)As (1.4 ev) as the middle cell, and GaInP ( ev) as the top cell. Efficiencies up to 41.6 % have been reported under 484 suns using this configuration 1. It has been shown that adding a fourth junction with a bandgap of 1 ev would allow better utilization of the solar spectrum and hence open the path to efficiencies greater than 50 % 2. The development of dilute nitride materials has attracted unique attention in that regard due to their ability to be grown lattice-matched to Ge and GaAs while independently tuning their bandgap can in the range 0.9 < E g < 1.3 ev by precisely controlling the alloy composition in the dilute nitrogen range (N < 5 %). For solar cells, the most commonly used material is GaInAsN. It has been used as a replacement material for the Ge bottom cell in conventional triple-junction solar cells and efficiencies as high as 44% have been reported under 947 suns 3. Despite the high efficiencies demonstrated using triple-junction solar cells, the use of dilute nitrides in tandem cells with 4 or more junctions has yet to be demonstrated. An interesting alternative to the most commonly used GaInAsN material for 1-eV solar cells is the GaAsSbN system which was first proposed by Ungaro et al. 4. As for GaInAs, adding a small amount of nitrogen to GaAsSb decreases both the bandgap and the lattice constant. However, due to a strong difference in band structure, and the fact that GaAsSb presents a stronger bandgap bowing than GaInAs, smaller bandgaps can potentially be achieved with GaAsSbN than with GaInAsN for the same nitrogen composition 5. This is a very attractive characteristic since the introduction of nitrogen is often associated with the creation of N-related defects that have shown to degrade the 2

3 minority-carrier lifetime in GaInAsN materials 6. Furthermore, by replacing In with Sb, nitrogen on a group-v site can bond only with Ga atoms, thus avoiding the formation of N-In defects and minimizing compositional modulation due to phase segregation 7. The introduction of Sb or N into GaAs has two very distinct signatures on the resulting band structure. On one hand, the addition of a few percent of Sb leads to a restructuring of the valence band. The heavy- and light-hole bands shift upward which induces a reduction of the fundamental bandgap 8. On the other hand, substituting group- V anions with a small amount of nitrogen leads to a splitting of the conduction band and a drastic reduction of the bandgap due to a strong interaction between the extended conduction band states and a narrow resonant band formed by the highly localized N states 9. The resulting band structures and bandgaps of GaAsSb and GaAsN are then calculated using a valence band-anticrossing (VBAC) model and a conduction bandanticrossing (CBAC) model respectively. In the case of the GaAs 1-x-y Sb x N y one needs to consider the contribution from both models and the band structure is then calculated using a double BAC model 10. Additionally, the lattice constant of the quaternary alloy can be calculated using Vegard s law by interpolating the lattice constant of the different binary compounds 11. By using this double anticrossing model in combination with Vegard s law, the GaAs 1-x-y Sb x N y alloy with a predicted bandgap close to 1 ev that is also lattice-matched to GaAs has Sb and N compositions of 6% and 2.3% respectively. In this paper we report preliminary results on the growth of the GaAsSbN alloy by solid-source molecular beam epitaxy (SSMBE) to approach the 1-eV bandgap while maintaining good material quality and lattice match to GaAs. We first focused our attention on optimization of the ternary GaAsSb and GaAsN alloys to ensure good 3

4 control of the Sb and N compositions independently, with the final objective of growing lattice-matched GaAs Sb 0.06 N on GaAs. II. EXPERIMENTAL DETAILS GaAsSb, GaAsN and GaAsSbN double heterostructures were grown on GaAs (001) substrates using a Veeco Gen III SSMBE system equipped with In, Ga and Al effusion cells and As, Sb and P crackers. The nitrogen was supplied by a Veeco Uni-bulb RF plasma source operated at 300 W. The nitrogen flow was controlled by a mass flow controller and the pressure in the chamber was monitored accordingly. The typical background pressure during plasma operation is Torr. The growth temperature was monitored using a pyrometer calibrated at the GaAs deoxidation temperature of 580 C while the Ga-induced oscillations were measured by reflection high energy electron diffraction (RHEED) to calibrate the growth rates. The V/III ratios were based on the beam equivalent pressure (BEP). Each sample was first subjected to a deoxidation step at elevated temperature for 15 min, after which a 300 nm thick GaAs buffer layer was grown at 600 C. For the GaAsSb structures, hereafter referred to as samples from set A, a substrate temperature of 500 C was chosen and 50 nm thick layers were grown at a growth rate of 0.7 µm/hr with a As/III ratio ~15. For the GaAsN and GaAsSbN samples, hereafter referred to as samples from sets B and C respectively, the plasma was ignited after the growth of the GaAs buffer layer while keeping the nitrogen shutter closed. In order to avoid any growth interruption and to provide enough time for the plasma to stabilize, a second GaAs buffer layer was grown with the plasma on and nitrogen shutter closed while the substrate temperature was reduced to the desired growth temperature, typically in the range C. This step generally took ~ 10 min. Once 4

5 the desired temperature was reached, the nitrogen shutter was opened and 100 nm thick GaAsN and GaAsSbN layers were grown. Additionally, a 11 nm thick GaAsSbN quantum well was also grown. The GaAsN layers were grown at growth rates of µm/hr with As/III ~10 while the GaAsSbN layers were grown at 1.0 µm/hr with As/III ratio ~10 and Sb/III ~0.16. All the samples were nominally undoped and capped with a 50 nm thick GaAs layer except for the quantum well sample which had a 100 nm thick cap layer to provide a slightly thicker well barrier. Table I summarizes the growth conditions used for all the samples investigated in this work. RHEED was monitored throughout the growth processes to control the growth surface morphology. High-resolution X-ray diffraction (HRXRD) was used to investigate the structural quality of the films and determine the alloy compositions while photoluminescence (PL) spectroscopy was used to assess their optical properties. The PL was excited by a 405 nm laser diode and detected using a liquid nitrogen cooled Ge detector. Furthermore, secondary ion mass spectrometry (SIMS) was used to obtain the composition of the quaternary alloys. TABLE 1. Summary of the growth conditions used for each sample presented in this work. T g corresponds to the growth temperature, R g to the growth rate and P F is the forward power of the nitrogen plasma. The presence of (-) indicates that there is no applicable data. Sample T g R g P As/III Sb/III F N % N % Sb % Sb % ( C) (µm/hr) (W) (XRD (SIMS) (XRD) (SIMS) A A A A

6 A A A A A A B B B B B B B B B C C III. RESULTS We first investigated the growth optimization of the GaAsSb and GaAsN alloys in order to obtain good control of the Sb and N compositions independently and to ensure that high crystal quality was obtained. Only after the optimum growth conditions were found for each ternary was the growth of lattice-matched GaAsSbN films considered. A. GaAsSb 1. Effect of Sb flux Five GaAsSb samples (A1 A5) were grown under the same growth conditions except for the Sb flux which was gradually increased in each sample. Figure 1a shows that there is an almost linear dependence between the Sb composition and the Sb beam equivalent pressure (BEP). The Sb compositions were evaluated by fitting the XRD ω-2θ scans presented in Fig. 1b. The Pendellösung fringes present in the XRD scans indicate that quite high crystal quality, specified by extended interference patterns and hence non- 6

7 deteriorated vertical coherence of epitaxial layers, was maintained over the whole range of Sb compositions grown ( % Sb). FIG. 1. (Color online) a) Dependence of the GaAsSb layer composition on the Sb beam equivalent pressure, b) XRD ω-2θ scans of the GaAsSb/GaAs heterostructures A1 to A5 with varying amount of Sb revealing high crystal quality for the range of compositions grown. The growth rate was kept constant at 0.7 µm/hr, the growth temperature at 500 C and the As/III at ~15 in all samples. 2. Effect of growth temperature Three additional GaAsSb structures (A6 A8) were grown at 475, 450 and 425 C to investigate the effect of the growth temperature on the Sb incorporation. All the other growth parameters, including the Sb flux ( Torr), were kept constant. Figure 2 shows that decreasing the growth temperature resulted in a noticeable increase in Sb incorporation. This phenomenon has been attributed to the differences in the atomic migration, sublimation energy and atomization energy of Sb and As 12. This suggests that in order to maintain the same Sb composition in structures grown at different temperatures, the Sb flux needs to be adjusted at each growth temperature. 7

8 FIG. 2. Effect of growth temperature on Sb incorporation. All the other growth parameters were kept constant (0.7 µm/hr growth rate, As/III ~15 and Sb/III ~0.23). Following these findings, two additional structures (A9 and A10) were grown at 460 and 420 C with different Sb fluxes (Sb/III ~ 0.32 and 0.28 respectively) while all the other growth parameters (As/III ~15 and growth rate of 0.7 µm/hr) were kept constant. The targeted Sb composition in both samples was 12 %. The results were then compared to the sample A5 (12.2 % Sb) grown at 500 C with Sb/III ~0.24. The XRD profiles presented in Fig. 3 revealed that similar Sb compositions (12.2, 12.8 and 12.6 %) were obtained in all three samples (A5, A9 and A10) by adjusting the Sb flux in each growth run. 8

9 FIG. 3. (Color online) XRD ω-2θ scans of GaAsSb/GaAs structures A5, A9 and A10 grown at different temperatures (500, 460 and 420 C). Similar Sb compositions were obtained in all three samples by adjusting the Sb flux in each growth run (Sb/III = 0.24, 0.32 and 0.28 respectively). The optical properties of the three GaAsSb/GaAs structures presented in Fig. 3 (A5, A9 and A10) were investigated using temperature-dependent PL. As shown in Fig. 4a the PL peak energy of the sample A5 grown at 500 C demonstrated a strong S-shape behavior at low temperatures. This type of PL behavior is characteristic of carrier localization effects induced by compositional fluctuations 13. At low temperatures, carriers become trapped in Sb-related localized states. As the temperature increases, carriers have enough thermal energy to escape these localized states and the PL peak energy follows the expected decrease with temperature. 9

10 FIG. 4. Temperature-dependent PL response measured on a) sample A1, b) sample A9 and c) sample A10. The S-shape behavior observed at low temperature is thought to be the result of carrier localization effects induced by compositional fluctuations. These localization effects are almost completely suppressed by reducing compositional fluctuation at the growth temperature 420 C. Lowering the growth temperature to 460 C resulted in a smaller S-shape behavior indicating a reduction in compositional fluctuations, as shown in Fig. 4b. This S-shape behavior disappeared when the temperature was further lowered down to 420 C, as shown in Fig. 4c, which indicates that growth at this lower temperature suppressed the compositional fluctuations and hence minimized the carrier localization. Similar findings were reported in mixed As-Sb alloys (namely InAlAsSb grown on InP) further indicating that lower growth temperatures seem to be favorable to diminish alloy fluctuations

11 B. GaAsN 1. 2D to 3D growth transition One of the main challenges when growing dilute nitride materials is to maintain smooth two-dimensional (2D) growth. In the presence of nitrogen, the GaAs surface is often subject to nitridation, i.e., the nitrogen accumulates at the surface and tends to form clusters 15. This generally translates to a spotty RHEED pattern which indicates threedimensional (3D) growth and an atomically rough surface. For high crystal quality 2D growth is critical. Our investigation of the effect of various growth parameters on the surface reconstruction of GaAsN structures revealed that the main factor controlling the transition from 2D to 3D growth is the background pressure in the chamber, which is a direct correlation of the nitrogen flow. In order to strike the plasma, the pressure in the growth chamber needs to be taken up to Torr while the forward power must be W. At this pressure the plasma ignites but is very dim. By reducing the background pressure to Torr the plasma becomes very bright. However when growing GaAsN films at this pressure, the RHEED pattern switches from a streaky (2 4) configuration to a spotty (1 1) pattern as soon as the nitrogen shutter is opened, as shown in Fig. 5a, indicating a fast transition to the undesired 3D growth mode. By gradually reducing the nitrogen flow rate until the background pressure hits Torr, the RHEED pattern remains streaky throughout the entire growth indicating that 2D growth is maintained (Fig. 5b). 11

12 a) b) FIG. 5. (Color online) RHEED pattern of two samples grown using different background pressure in the growth chamber a) Torr and b) Torr. A spotty pattern is indicative of 3D growth while a streaky pattern indicates smooth 2D growth. 2. Effect of nitrogen flux, plasma power and growth temperature After ensuring that 2D growth mode was maintained throughout the growth, the effect of the growth temperature, growth rate and plasma forward power on the nitrogen incorporation was investigated. The nitrogen flow was kept constant throughout the various growths with a corresponding background pressure of Torr. As shown in Fig. 6a, similar nitrogen compositions were obtained in samples grown at temperatures between 420 and 480 C indicating that N incorporation is temperature independent in this range. These results are consistent with other published results 16,17. Figure 6b reveals that increasing the group-iii growth rate results in lower N incorporation, which also agrees with the literature 18. On the other hand increasing the forward power seems to increase the N incorporation, most especially at lower growth rates. For higher growth rates varying the forward power does not seem to affect the N incorporation. These results are summarized in Table 1. The N composition was determined from XRD using Vegard s law. SIMS was also performed on three GaAsN structures (B4, B5 and B6) in order to compare the 12

13 results with XRD. It can be seen from Table 1 that the N compositions obtained from XRD differ somewhat from that obtained by SIMS, within the range expected from different analytical techniques. Several uncertainties arise from either technique which can result in small discrepancies. The Sb and N compositions obtained from SIMS were calibrated using an InGaAsSbN standard sample and therefore, the results depend on the accuracy of the compositions originally determined in the reference material. Furthermore, XRD represents a measure of the lattice parameter which is directly related to the composition. A slight deviation from Vegard s law as used in this work will also induce error bars in the measurements. Consequently additional care must be taken when comparing results from different analytical techniques. FIG. 6. (Color online) Nitrogen composition dependence on a) the growth temperature and b) the growth rate and forward power of the plasma. Results are also summarized in Table 1. C. GaAsSbN 13

14 Two GaAsSbN/GaAs structures were grown under the same growth conditions as the GaAsN samples but with different thicknesses, C1 (100 nm heterostructure) and C2 (11 nm quantum well). As shown in Table 1, other than the introduction of a Torr Sb flux, the growth conditions were identical to that of sample B5 (100 nm GaAsN heterostructure) which served a baseline to analyze the effect of Sb on the N incorporation. Figure 7a shows the XRD double-crystal -2 scans and the calculated profiles of the two samples C1 and C2. The GaAsSbN layers appear to be approximately 1.2 thicker than expected, and to be under slight compressive strain with a corresponding lattice mismatch of approximately % at room temperature. SIMS was performed on these structures in order to determine the Sb and N compositions. The SIMS profile of the thicker GaAsSbN sample C1 is presented in Fig. 7b showing compositions of 7.6 and 1.8 % for the Sb and N respectively and a thickness of 120 nm. Although grown under the exact same growth conditions, the SIMS profile of the quantum well sample C2 (not shown) revealed slightly lower Sb and N incorporation (7.3 and 1.5 %). This appears to be an artifact due to slight roughening as a function of depth occurring during the measurement which caused some amount of profile broadening in the quantum well sample, i.e., the Sb and N compositions are underestimated in the quantum well region. In that regard and considering that both samples were grown one after the other under the same growth conditions, it is very likely that the composition in both samples are closer to each other than what has been measured by SIMS. Although based on a small sample size, the N composition determined from SIMS in sample B5 14

15 (GaAsN ) and samples C1 (GaAsSb N ) and C2 (GaAsSb N ) suggest that the N composition does not vary with the introduction of Sb within experimental error. The SIMS profile in Fig. 7b also reveals that even though the nitrogen shutter was closed during the growth of the GaAs buffer layer, up to 0.1% N was unintentionally incorporated into the GaAs buffer. Furthermore, we also observed that the Sb and N profiles do not end at the same depth although both shutters were closed at the same time. This has been observed by other groups as well 19,20 and has been attributed to an accumulation of Sb atoms on the growth front which continues to incorporate in the layer even after the Sb shutter is closed until all the atoms have been desorbed. FIG. 7. (Color online) a) XRD -2 scans of the 100 nm thick GaAsSb N heterostructure (HS) and 11 nm thick GaAsSb N quantum well (QW). The dashed lines represent the calculated profiles. b) SIMS profile of the HS sample C1. The PL was measured at 10 K after rapid thermal annealing (RTA) both structures (C1 and C2) at 700 C for 5 min under N 2 flow. Post-growth annealing has been shown to reduce the density of N-related defects and as a result improve the optical properties of 15

16 dilute nitride materials 21. It can be seen from Fig. 8 that whereas only one emission line was observed at 1.06 ev for the thicker structure C1, the quantum well structure C2 demonstrated two PL peaks at 1.09 and 1.44 ev. It is believed that the main peak at 1.09 ev corresponds to the GaAsSbN layer while the second peak at 1.44 ev corresponds to the GaAs buffer layer and its top N-contaminated part. Due to the thicker Sb-N layer in structure C1 and the high energy excitation of the laser, most of the incoming power is absorbed within the GaAsSbN layer of the film and as a result the higher energy peak corresponding to the GaAs:N buffer layer is not observed. These results are encouraging considering that the targeted room temperature bandgap is 1 ev however, further tuning of the growth parameters will be necessary to reach the desired compositions for 1-eV lattice-matched GaAsSb 0.06 N Furthermore, N contamination of the GaAs buffer layer is undesirable and will need to be properly addressed in order to improve the material quality. FIG. 8. (Color online) Low-temperature PL of sample C1 (100 nm thick GaAsSb N ) and C2 (11 nm thick GaAsSb N quantum well) after RTA at 16

17 700C for 5 min in an N 2 atmosphere. Both sample structures are shown in the inset where the grey areas represent the GaAsSbN layer. IV. DISCUSSION AND CONCLUSIONS The incorporation behavior of group-v elements in GaAsSbN is in general still not fully understood. Several groups reported that Sb favors a higher N incorporation 7,22,23 whereas Ma et al. 12 found that increasing the Sb flux did not change the N composition in the temperature range 420 to 490 C. The SIMS analysis performed in this work tend to agree with the later; the N incorporation does not seem to vary with the presence of Sb, within experimental error. However, it should be noted these observations are based on a very limited sample size and additional investigations will need to be performed to fully understand the incorporation process of the various elements in our materials. SIMS analysis also revealed that unintentional nitrogen of up to 0.1% is incorporated in the GaAs buffer layer even when the nitrogen shutter is closed. This is undesirable since N incorporation might result in the introduction of point defects into the starting GaAs buffer which would in turn degrade the crystal quality, and hence, the device performance. Freundlich et al. have reported similar contamination of their GaAs material and have suggested using a gate valve that allows the plasma to operate continuously during the growth without interfering with the nitrogen-free layers 24. Wistey et al. proposed using an As capping layer in order to protect the wafer surface during plasma ignition and stabilization 25 and to additionally use deflection plates in order to reduce the defect density by driving nitrogen ions out of the plasma 26. These different 17

18 options represent opportunities to avoid degradation of the material quality and/or unintentional nitrogen doping in future work. In conclusion, we reported on the growth demonstration of GaAs 1-x-y Sb x N y /GaAs heterostructures for 1-eV solar cell applications. The Sb and N compositions were calibrated independently of each other by optimizing the growth of GaAsSb and GaAsN ternary alloys. Carrier localization effects resulting from compositional fluctuations were observed in GaAsSb structures grown at temperatures above 420 C. The background pressure in the MBE chamber, which is directly related to the nitrogen gas flow, was found to have a major effect on the transition from 2D to 3D growth of GaAsN. During growth, chamber background pressures below T resulted in smooth 2D growth whereas higher pressure typically resulted in 3D growth. The 100 nm thick GaAsSb N structure demonstrated a PL peak energy of 1.06 ev at 10 K while the 11 nm GaAsSb N quantum well sample demonstrated a PL peak at 1.09 ev. In both samples, XRD revealed a small lattice mismatch of % while SIMS analysis showed that up to 0.1% N incorporated into the GaAs buffer layer during the nitrogen plasma ignition and stabilization step. Possible solutions to avoid this undesirable N contamination and further analysis of the interaction between Sb and N in GaAsSbN are yet to be investigated. ACKNOWLEDGMENTS This work was primarily supported by the Engineering Research Center (ERC) Program of the National Science Foundation (NSF) and the Office of Energy Efficiency and Renewable Energy of the Department of Energy (DOE) under NSF Cooperative 18

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