Modulation-doped quantum wells (QWs) of GaSb clad by AlAsSb were grown by molecular beam epitaxy on InP substrates. By virtue of quantum confinement and compressive strain of the GaSb, the heavy- and light-hole valence bands in the well are split and the hole mobility is thereby significantly enhanced. Room-temperature Hall mobilities as high as 1200–1500 cm2/V s were achieved for 5–10 nm QWs and biaxial strains of 1–3%. This contrasts with earlier work on GaSb/AlGaAsSb QWs on GaAs substrates in which the mobilities were found to fall off above 1% strain. Moreover, unlike in comparable InGaSb and InSb QWs, the high mobilities were maintained out to sheet densities of 3.5 × 1012/cm2. As a result, the sheet resistivities observed in the GaSb/AlAsSb wells reached record levels as low as 1500 Ω/□. Modeling indicates that this performance gain is due to the larger valence band offset of the GaSb QWs and the consequent reduction in scattering because of the better confinement and the lower doping levels needed for a given sheet charge.
► GaSb quantum wells were grown on InP substrates with AlAsSb buffer layers. ► Compressive strain enhances hole mobility in the GaSb. ► Hole mobilities of 1500 cm2/V s at 300 K were achieved. ► Sheet resistivities as low as 1500 Ω/□, world record for p-type III–V quantum well. ► Could lead to better performance in p-channel FETs and applications in III–V CMOS.
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A structure of gallium antimonide (GaSb) and gallium arsenide (GaAs) wafers is built to modulate light reflectivity at CO2 laser wavelength. A quantum well composed of GaSb/GaAs heterojunction with highly doped GaAs up to 3×1018 cm-3 is inserted inside a layer structure. A grating of periodic structure of GaAs and gold layer is added just below the substrate. Gsolver software is used to determine the reflectivity of incident light with the existence of free carriers. A voltage is applied to the doped layer to deplete the free electrons and the reflectivity is determined again. The significant difference in reflectivity between the two cases can be used to build a light reflectivity modulator device.
Figure 4 shows the weak but directly proportionate relationship between GaSb incorporation and Sb growth flux. It can be seen that the GaSb mole % is significantly reduced in the Ga-rich samples compared to the N-rich samples most likely due to the large quantities accumulated on the surface.
RBS data were measured from etched and pre-etched samples. For pre-etched material a value for the GaSb content in the bulk was established from plateau regions, which are clearly visible beyond the accumulated surface metal droplets, as shown in figure 5(a). Similar traces from etched samples gave a very close match for the GaSb composition in this region. The RBS data show a similar trend to the WDX data, but with larger GaSb content, shown in figure 5(b).
Figure 5. RBS measurements showing (a) Ga-rich sample, before and after removal of the metal drops and (b) GaSb mole percentage against Sb growth flux.
To compare the difference in measured composition between the WDX and RBS, more samples were analysed which had higher measured GaSb fractions, represented by the triangular data points shown in figure 6. The solid line is a guide to the eye of the relationship between the WDX and RBS measurements and shows the techniques agree well, diverging only in their estimates of the very dilute-Sb samples composition. The dashed line is a one to one correlation between RBS and WDX. The square and circular data points are the measured Ga-rich and N-rich results, respectively. The most dilute, Ga-rich, samples show a divergence where RBS predicts higher GaSb mole concentrations. The exact reason for this is unknown, however it is possible unknown factors are affecting the WDX ZAF iterative routines due to the large difference between the composition of the GaSb standard relative to these very dilute GaN1 − xSbx samples.
Figure 6. Comparison of the WDX and RBS technique's measurement of GaSb composition. Three sample series are shown including the Ga-rich (squares) and N-rich (circles) samples, as well as samples from Army research labs (ARL) which have a higher GaSb mole fraction (triangles). The solid line is a guide to the eye and the dashed line is the one to one correlation between RBS and WDX.
Absorption measurements from pre-etched Ga-rich samples were performed on samples with various GaSb mole percentages, determined by RBS. Figure 7 shows the absorption coefficient (α) as a function of the energy for a series of GaN1 − xSbx samples grown with 0 ≤ x ≤ 0.1%. The figure presents the clear observation of sub-gap absorption (<3.4 eV) for Ga-rich samples which had Sb present during growth. For the sample with no measured GaSb there is no observed sub-gap absorption. The sub-gap absorption can be seen at very low GaSb contents, which increases as the GaSb content increases.
Figure 7. Absorption coefficient (α) against energy for Ga-rich, GaN1 − xSbx samples grown with and without Sb.
Room temperature PL and CL spectra were measured for Ga-rich GaN1 − xSbx samples, as seen in figure 8. The CL samples were fully etched and the PL samples were unetched. For these Ga-rich samples strong luminescence was observed. Using a 5 kV, 20 nA, focused electron beam and 5 s acquisition time, point CL was performed at a number of points for each sample. Monte–Carlo simulations show the 5 kV electron beam deposits 90% of its energy within ≈100 nm of the surface for the compositions measured for Ga-rich samples.
Figure 8. (a) Typical room temperature CL and PL spectra for samples with Ga flux = 2.3 × 10−7 Torr, with and without Sb and (b) room temperature CL spectra for samples with various Ga flux and fixed Sb flux = 3 × 10−8 Torr.
The CL measurements in figures 8(a) and (b) show a strong GaN band-edge luminescence peak with 3.4 eV centre energy. Excitation studies, where the intensity of the electron beam excitation source was increased, show this peak to increase proportionately with beam intensity. The PL spectrum for the sample in figure 8(a) does not show a GaN band-edge peak but it should be noted that most other samples in this series did show a PL peak at 3.4 eV. A broad luminescence peak near 2.2 eV was also observed in Ga-rich samples (Ga flux > 2.3 × 10−7 Torr) where there was Sb present during growth. There was no 2.2 eV peak observed in samples grown under the same conditions, but with no Sb, however in this sample there is still a strong 3.4 eV peak. As discussed above the substitutional Sb is expected to introduce a localized energy level at ≈1.1 eV above the VBM, providing a possible explanation as to the origin of the broad 2.2 eV peak that could be attributed to the optical transitions from the CBM to the Sb level. It should be noted however that the observed peak energy coincides with the yellow luminescence peak very often observed in GaN. Figure 9 shows the plot of normalised 2.2 eV peak height versus Sb flux. The clearly observed increase of the peak intensity with Sb content support the notion that the localized Sb levels contribute to the emission in the 2.2 eV range.
Figure 9. Normalised CL (Pre and Post-etch) and PL (Pre-etch) 2.2 eV peak heights against Sb flux, with fixed Ga flux = 2.3 × 10−7 Torr.
There is a strong relationship between the Ga growth flux and the peak intensity, which increases linearly. Point CL observes a small variation (≈10 meV) of the centre energy of the 2.2 eV peak with position probed. There is no observed correlation between this and the growth conditions, possibly due to a small degree of lateral compositional inhomogeneity. Point CL also showed a large variation of the peak height with probing position, therefore CL maps were performed to see the extent of the luminescence inhomogeneity, seen in figure 10.
Figure 10. CL intensity map, showing the peak height in the range 2.0–2.4 eV, taken using an 8 kV, 10 nA, focussed beam for a sample containing Sb.
The CL map shown in figure 10 was performed with a 8 kV, 10 nA, focused electron beam, with 2 s acquisition time per pixel. The mapping area was 50 × 50 µm2. The map shows several bright features brighter than the mean value. Due to the very dilute nature of GaSb within these samples, the characteristic x-ray intensities are very low, which precludes a simultaneous map of Sb x-ray intensity
The compositional and optical characterisation of three series of dilute-Sb GaN1 − xSbx alloys grown with various Sb flux, under N and Ga-rich conditions, were presented. WDX and RBS measurements show that for the same growth conditions more GaSb is incorporated during the growth under the N-rich rather than the Ga-rich conditions. The optical properties of the Ga-rich samples were measured using room temperature CL, PL and absorption measurements, on etched and pre-etched samples. The strong Sb content dependent luminescence with a peak at 2.2 eV is attributed to the optical transition from the conduction band to the localized Sb levels. CL mapping revealed a large spatial variation of the peak intensity of this 2.2 eV peak. The strong luminescence from these samples continues to suggest that dilute-Sb GaN1 − xSbx alloys are an excellent material system for use in solar energy conversion devices.
This work was undertaken with support from the EPSRC UK, under grant numbers EP/I004203/1 and EP/I00467X/1. The MBE growth at Nottingham was also supported by the US Army Foreign Technology Assessment Support (FTAS) program (grant W911NF-12-2-0003). The characterization work performed at LBNL was supported by the Director, Office of Science, Office of Basic Energy Sciences, Materials Sciences and Engineering Division, of the US Department of Energy under Contract No. DE-AC02-05CH11231. Data associated with research published in this paper can be accessed by contacting the corresponding author.