Long-Wavelength Photoluminescence from Stacked Layers of High

Jan 17, 2013 - and Anthony Krier. †. †. Department of Physics, Lancaster University, Lancaster LA1 4YB, U.K.. ‡. Department of Physics, Warwick ...
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Long-Wavelength Photoluminescence from Stacked Layers of HighQuality Type-II GaSb/GaAs Quantum Rings Peter J. Carrington,*,† Robert J. Young,† Peter D. Hodgson,† Ana M. Sanchez,‡ Manus Hayne,† and Anthony Krier† †

Department of Physics, Lancaster University, Lancaster LA1 4YB, U.K. Department of Physics, Warwick University, Coventry CV4 7AL, U.K.

‡

ABSTRACT: We report the successful molecular-beam epitaxial growth of 10 stacked layers of GaSb/GaAs quantum rings using a new procedure. Exact control of the arsenic flux during capping helps to reduce the strong group-V As−Sb exchange reactions, enabling the rings to be capped at the same growth temperature (480 °C) without dissolution. X-ray diffraction and transmission electron microscopy indicate excellent structural quality and uniformity with no threading dislocations. This is due to the reduction in the average strain through the quantum-ring formation. The total ring density in the stacks is 1 × 1011 cm−2. An unusually long-wavelength quantum-ring photoluminescence peak of 1.3 μm is observed at low temperature, which is attributed to a reduction in the quantum-ring charging due to lower unintentional p-doping in the GaAs cap layer. The impact that this effect will have on future device designs in solar cells and lasers is also discussed.

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INTRODUCTION There is increasing interest in type-II GaSb/GaAs quantum dots (QDs) in which holes are strongly confined in the QDs while electrons are Coulomb bound in the GaAs matrix. The spatial separation of the electron−hole pairs increases the exciton lifetime, which is attractive for applications in solar cells (SCs),1 while the deep confinement potential for holes could provide room-temperature charge storage for memory devices.2 In an SC, the introduction of QDs into the active region extends the absorption edge into the infrared spectral range, which increases the short-circuit current, but it is necessary to stack multiple layers to increase the light absorption. However the large mismatch (7.8%) makes stacking GaSb QDs difficult, leading to the formation of relaxed QDs containing defects and to the generation of dislocations due to build up of internal strain. Previous work has focused on the interfacial misfit (IMF) array growth, which relieves the strain energy by the formation of misfit dislocations. Recently, we presented an alternative strategy for stacking multiple layers through the use of GaSb quantum rings (QRs), which reduces the net strain without generating lateral or threading dislocations.3 The resultant SCs displayed an extended infrared spectral response where the 1 sun short-circuit current was enhanced by nearly 6.0% compared to a GaAs control cell. GaSb nanostructures grown by molecular beam epitaxy (MBE) tend to preferentially form QRs rather than QDs, due to a combination of the large strain and strong Sb segregation.4 The latter is enabled by strong group-V As−Sb exchange reactions during the capping procedure, which can result in significant ring intermixing and dissolution;5 that is, the volume © 2013 American Chemical Society

of the rings is reduced substantially, resulting in a thin GaAsSb layer and weak photoluminescence (PL). The As−Sb exchange also creates a Sb floating layer at the growth surface, which can aggregate into further layers of rings as they are stacked. This creates an inhomogeneous ring distribution where the size and density of the rings increases up the stack, increasing the strain which generates threading dislocations.6 The use of a valved cracker cell for As (which supplies predominantly As2 instead of As4) will also increase the As−Sb exchange since As2 is much more aggressive than As4. The introduction of a “cold cap”7 procedure, in which the QRs were capped at a lower growth temperature (at 430 °C) to reduce As−Sb exchange, helped to preserve the rings and improved the QR PL intensity. However, this growth temperature is much lower than the optimum for high-quality GaAs (580 °C), which results in poor quality material, which increases the background carrier concentration.8 In an SC, this could increase the dark current and, in turn, reduce the open-circuit voltage. Here, we present an optimum growth technique that avoids the cold cap procedure to produce stacked layers of high-purity GaSb/GaAs QRs. The sample exhibits excellent crystalline quality and long-wavelength photoluminescence (PL) peaking around 1.31 μm at 1.4 K. Analysis of the PL spectra indicates that a GaAs cold cap increases the number of carbon acceptors, providing extra holes to the QRs, which then induce a blue shift of the emission energy due to capacitive charging. Received: November 14, 2012 Revised: December 22, 2012 Published: January 17, 2013 1226

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Table 1. Summary of Sample Growth Details and Propertiesa sample

growth procedure

R

Δθ (arcsec)

ε⊥

PL peak at 1.4 K (eV)

A B

hot cap (480 °C) cold cap (430 °C)

0.90 0.78

−205 −240

1.53 × 10−3 1.79 × 10−3

0.945 1.00

a Segregation coefficient, R; zeroth-order satellite peak, Δθ; average perpendicular strain, ε⊥; and the 1.4 K PL peak position at a laser power of 0.2 mW.

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EXPERIMENTAL SECTION

Two samples were grown on GaAs(100) substrates using a VG-V80 H MBE reactor. Details and characteristics of the samples are listed in Table 1. A thermal effusion K-cell was used to supply the Ga flux, and two (Veeco) valved cracker cells were used to provide Sb2 and As2. Each sample consisted of a 0.5 μm GaAs buffer layer that was grown at 580 °C, followed by 10 stacks of GaSb QRs and completed with a 100 nm thick GaAs cap layer. The growth process for the GaSb QRs is described as follows. First, the growth temperature is reduced under an As flux to 480 °C, and then the GaAs surface is soaked under an Sb flux (of 5 × 10−7 mbar) for 30 s, where the RHEED pattern changed from a 2 × 4 to a 1 × 3 reconstruction, indicating a GaSb-rich surface. Next, GaSb was directly deposited using a growth rate of 0.3 ML/s (corresponding to a Ga flux of 6 × 10−8 mbar), producing a total thickness of 2.1 ML. The RHEED pattern changed from streaky to spotty after the deposition of 1.3 ML of GaSb, indicating the formation of GaSb QRs. In sample A, the QRs were capped with a 10 nm GaAs layer at the same growth temperature of 480 °C (hot cap). In sample B, the QRs were cold capped, meaning that the growth temperature was reduced to 430 °C under an Sb flux for 4 min, and then capped with a 5 nm thin GaAs layer. Exact control of the As flux during the cap growth for sample A is crucial to preserving the rings; we found that an As/Ga flux ratio of 5 (corresponding to an As flux of 3 × 10−7 mbar) is optimal to prevent ring dissolution by As−Sb exchange while still maintaining high-quality GaAs growth. After the initial cap layer, the temperature was then increased to 580 °C under an As flux to grow the GaAs spacer layers (total thickness of 40 nm). The structural quality of the epilayers were characterized by crosssectional transmission electron microscopy (TEM) using a JEOL 2000FX operating at 200 kV. Images were taken under dark-field 002 imaging conditions. High-resolution X-ray diffraction (HRXRD) measurements were taken using a BEDE QC200 diffractometer. Low-temperature (4.2 K) photoluminescence (PL) measurements were performed using a frequency-doubled diode-pumped solid-state laser emitting at 532 nm and a 30 cm focal-length spectrometer combined with a Peltier-cooled InGaAs diode array. Optical fibers were used to transmit light to and from the sample with an illumination area of ∼2 mm2, giving laser power densities in the range from 1 × 10−5 to 100 W/cm2.

Figure 1. Cross-sectional TEM images from sample A (a) and sample B (b) that contain 10 sheets of GaSb QRs. The total ring density in the stack is ∼1 × 1011 cm−2. The total Sb content inside the white boxes is evaluated and is shown in Figure 2. (c) Sketch illustrating that the paired features represent a single QR cleaved through its center. (d) Magnified view of a single GaSb QR.

Table 2. Sizes and Densities of the GaSb QRs Determined from TEM Measurements sample

diameter (nm) outer/inner (±5 nm)

height (nm) (±0.5 nm)

density (cm−2) (per layer)

A B

23/9 24/11

1.7 1.6

∼1 × 1010 ∼1 × 1010

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RESULTS AND DISCUSSION Figure 1 shows cross-sectional TEM images from both samples where the formation of GaSb QRs can be clearly seen in each layer. The paired features represent a cleaved cross section through a QR, as illustrated in Figure 1c. The sizes of the QRs are nearly identical in each sample, and the density is approximately 1 × 1010 rings/cm2 per layer (Table 2). The average distance between the quantum rings is approximately 75 nm. The composition of the QRs are close to 100% GaSb with very pure GaAs centers (determined from previous STM measurements9). No threading dislocations were observed due to a reduction in the net strain through the QR formation. Since dark-field 002 images are primarily sensitive to composition in these materials and the contrast is related to the difference in mean group III and V atomic number, the Sb concentration can be estimated from the relative intensity of the different materials. Figure 2 shows the Sb content along the growth direction (z) for each sample, through the individual

Figure 2. Cross-sectional dark-field images showing the first four GaSb WL stacks from sample A (a) and sample B (b). The graphs plot the Sb content along the growth direction, z, which is evaluated using the integrated intensity from the TEM images. The peaks correspond to the individual WLs. The green curves indicate the segregation fits. 1227

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the satellite peaks is only ∼70 arcsec in both samples. The average perpendicular strain, ⟨ε⊥⟩, can be determined by11

wetting layers (WLs). In contrast to the QRs, the wetting layer (WL) is strongly intermixed, where the Sb content varies between 30 and 40%. This is presumably due to the different strain distributions in the QRs and WL.4 During capping, the strain will be extremely high at the QD peak so that Sb will be redistributed toward the edge, leading to the QR formation with very pure GaAs centers, whereas the WL is constantly diluted through As−Sb exchange. The Sb segregation profile can be evaluated by analyzing the Sb content above the wetting layer base, as shown in Figure 2 (the areas analyzed are highlighted by the white boxes in Figure 1). A clear exponential decay can be observed, which can be fitted using4 ⎛ −z ⎞ c(z) = A exp⎜ ⎟ + c0 ⎝ l0 ⎠

⟨ε⊥⟩ =

(3)

where θB is the Bragg angle of the GaAs substrate and Δθ is the zeroth-order satellite peak. The calculated values are listed in Table 1. These values are about an order of magnitude lower than those of conventionally grown strain uncompensated InAs/GaAs QD stacks (ε = 1.49 × 10−2)12 and approximately half of that for GaSb/GaAs IMF grown QDs (ε = 3.54 × 10−3).13 The slightly higher compressive strain in sample B may arise from the higher Sb content in the WL, which increases the strain and lattice mismatch. Figure 4a shows the PL spectra measured with an excitation laser power of 0.2 mW. In sample B, we observe a WL peak at

(1)

where l0 is the decay constant or segregation length, and c0 represents a weak Sb background content of ∼0.22%. Values of l0 equal to 2.7 ± 0.5 and 1.1 ± 0.2 nm were obtained for samples A and B, respectively, indicating that the Sb segregation length increases when the growth temperature increases, in agreement with previous work on InSb/InAs QD structures.10 This is due to stronger As−Sb exchange at higher growth temperatures. The capping layer thickness (see Table 1) could also influence l0, but we note that the Sb content in sample B falls to near the background count within the 5 nm cold cap. The segregation coefficient, R, can be calculated from eq 1 by ⎛ −a ⎞ R = exp⎜ ⎟ ⎝ 2l0 ⎠

sin(θB) −1 sin(θB + Δθ )

(2)

where a is the GaAs lattice constant, giving values of R equal to 0.90 ± 0.02 and 0.78 ± 0.03 for samples A and B, respectively. This means that the floating Sb layer is depleted more slowly in sample A (which increases l0). Figure 3 shows the HRXRD curves from the samples where several pronounced high-order satellite peaks as well as numerous Pendellosung fringes are observed, indicating high structural quality. The full width at half-maximum (fwhm) of

Figure 4. (a) Photoluminescence spectra measured at 1.4 K using an incident laser power of 0.2 mW where the transitions corresponding to the GaSb/GaAs quantum rings and wetting layer are identified. (b) QR PL peak energy as a function of laser power.

1.32 eV and a peak from the GaSb rings at 1.00 eV with a full width at half-maximum (fmhm) of 98 meV. This is consistent with previous results on single-layer GaSb QRs grown with a cold cap.14 In sample A, the QR peak is red shifted to 0.945 eV, with a narrower fwhm of 73 meV. This emission energy is lower than all previous reports on GaSb/GaAs QR/QD nanostructures. We also observe that the WL PL peak in sample A is at a higher emission energy (1.37 eV). Figure 4b plots the change in emission energy for the two samples over 6 orders of magnitude of laser power. In both cases, a strong blue shift of the QR emission peak is observed. Such blue shifts are

Figure 3. High-resolution X-ray diffraction spectra obtained from the GaSb/GaAs QR samples where several pronounced high-order satellite peaks are observed. The fwhm values of the satellite peaks are approximately 70 arcsec. 1228

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sample A. The transition and localization energies are E = 0.925 and 0.595 eV, respectively, in excellent agreement with those calculated in ref 16. Figure 5b plots the change in emission energy as a function of QR height. The height is the main geometric parameter that controls the emission energy, and the range is chosen to be roughly equal to the uncertainty on the TEM measurements. The emission energy varies by 32 meV, which is less than the difference observed between the PL peak energies (43 meV). The severity (defined by the ratio of the ring thickness to the overall size of the ring) of the QRs can also influence the PL energy;18 however, this can be discounted since, in these samples, the dimensions of the QRs are nearly identical. One possible explanation is that the higher Sb composition in the GaAs matrix would reduce the emission energy, but we note that the GaAs content inside and near the top of the QRs where the electrons reside is very pure in both samples. A likely alternative explanation for the large difference in the emission energy between the two samples is a reduction in the QR charging in sample A due to lower unintentional pdoping in the GaAs capping layer, which is justified below. The calculated values in Figure 5a,b correspond to the energy when no holes are present in the QRs. Reference 14 presented a single-layer GaSb QR system that exhibited discretely charged hole states, where an average charging energy of 21 meV was deduced to represent the energy required to overcome the repulsive Coulomb potential to add one heavy hole to the QR. This enables us to estimate the average number of additional holes per ring in each sample due to unintentional p-doping. We can define the minimum transition energy, (E0), as the energy between the QR ground state and the GaAs conduction band edge when the incident laser power is zero. The E0 values were found from Figure 4b by extrapolating a linear fit of the lowest few laser power data points to zero laser power. In sample A, as the laser power tends toward zero, the PL peaks near 0.944 eV (E0), which corresponds to an occupancy of one additional hole per ring, close to the calculated value of 0.946 eV for one hole (0.925 + 0.021 eV). In sample B, the PL peaks near 0.987 eV (E0), which corresponds to three additional holes per ring (0.988 eV). Calculations aside, the 43 meV difference between the QR PL peak energies of the two samples corresponds to two holes per ring. The additional holes in sample B are due to a higher level of unintentional incorporation of carbon acceptors in the GaAs cold cap, which supply additional holes to the QRs in the dark. A schematic illustrating this effect is shown in Figure 5c. It is wellknown that carbon is an electrically active impurity in GaAs and occupies substitutional sites of As, acting as a shallow acceptor. The presence of carbon in the samples was also confirmed by secondary ion mass spectroscopy. Using a ring density of 1 × 1010 cm−2 and assuming no carbon in the high-growthtemperature (∼ 580 °C) GaAs spacer layers, we can estimate the p-doping concentration in the capping layers (the hot and cold caps) to be approximately 1 × 1016 and 6 × 1016 cm−3 in samples A and B, respectively, which are of the same order of magnitude as those reported in ref 8.

characteristic features of type-II systems and are attributed to a combination of band bending and capacitive Coulomb charging, although the latter is more important in GaSb/ GaAs QDs and QRs.2,15 We note that the peak emission energies observed in Figure 4 predominately correspond to emission from the QR heavy-hole ground state. Recombination from excited heavy-hole states and light-hole states is likely to be present in the PL, but not resolved. The quantum-confined hole levels inside the GaSb QRs were calculated using a single-band Schrodinger solver using COMSOL software. The band alignments were estimated using the model-solid theory by Van De Walle and Martin with an unstrained valence band offset of 580 meV between GaAs and GaSb.16 Material parameters were taken from the review by Vurgaftman et al.17 Figure 5a shows the calculated band structure along with the first confined heavy-hole (hh) state for

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CONCLUSION In summary, we have presented a new growth procedure to produce stacked layers of high-quality GaSb/GaAs QRs, where the QRs were capped at the same growth temperature (480 °C), avoiding the use of a GaAs cold cap. This was accomplished through fine control of the As flux during capping, to reduce the strong group-V As−Sb exchange

Figure 5. (a) Calculated band structure of the GaSb/GaAs QR structure showing the first confined heavy-hole state. (b) QR peak energy as a function of the QR height. (c) Schematic illustrating the mechanism for charging the rings at zero laser power (in the dark) where holes are only supplied from carbon acceptors present in the GaAs cap. This induces a shift in the confined hole ground state. The additional holes in sample B are due to the higher level of carbon acceptors present in the cold cap. 1229

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(10) Semenov, A.; Lyublinskaya, O. G.; Solov’ev, V. A.; Meltser, B. Y.; Ivanov, S. V. J. Cryst. Growth 2007, 301−302, 58−61. (11) Nuntawong, N.; Birudavolu, S.; Hains, C. P.; Huang, S.; Xu, H.; Huffaker, D. L. Appl. Phys. Lett. 2004, 85, 3050. (12) Tatebayashi, J.; Nuntawong, N.; Wong, P. S.; Xin, Y. C.; Lester, L. F.; Huffaker, D. L. J. Phys. D: Appl. Phys. 2009, 42, 073002. (13) Tatebayashi, J.; Khoshakhlagh, A.; Huang, S. H.; Dawson, L. R.; Balakrishnan, G.; Huffaker, D. L. Appl. Phys. Lett. 2006, 89, 203116. (14) Young, R. J.; Smakman, E. P.; Sanchez, A. M.; Hodgson, P.; Koenraad, P. M.; Hayne, M. Appl. Phys. Lett. 2012, 100, 082104. (15) Hodgson, P. D.; Young, R. J.; Ahmad Kamarudin, M.; Carrington, P. J.; Krier, A.; Zhuang, Q. D.; Smakman, E. P.; Koenraad, P. M.; Hayne, M. Blueshifts of the emission energy in type-II quantum dot nanostructures. Manuscript in preparation. (16) Nowozin, T.; Marent, A.; Bonato, L.; Schliwa, A.; Bimberg, D.; Smakman, E. P.; Garleff, J. K.; Koenraad, P. M.; Young, R. J.; Hayne, M. Phys. Rev. B 2012, 86, 035305. (17) Vurgaftman, I.; Meyer, J. R.; Ram-Mohan, L. R. J. Appl. Phys. 2001, 89, 5815. (18) Ahmad Kamarudin, M.; Hayne, M.; Young, R. J.; Zhuang, Q. D.; Ben, T.; Molina, S. I. Phys. Rev. B 2011, 83, 115311.

reactions that cause ring dissolution, while maintaining highquality GaAs growth. The two samples investigated were grown under identical conditions, except for the capping procedure; sample A was capped with a 10 nm GaAs layer at the same growth temperature of the QRs (480 °C), whereas sample B was cold capped (430 °C) with a 5 nm GaAs layer. Analysis of the XRD spectra showed that the ring formation reduces the average perpendicular strain, enabling multilayers to be stacked without strain compensation. TEM studies revealed that the size and density, ∼1.0 × 1010 cm−2, of the QRs were nearly identical in each sample, but Sb segregation was higher in sample A. Despite the increased Sb segregation, the QR PL peak from sample A (0.945 eV) was much lower than that from sample B (1.0 eV) with a narrower fwhm (73 meV compared with 98 meV). This is attributed to a reduction in the QR charging in sample A as a result of lower unintentional pdoping in the GaAs cap layer. This means that the QRs in sample A could extend the absorption edge to longer wavelengths in an SC, while the lower unintentional doping could decrease the dark current and improve the open-circuit voltage. The lower number of holes could also reduce Auger recombination and hence the threshold current in GaSb/GaAs QD/QR lasers. In general, this work shows that exact control of the group-V fluxes and growth temperature is necessary to produce high-quality GaSb/GaAs QRs for future applications in SCs, memory devices, and lasers.

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AUTHOR INFORMATION

Corresponding Author

*Telephone: +44 01524 592780. E-mail: peterjcarrington@ gmail.com. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Financial support for this work was provided from EPSRC (grant EP/G070334/1) and in the framework of the QD2D project [grant EP/H006419], by the Royal Society - Brian Mercer Feasibility Award and QinetiQ (Agreement No.: 3000127730).

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REFERENCES

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dx.doi.org/10.1021/cg301674k | Cryst. Growth Des. 2013, 13, 1226−1230