The GaSb-incorporation in the samples grown under N-rich conditions was studied using RBS and WDX. Figure 2 shows the Sb profile measured by RBS from the sample with an Sb flux of 3.4 × 10−8 Torr and also the Monte–Carlo simulation of x-ray generation under 8 kV electron beam excitation.
Figure 2. RBS measured depth profile of the GaSb mole fraction for a typical N-rich GaN1 − xSbx sample (solid line) and Monte–Carlo Sb x-ray generation with depth (dashed line).
Due to the WDX surface sensitivity, Monte–Carlo simulations were used to estimate the x-ray generation rate with depth. A weighted average was then performed between the Monte–Carlo x-ray intensity with depth curve, and the measured RBS GaSb depth profile, shown in figure 2. This allowed a weighted average GaSb percentage to be determined for direct comparison of the RBS and WDX results. Figure 3 shows the WDX and RBS measurements of GaSb mole % incorporation with Sb growth flux for N-rich samples. WDX shows the lowest measured GaSb mole% to be (0.27 ± 0.01)% and the highest measurement to be (0.66 ± 0.02)%, assuming a systematic error of 1% of the measured value.
Figure 3. Plot of WDX and RBS GaSb mole percentage against Sb growth flux for N-rich GaN1 − xSbx layers.
An 8 keV, 40 nA electron beam was used to search for room temperature CL from these N-rich GaN1 − xSbx samples. No GaN1 − xSbx related luminescence peaks were observed in the range 330–850 nm.
The GaSb incorporation was found to be much lower in the Ga-rich GaN1 − xSbx samples. Due to the very small amounts of Sb extra care and analysis were required to quantify the GaSb content using WDX. To maximise the signal to noise a 7 kV electron beam, large counting times (240 s for the Sb L peak) and high currents (150 nA) were used. For each sample 10 random points were probed across the surface using a 10 µm defocused electron beam. In some cases the measured Sb x-ray counts were below the measured background for some of the data points and a negative value was then used in the calculation of the average GaSb atomic percentage. The resulting GaSb mole percentages are plotted against Sb flux in figure 4 which shows the lowest non-zero measurement to be (0.004 ± 0.002)% and the highest measured GaSb mole% to be (0.017 ± 0.001)%, assuming a systematic error of 5% of the measured value due to the large composition difference between the standards and the sample. At such low concentrations there may be additional uncertainties due to the correction procedures applied in the analysis of the WDX data.
In this work the compositional and optical characterization of three series of dilute-Sb GaN1 − xSbxalloys grown with various Sb flux, under N and Ga-rich conditions, are presented. Using wavelength dispersive x-ray microanalysis and Rutherford backscattering spectroscopy it is found that the N-rich samples (Ga flux < 2.3 × 10−7 Torr) incorporate a higher magnitude of GaSb than the Ga-rich samples (Ga flux > 2.3 × 10−7 Torr) under the same growth conditions. The optical properties of the Ga-rich samples are measured using room temperature cathodoluminescence (CL), photoluminescence (PL) and absorption measurements. A broad luminescence peak is observed around 2.2 eV. The nature and properties of this peak are considered, as is the suitability of these dilute-Sb alloys for use in solar energy conversion devices.
Highly Mismatched Alloys (HMAs) are semiconductor alloys where the substitutional atoms have very different atomic radii and/or electronegativity . Examples include GaNAs , GaNBi  and InNAs . Conventional semiconductor growth mechanisms have meant that such compounds could not be grown over a large range of compositions and were immiscible. Recently plasma assisted molecular beam epitaxy (PA-MBE) has been used to synthesise a number of HMAs over a large (or complete) range of compositions . For example, GaN1 − xAsx has been synthesised over the entire composition range and GaN1 − xBix which has been grown with a GaBi concentration up to ≈11% using very low growth temperatures [2, 3]. This allows their properties to be explored experimentally and compared with theoretical studies [4, 5].
HMAs display a large bowing of their bandgap with composition and their electronic structure is drastically different from their constituent binary materials. The Band Anti-Crossing (BAC) model has been successfully used to describe the electronic structure of the conduction and valence bands of HMA s in the dilute alloy limit [6, 7]. For the HMA GaN1 − xAsx the BAC model predicts a bandgap range of 3.4–0.7 eV with considerable bowing below the GaAs bandgap [2, 8]. Even stronger modifications of the band structure are expected for more extremely mismatched GaN-based alloys, such as GaN1 − xSbx and GaN1 − xBix. The large bandgap range and controllable position of the conduction and valence bands make these materials promising systems for use in solar energy conversion devices .
For example, theoretical calculations predict that the addition of As or Sb to GaN at concentrations below ≈10% can substantially lift the valence band edge and thus reduce the fundamental bandgap [8, 9]. The modification of the band structure enables the material to capture more photons from the solar spectrum while still maintaining the favourable alignment of the GaN band-edges with the redox potential of water for spontaneous hydrogen production by water splitting [4, 10]. Such materials can be used as the photoelectrode within a photoelectrochemical (PEC) cell .
The electronic band structure of HMAs is determined by the anticrossing interaction between the localized level of the mismatched anion and the extended states of the host matrix (i.e.: Sb in GaN1 − xSbx). The energy of the localized state can be deduced from the known location of the state in another III–V compound and the assumption that the energy of the state remains constant relative to the vacuum level. It has been found previously that the Sb level in GaAs is located at 1.0 eV below the valence band edge . This locates the Sb level at 1.0 eV above the valence band edge of GaN as the valence band offset between GaAs and GaN equals 2 eV [11, 12].
Transitions from the conduction band edge to this level would result in photons being emitted at an energy in the range 2.2–2.3 eV, namely 1.0 eV less than the GaN bandgap of 3.4 eV and including a Stokes' shift of about 0.1–0.2 eV. A similar situation has been observed in the GaN1 − xAsx HMA system, where the localized EAs level was determined by BAC fitting to be at 0.6 eV above the valence band edge . This was associated with a blue emission at about 2.6–2.7 eV [13, 14] and used to explain the absorption edge shift with increasing As content . Observation of the dilute Sb induced level would help to confirm the use of the BAC within the GaN1 − xSbx system and other similar alloys and would be useful in supporting the potential of these alloys for use in solar energy conversion devices.
GaN1 − xSbx is one important HMA and although there has been extensive study of the Sb-rich case  of GaSb alloyed with dilute amounts of N, there has been comparatively less reported on the dilute-Sb GaN1 − xSbx system. We have studied a wide range of growth temperatures for Sb doped GaN—from 10° C up to approximately 500° C . The incorporation of Sb increases with decreasing growth temperature. In this paper we concentrate on low GaSb concentrations as it is better to grow GaN layers doped with Sb at temperatures that are as high as possible in order to increase the quality of the layers. In this work the GaSb contents in several series of dilute-Sb GaN1 − xSbx layers are accurately quantified, mapped and correlated with the strength of luminescence peaks observed in cathodoluminescence (CL) and photoluminescence (PL) spectra.
2. Experimental details
The dilute-Sb GaN1 − xSbx epilayers were grown at ≈500° C using plasma assisted molecular beam epitaxy (PA-MBE) in a MOD-GENII system on two-inch diameter sapphire substrates. The active nitrogen for the growth of the group III-nitrides was provided by an HD25 RF activated plasma source. Standard Veeco effusion sources were used for Ga and Sb. In order to increase uniformity across the wafer, all films were grown with substrate rotation of ≈10 rpm. In MBE the substrate temperature is normally measured using an optical pyrometer, however, because uncoated transparent sapphire was used, the pyrometer measures the temperature of the substrate heater, not that of the substrate. Therefore in this study estimates of the growth temperature are based on thermocouple readings [4, 17].
Prior to growth the sapphire wafers were heated to ≈700° C and annealed for 20 min. After annealing, the substrate was cooled down to the growth temperature over a 20 min period under a reduced active nitrogen flux and growth was started by simultaneous opening of the Ga and N shutters. The Sb shutter was opened after a 1 min delay in order to avoid the deposition of any Sb on the sapphire surface before GaN growth. The growth time was kept at 2 h for all layers. The growth temperature was approximately 500° C resulting in low levels of GaSb incorporation. Studies of similar systems of GaN alloyed with group V anions , such as Bi, showed that in order to incorporate large amounts of the mismatched anions even lower growth temperatures were required, ranging from approximately 500° C to 80° C. The thicknesses of the GaN1 − xSbx layers are estimated to be approximately 500 nm.
A separate series of samples with higher GaSb contents were studied for comparison. These were grown at lower growth temperatures (275–375° C) in a second reactor. The material was grown using a Gen II Veeco solid-source MBE system equipped with Sb valved cracker sources, solid sources for Ga and a Uni-bulb plasma-source supplied the N. The samples were rotated at 5 rpm. The 2 inch sapphire substrates were outgassed at 800° C for 30 min. The growth time was 20 h.
There are three main growth regimes for PA-MBE of GaN ; N-rich growth (here the active nitrogen flux is larger than the Ga-flux), Ga-rich growth (here the active nitrogen flux is less than the Ga-flux) and strongly Ga-rich growth (here the active nitrogen flux is much less than the Ga-flux and Ga droplets are formed on the surface). In this study three dilute-Sb series were grown under the N-rich and Ga-rich regimes; with Sb fluxes varied up to 6.5 × 10−7 Torr. With the N supply held constant, the Sb and Ga growth fluxes for the samples in these three series are shown in figure 1.
Figure 1. Sb and Ga growth flux, representing three growth series: N-rich with various Sb flux; Ga-rich with various Sb flux; and fixed Sb flux with Ga flux which extends from N to Ga-rich. The dotted line indicates the region of the transition from N-rich to Ga-rich growth conditions.
Compositional studies were performed using Rutherford backscattering spectrometry (RBS); and by electron probe micro-analysis (EPMA) using a Cameca SX100 apparatus. A 3.04 MeV He2+ ion beam was used for RBS measurements to probe near the surface  and spectral fitting of the RBS data was performed using the SIMNRA  and SIMtarget  codes to obtain composition with depth information.
The EPMA has three wavelength dispersive x-ray (WDX) spectrometers, as well as the addition of an optical spectrograph and Silicon CCD array for Cathodoluminescence (CL) measurements [22, 23]. The samples are mounted on a precision stage which allows for micron-scale simultaneous mapping of WDX and CL signals . For quantitative WDX the peak-to-background signals were compared with GaN and GaSb standards to obtain the experimental k-ratios (sample intensity/standard intensity). The k-ratios were converted to atomic percentages using standard ZAF correction iterative procedures [22, 25]. For each WDX measurement 10 points were probed along a length of about 5 mm. For quantitative measurements electron beam energies of 7 or 8 kV were used, with currents between 100 nA and 150 nA. Larger currents and acquisition times were needed for the very dilute samples. Monte–Carlo simulations  show that 90% of the energy in a 7 kV beam is deposited within a depth of 165 nm for GaN1 − xSbx(x = 0.016%), corresponding to 90% of the Sb x-rays being generated within 100 nm of the surface. A higher beam energy of 8 kV was used for the N-rich samples. Monte–Carlo simulations show 90% of the 8 kV beam energy is deposited within a depth of 175 nm, which corresponds to an Sb x-ray generation depth of 125 nm.
For point CL a 5 kV, 20 nA, focused beam was used with a 5 s acquisition time to probe multiple points. For mapping the conditions were changed to a 8 kV, 10 nA, focused beam with a 2 s acquisition time to improve image resolution. Room temperature photoluminescence (PL) measurements were performed using a ≈ 5.6 eV CW laser. The absorption spectra were measured using a LAMBDA-950 UV/vis/NIR spectrophotometer over the range 190–3300 nm .
For samples grown under Ga-rich conditions there was a surface layer of metal droplets, composed of Ga accompanied in some cases by Sb. These were removed by etching for approximately 20 mins using concentrated Hydrochloric (HCl) acid in an ultrasonic bath. Confirmation of the removal of the metal droplets was performed using Secondary electron (SE) and back-scattered electron (BSE) imaging and by WDX mapping.
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We report on the modeling, growth, processing, characterization and integration in a gas detection setup of side wall corrugated distributed feed-back antimonide diode lasers emitting at 2.28 and 2.67 μm. The laser structures were grown by molecular beam epitaxy on GaSb substrate. Ridge lasers were fabricated from the grown wafers according to the following process: a second order Bragg grating was defined on the sides of the ridges by interferometric lithography, optical lithography and etched in a Cl-based inductively coupled plasma reactor. The devices exhibit a power reaching 40 mW, a side mode suppression ratio better than 28 dB and a tuning range of 3 nm at room temperature. One of these devices was successfully integrated in a tunable diode laser absorption spectroscopy setup, thus demonstrating that they are suitable for gas analysis.
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In this study, the feasibility of using wafer-bonding technology to fabricate a GaSb semiconductor on GaAs substrates for potentially creating a GaSb-on-insulator structure has been demonstrated. A GaSb wafer has been bonded on two types of GaAs substrates: (1) a regular single crystal semi-insulating GaAs substrate and (2) the GaAs wafers with pre-deposited low-temperature amorphous α-(Ga,As) layers. The microstructures and interface adhesion studies have been carried out on these wafer-bonded semiconductors. It has been found that the GaSb-on-α-(Ga,As) wafers have shown enhanced interface adhesion and lower temperature bonding capability.
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