Letter pubs.acs.org/NanoLett
Selective-Area Epitaxy of Pure Wurtzite InP Nanowires: High Quantum Efficiency and Room-Temperature Lasing Qian Gao,*,† Dhruv Saxena,† Fan Wang,† Lan Fu,† Sudha Mokkapati,† Yanan Guo,† Li Li,‡ Jennifer Wong-Leung,†,§ Philippe Caroff,† Hark Hoe Tan,† and Chennupati Jagadish† †
Department of Electronic Materials Engineering, Research School of Physics and Engineering, ‡Australian National Fabrication Facility, Research School of Physics and Engineering, and §Centre for Advanced Microscopy, The Australian National University, Canberra, ACT 0200, Australia S Supporting Information *
ABSTRACT: We report the growth of stacking-fault-free and taper-free wurtzite InP nanowires with diameters ranging from 80 to 600 nm using selective-area metal−organic vapor-phase epitaxy and experimentally determine a quantum efficiency of ∼50%, which is on par with InP epilayers. We also demonstrate room-temperature, photonic mode lasing from these nanowires. Their excellent structural and optical quality opens up new possibilities for both fundamental quantum optics and optoelectronic devices. KEYWORDS: III−V semiconductors, nanowires, wurtzite, quantum efficiency, nanowire laser, selective-area metal−organic vapor-phase epitaxy
N
leads to the formation of defects, including twinning and polytypism,14 which limits applications requiring larger diameters, such as photonic NW lasers.15 Moreover, Au atoms can be directly incorporated into the NWs grown by Au-seeded VLS growth,16 which may be detrimental to their optical performance.17,18 Chemically removing gold after growth and overgrowing the NWs to form a larger diameter adds avoidable complexity and could compromise their optical quality.19 In contrast, selective-area epitaxy (SAE) growth of NWs is intrinsically free of metal particles.20 In addition, it has many other advantages over self-seeded and metal-seeded VLS growth, such as the ability to eliminate tapering altogether, reduce impurity concentrations due to higher growth temperature, and produce NWs with large diameters. Selective-area MOVPE (SA-MOVPE) has been successfully used to grow InP NWs with good crystal quality.21 The high quality of such SAEgrown NWs has been confirmed in devices such as highefficiency solar cells.21,22 Nonetheless, the QE of the SAEgrown III−V NWs, which is the critical performance indicator when comparing with their planar counterparts, has not been reported yet. In this Letter, we demonstrate the successful growth of stacking-fault-free wurtzite (WZ) InP NWs with wide range of diameters, varying from 80 to 600 nm, using SA-MOVPE. We quantify the QE of the WZ NWs to be equivalent to the best quality 2D layers. Finally, as a result of the excellent structural
umerous optoelectronic devices use direct bandgap III−V semiconductors. In terms of reduced structural defects, low impurity levels, and low surface recombination velocities (SRVs), excellent material quality is necessary to obtain high efficiency devices. Down-scaling the footprint of devices to the nanoscale together with their integration onto silicon has significant technological outcomes.1 III−V semiconductor nanowires (NWs) grown using the bottom-up approach are promising for the realization of integrated nanoscale devices.2−4 However, material quality and control issues for NWs still remain as the main challenges, preventing them from even matching the quality and the efficiency of their planar counterparts. InP, being a direct bandgap semiconductor with very low SRV,5 is highly relevant as a platform for optical telecommunications,6 solar cells,3 and high-speed electronics.7 Growing high quality InP NWs has been the subject of intense research efforts.8−11 A variety of growth mechanisms have been explored, because each technique/growth mechanism has some advantages not covered by others. However, no report has so far shown diameter-independent crystal phase perfect InP NWs with proven high quantum efficiency (QE), low SRV, and perfectly controllable dimensions in taper-free NW arrays, which would be necessary for covering a wide range of device applications. In recent years, small diameter Au-seeded InP NWs with high crystalline quality and limited tapering have been demonstrated by chemical beam epitaxy,12 molecular beam epitaxy,13 and metal−organic vapor-phase epitaxy (MOVPE) via vapor−liquid−solid (VLS) growth.10 However, up-scaling diameters by using metal catalysts in epitaxial NW growth often © XXXX American Chemical Society
Received: June 9, 2014 Revised: August 4, 2014
A
dx.doi.org/10.1021/nl5021409 | Nano Lett. XXXX, XXX, XXX−XXX
Nano Letters
Letter
Figure 1. Electron micrographs of InP NWs grown by SA-MOVPE. (a−c) SEM images at 45° tilt view of the NWs grown at (a) 650, (b) 700, and (c) 730 °C under otherwise identical growth conditions. Insets show the corresponding top view SEM images. (d) Bright-field TEM image showing the morphology of a typical InP NW grown at 730 °C. (e−g) High-resolution TEM (HRTEM) images taken along [2̅110] zone axis from the top, middle, and bottom regions of the InP NW shown in (d), respectively. The HRTEM images show that the NW has a WZ structure and are free of stacking faults. (h) Selected area electron diffraction pattern of the NW shown in (d) confirming the WZ phase. Detailed structural information about the samples shown in (a) and (b) is provided in Supporting Information.
growth temperature increases to 700 °C (Figure 1b). The measured average diameter is 225 nm at the base with a tapering rate less than 7 nm/μm. When grown at 730 °C, the NWs show significantly improved uniformity in diameter (Figure 1c). They have hexagonal cross sections with minimal tapering with a diameter of 200 nm. Typically, tapering is related to direct vapor−solid growth on the NW sidewalls and it has been shown that different crystal structures lead to local changes in preferred sidewall facets and radial growth rate, leading to different final morphologies.24,25 Tapered InP NWs grown by SA-MOVPE have been reported with polytypic structure of zinc blende (ZB) and WZ.26 In the WZ segment, {-2110} facets are vertical to the (111)A substrate, while in the ZB segment, {-111} facets are inclined 19.5° from ⟨111⟩ vertical direction to the InP (111)A substrate. Therefore, a higher proportion of ZB segments will result in a more tapered morphology of the NWs. In our case, the NWs became less tapered and more uniform with increasing growth temperature, which could indicate that the crystal quality improves with increasing growth temperature. From high-resolution TEM examination along the length of the NW, all three NWs are found to have predominantly WZ phase with different densities of stacking faults (including ZB segments). The density of thin ZB segments is 110/μm for the sample grown at 650 °C and 22/μm for the sample grown at 700 °C (see Supporting Information). High-resolution images taken along a NW grown at 730 °C show a pure WZ crystal phase, evident by three typical high-resolution TEM images taken from the top, middle, and bottom sections of the NWs shown in Figure 1e−g, respectively. The WZ crystal phase was also verified by selected area electron diffraction (SAED) (Figure 1h). The degree of tapering observed with growth
and optical quality of the NWs room-temperature lasing from conventional guided modes is obtained. To grow the NW arrays, 30 nm of SiO2 was first deposited on (111)A InP substrates and electron beam lithography (EBL) was used to pattern these substrates with hexagonal arrays of circles, followed by chemical etching through the pattern to form the arrays of holes. The patterned substrates were then placed in a horizontal flow low-pressure MOVPE system, using trimethylindium (TMIn) and phosphine (PH3) as the precursors for NWs growth.23 Complete growth details are provided in Supporting Information. The NWs were structurally characterized using scanning electron microscopy (SEM) and transmission electron microscopy (TEM). The top panel of Figure 1 shows the SEM images of the ensemble NW arrays grown on a pattern with an opening diameter of 200 nm and a pitch size of 500 nm. NWs shown in Figure 1a−c were grown at 650, 700, and 730 °C, respectively. All other growth parameters were the same for these three NW samples (see Supporting Information). The average length of the NWs is 2.5, 4, and 5.7 μm for the arrays grown at 650, 700, and 730 °C, respectively. Therefore, we observed an increase of NW growth rate with increasing growth temperature, consistent with higher adatom diffusion length at higher temperatures. It is also interesting to note that the morphology of the NWs varies with growth temperature, as can be seen in the top panel of Figure 1. The NWs grown at 650 °C are tapered and have wide distribution in diameter; the measured average diameter is 240 nm at the base with a tapering rate up to 25 nm/μm. These NWs have triangular/hexagonal cross sections, as can be seen in the top view SEM image shown in the inset. The NWs become less tapered and more uniform in diameter as the B
dx.doi.org/10.1021/nl5021409 | Nano Lett. XXXX, XXX, XXX−XXX
Nano Letters
Letter
temperature is consistent with a higher density of ZB segments.26 Further details on the stacking-fault-free WZ InP NWs grown under this condition with NW diameter up to 600 nm can be found in the Supporting Information. To better control the NW crystal structure, other growth parameters including V/III ratio and molar fractions of the reactants were also investigated. The density of structural defects for NWs grown at 730 °C increases with increasing V/III ratio (by increasing the flow rate of PH3), which is in agreement with previous work,27 but also when increasing the molar fractions of both reactants simultaneously (see Supporting Information). The growth parameter dependence of the crystal structure of III−V NWs grown via VLS has been well studied with higher growth temperature, lower V/III ratio, and lower molar fractions of reactants, generally leading to a structural transition from ZB to WZ.28,29 We further note that all experimental works about gold-seeded III−V NWs have reported an evolution of the crystal structure or defect density with diameter, which was tentatively explained by various theoretical models.28 In addition, there have also been reports on the dependence of NW crystal structure on the array parameters (hole diameter and pitch size) in SA-MOVPE growth of III−V NWs, where larger diameters of NWs often result in fewer stacking faults.30 Surprisingly, our NWs do not display such dependence. Indeed the NWs grown at the optimal conditions, that is 730 °C and V/III ratio of 80, remain stacking-fault-free WZ for all array patterns, with diameters varying from 80 to at least 600 nm (grown out of a 200 nm opening for the latter, see Supporting Information). As current theoretical understanding have not predicted occurrence of such large pure WZ InP crystals, new theoretical development would be required, because large diameter, taper-free, and stacking-fault-free NWs are required in many applications such as photonic lasers for a better confinement of the photonic mode to the gain region.15,31 The optical quality of our NWs was assessed using photoluminescence (PL) and time-resolved PL (TRPL) measurements. These measurements were performed on asgrown NW arrays standing on the substrate. The NWs were excited using a 522 nm pulsed laser with a pulse duration of 300 fs and 20.8 MHz repetition rate. The minority carrier lifetime in the NWs was extracted from the single exponential fitting of TRPL decay curve that is measured by a time-correlated single photon counting (TCSPC) PL system. See Supporting Information for further details on the optical setup. Figure 2a shows the room-temperature PL spectra from the three NW arrays shown in Figure 1a−c. The spectra from all three samples show a peak with a shoulder at higher energy. We fit the spectra with two Gaussian peaks (see Supporting Information). The lower energy Gaussian peak is centered at 1.413 eV and the higher energy peak at 1.44 eV. Temperaturedependent PL spectra of InP NWs grown at 730 °C shows that the emission peak at 1.413 eV shifts to higher energy and the shoulder at 1.44 eV disappears with decreasing temperature (see Supporting Information). The shift in the peak position (the peak at 1.413 eV at room temperature) with temperature fits well to Varshni equation. Hence, we attribute this peak to the band edge emission from WZ InP NWs and the higher energy peak to the split off valence band.32,33 The room-temperature PL spectra in Figure 2a are not normalized, and the peak intensity is an indication of the optical quality of the NWs. The PL peak intensity increases as the growth temperature increases. The full width at half-maximum
Figure 2. (a) Typical room-temperature photoluminescence (PL) spectra from NWs grown at different temperatures. The spectra shown are not normalized and are taken at the excitation power density of 0.24 μJ/cm2/pulse. (b) Room-temperature minority carrier lifetimes for NWs grown at different temperatures. The lifetimes were extracted from room-temperature time-resolved PL (TRPL) measurements from several different NWs. The inset shows typical TRPL decays, measured at the peak emission wavelength, from NWs grown under the three different growth temperatures.
(fwhm) of the band edge emission peak after fitting the PL spectra with two Gaussian peaks are 27, 17, and 16 nm for the arrays grown at 650, 700, and 730 °C, respectively. A narrower line width of band edge emission is also an indication of better optical quality. The room-temperature minority carrier lifetimes, τ mc extracted by fitting the TRPL spectra with monoexponential decays are shown in Figure 2b. The TRPL spectra used for extracting τmc are shown in the inset of Figure 2b. τmc is defined as (1/τmc) = (1/τnr) + (1/τr), where τnr is the nonradiative lifetime and τr is radiative lifetime of carriers. NWs grown at 650 °C have τmc of 0.8 ns and this increases to 1.6 ns for NWs grown at 730 °C. The room-temperature τmc measured at low excitation power density is limited by τnr.34 Long τmc measured at room temperature is an indication of low bulk defect incorporation and good surface quality/low SRV. We measured τmc for different diameter NWs to estimate the SRV using the formula (1/τmc) = (1/τbulk) + (4SRV/d), where τbulk is the bulk minority carrier lifetime and d is NW diameter (see Supporting Information). From our fit, we estimate τbulk of 1.57 ns and SRV of 161 cm/s for the NWs grown at 730 °C. The SRV for our NWs is significantly lower than that of AlGaAs passivated GaAs NWs35 but comparable to the value measured for polytypic InP NWs grown by VLS using terahertz spectroscopy.9 QE gives a quantitative estimate of the optical quality of the NWs. The QE of NWs can be extracted from the variation in PL emission intensity with excitation power density, following the approach of Yoo et al.36 Here, power-dependent PL data was acquired at room temperature using a 532 nm continuous wave laser with power density ranging from 3.6 to 30.1 kW/ cm2. A 50× (NA = 0.55) objective lens was used to focus the laser beam to a spot size of 1 μm on single NWs transferred onto ITO-coated glass substrate. Figure 3a shows the PL C
dx.doi.org/10.1021/nl5021409 | Nano Lett. XXXX, XXX, XXX−XXX
Nano Letters
Letter
larger than that of the wafer and is comparable to that of the epilayer, which is an indication that they have excellent optical quality. It is worth noting that QEs of up to 25% have been evaluated for WZ InP nanoneedles at room temperature.37,38 We attribute our higher QE to the better crystal quality of our NWs. Because the growth temperature is higher in our NWs than in the nanoneedles case (400−450 °C), our NWs are also less susceptible to impurity incorporation.39 Impurities can act as nonradiative recombination centers, hence lower level of impurities will increase τnr of the NWs, resulting in a higher QE. The excellent structural and optical quality of the InP NWs grown at 730 °C for a broad range of diameters makes them well suited for nano-optoelectronic device applications, especially as lasers that require a larger diameter to support the optical mode and reduce lasing threshold. NWs transferred onto a low index substrate serve as a Fabry−Perot cavity; supporting guided modes along the NW and providing optical feedback due to the large refractive index contrast at NW ends. The SAE grown NWs have flat facets on top (Figure 1c), and the SEM image of a transferred NW show a relatively flat bottom end facet (see Supporting Information). Thus, these NWs provide both a good optical cavity and high quality gain medium. We performed optical pumping experiments at room temperature to test lasing behavior of these NWs. Figure 4
Figure 3. (a) Room -temperature PL spectra of a single InP NW from the array grown at 730 °C for five different excitation power densities. (b) Quantum efficiency (QE) as a function of excitation power density for NWs grown at 730 and 700 °C. QE for an InP wafer and InP epilayer (see Supporting Information for growth condition) are also shown for comparison. The data points are calculated from fitting the variation in PL intensity with excitation power density using the method described by Yoo et al.,36 and the lines are calculated using an alternative approach based on the ratio of minority carrier and radiative lifetimes, QE = τmc/τr (see Supporting Information for further details). The error bars shown are based on QE estimation from PL versus power density measurements on three individual NWs.
spectra at different excitation power densities for a single NW grown at 730 °C. The integrated PL intensity as a function of excitation power density is fitted using a model based on rateequations, and the QE is estimated from parameters derived from this fit (see Supporting Information). The estimated QE is shown as points as a function of excitation power density in Figure 3b. The error bar shows the variation in QE from measurements on three single NWs grown at the same condition. The QE of NWs can also be calculated by estimating the steady-state carrier density (N) within the NW at each excitation power density (see Supporting Information). The QE can be expressed as QE = τr−1/(τr−1 + τnr−1). We obtain τr from the estimate of N using the relation τr = 1/BN, where B is radiative recombination coefficient, and τnr can be estimated from the measured lifetime value. The QE calculated from τr and τnr as a function of excitation power density is shown by the solid lines in Figure 3b. The QE calculated using this approach agrees very well with the QE estimated by fitting the PL emission intensity with excitation power density. The QE for the NWs grown at 730 °C is 26% at an excitation power density of 3.6 kW/cm2 and increases to 50% at 30.1 kW/ cm2. The increase in the estimated QE with excitation power density can be attributed to a reduction in τr with increasing photogenerated carrier density. The QE data for the NWs grown at 700 °C is also shown in Figure 3b. The QE of the NWs is higher for higher growth temperature, consistent with the higher crystal quality of that sample. The QE for an InP epilayer and an InP wafer are also shown in Figure 3b for comparison. The wafer and the epilayer have QE of 5 and 24% at 3.6 kW/cm2 and 20 and 51% at 30.1 kW/ cm2, respectively. The QE of NWs grown at 730 °C is much
Figure 4. Room-temperature lasing from an optically pumped single InP NW grown at 730 °C. Emission intensity from a single NW as a function of pump fluence (filled circles) shows an “s” type behavior on the log−log scale. The gray region highlights the lasing threshold region. The fwhm of the NW emission spectrum (filled triangles) dramatically reduces from 70 nm to ∼1 nm at threshold. The inset shows the NW emission spectra on a log-scale at three different excitation power densities (110, 130, and 150 μJ/cm2/pulse); corresponding to spectra below threshold (green), at threshold (orange), and above threshold (brown), respectively. The right inset shows the optical microscope image of the NW emission above threshold. The rectangle highlights the outline of the InP NW. The two bright spots correspond to the NW ends. The interference pattern observed in the image is due to coherent emission from the NW ends.
shows the intensity of NW emission as a function of pump fluence from a single NW with diameter of 480 nm and length of 6.5 μm, lying on ITO-coated glass substrate. The data shows s-type behavior on the logarithmic plot of light input versus light output, typical of lasers. The shaded region indicates transition from spontaneous emission to stimulated emission. The transition is accompanied by a narrowing of the emission spectrum, as shown by the triangular points in Figure 4, where the fwhm of the NW emission reduces from 70 to ∼1 nm. The D
dx.doi.org/10.1021/nl5021409 | Nano Lett. XXXX, XXX, XXX−XXX
Nano Letters
Letter
Cincinnati for fruitful scientific discussions during the preparation of this manuscript. We also thank Drs. Fouad Karouta, Kaushal Vora, Naeem Shahid, and Animesh Basak for valuable technical support.
emission spectrum below threshold, at threshold, and above threshold is shown in the left inset of Figure 4. A narrow lasing peak at 1.38 eV appears on the shoulder of the broad spontaneous emission spectrum at pump fluence of 130 μJ/ cm2/pulse. The intensity of this narrow peak continues to increase with increasing pump fluence, whereas the intensity of the spectrally broad spontaneous emission is clamped, indicating the onset of lasing. The right inset in Figure 4 shows the optical image of the NW emission taken above threshold at pump fluence of 170 μJ/cm2/pulse with the excitation laser filtered out. Intense emission is observed from the NW ends (two bright spots in the image), which confirms that lasing is due to guided modes supported along the NW axis. The interference pattern observed in the image is a consequence of coherent laser emission from the NW ends and is yet again another characteristic of lasing. Excellent material and structural quality is an important requirement for achieving low threshold lasing. Room-temperature optical pumping measurements on other stacking-faultfree InP NWs grown at 730 °C with similar dimensions showed lasing at similar pump intensities. The lasing threshold of our stacking-fault-free InP NWs is lower than that for previously demonstrated GaAs NW lasers15,40 due to higher QE, significantly lower SRV and good morphology, that is, taperfree along NW axis and flat end facets. It should be noted that our NWs do not have any surface passivation. In comparison with previously reported InP NW lasers,38,41 which lase from whispering gallery modes, our lasers have smaller diameters, and lase from low order waveguide modes. In summary, we have reported the growth of stacking-faultfree WZ InP NWs by SA-MOVPE with wide range of diameters (80−600 nm). The high-quality InP NWs were grown epitaxially on patterned (111)A substrates with optimized growth parameters in terms of growth temperature and precursor flow rate. We note that maintaining a pure WZ phase up to very large diameters challenges current theoretical growth understanding. The NWs have an excellent QE of ∼50%, which is on par with InP epilayer. Finally, the structural uniformity and high QE of the InP NWs enables low threshold, room-temperature lasing from single NWs under optical pumping, which is promising for future device applications.
■
■
ASSOCIATED CONTENT
S Supporting Information *
Substrate processing and growth, Section 1. Additional TEM images, Section 2. PL and lifetime measurements, Section 3. QE estimation, Section 4. NW laser characterization, Section 5. This material is available free of charge via the Internet at http://pubs.acs.org.
■
REFERENCES
(1) Chau, R.; Doyle, B.; Datta, S.; Kavalieros, J.; Zhang, K. Nat. Mater. 2007, 6 (11), 810−812. (2) Tomioka, K.; Yoshimura, M.; Fukui, T. Nature 2012, 488 (7410), 189−192. (3) Wallentin, J.; Anttu, N.; Asoli, D.; Huffman, M.; Åberg, I.; Magnusson, M. H.; Siefer, G.; Fuss-Kailuweit, P.; Dimroth, F.; Witzigmann, B.; Xu, H. Q.; Samuelson, L.; Deppert, K.; Borgström, M. T. Science 2013, 339 (6123), 1057−1060. (4) Tchernycheva, M.; Messanvi, A.; de Luna Bugallo, A.; Jacopin, G.; Lavenus, P.; Rigutti, L.; Zhang, H.; Halioua, Y.; Julien, F. H.; Eymery, J.; Durand, C. Nano Lett. 2014, 14 (6), 3515−3520. (5) Rosenwaks, Y.; Shapira, Y.; Huppert, D. Appl. Phys. Lett. 1990, 57 (24), 2552−2554. (6) Bai, Y.; Bandyopadhyay, N.; Tsao, S.; Slivken, S.; Razeghi, M. Appl. Phys. Lett. 2011, 98 (18), 181102. (7) Flückiger, R.; Lövblom, R.; Alexandrova, M.; Ostinelli, O.; Bolognesi, C. R. Appl. Phys. Express 2014, 7 (3), 034105. (8) Paiman, S.; Gao, Q.; Tan, H. H.; Jagadish, C.; Pemasiri, K.; Montazeri, M.; Jackson, H. E.; Smith, L. M.; Yarrison-Rice, J. M.; Zhang, X.; Zou, J. Nanotechnology 2009, 20 (22), 225606. (9) Joyce, H. J.; Wong-Leung, J.; Yong, C.-K.; Docherty, C. J.; Paiman, S.; Gao, Q.; Tan, H. H.; Jagadish, C.; Lloyd-Hughes, J.; Herz, L. M.; Johnston, M. B. Nano Lett. 2012, 12 (10), 5325−5330. (10) Vu, T. T. T.; Zehender, T.; Verheijen, M. A.; Plissard, S. R.; Immink, G. W. G.; Haverkort, J. E. M.; Bakkers, E. P. A. M. Nanotechnology 2013, 24 (11), 115705. (11) Kitauchi, Y.; Kobayashi, Y.; Tomioka, K.; Hara, S.; Hiruma, K.; Fukui, T.; Motohisa, J. Nano Lett. 2010, 10 (5), 1699−1703. (12) Dalacu, D.; Mnaymneh, K.; Lapointe, J.; Wu, X.; Poole, P. J.; Bulgarini, G.; Zwiller, V.; Reimer, M. E. Nano Lett. 2012, 12 (11), 5919−5923. (13) Alouane, M. H. H.; Chauvin, N.; Khmissi, H.; Naji, K.; Ilahi, B.; Maaref, H.; Patriarche, G.; Gendry, M.; Bru-Chevallier, C. Nanotechnology 2013, 24 (3), 035704. (14) Dubrovskii, V. G. Nucleation Theory and Growth of Nanostructures; Springer: Berlin Heidelberg, 2014 (15) Saxena, D.; Mokkapati, S.; Parkinson, P.; Jiang, N.; Gao, Q.; Tan, H. H.; Jagadish, C. Nat. Photonics 2013, 7 (12), 963−968. (16) Bar-Sadan, M.; Barthel, J.; Shtrikman, H.; Houben, L. Nano Lett. 2012, 12 (5), 2352−2356. (17) Breuer, S.; Pfüller, C.; Flissikowski, T.; Brandt, O.; Grahn, H. T.; Geelhaar, L.; Riechert, H. Nano Lett. 2011, 11 (3), 1276−1279. (18) Jiang, N.; Gao, Q.; Parkinson, P.; Wong-Leung, J.; Mokkapati, S.; Breuer, S.; Tan, H. H.; Zheng, C. L.; Etheridge, J.; Jagadish, C. Nano Lett. 2013, 13 (11), 5135−5140. (19) Heurlin, M.; Hultin, O.; Storm, K.; Lindgren, D.; Borgström, M. T.; Samuelson, L. Nano Lett. 2014, 14 (2), 749−753. (20) Ikejiri, K.; Noborisaka, J.; Hara, S.; Motohisa, J.; Fukui, T. J. Cryst. Growth 2007, 298 (0), 616−619. (21) Yoshimura, M.; Nakai, E.; Tomioka, K.; Fukui, T. Appl. Phys. Express 2013, 6 (5), 052301. (22) Yoshimura, M.; Nakai, E.; Tomioka, K.; Fukui, T. Appl. Phys. Lett. 2013, 103 (24), 243111. (23) Gao, Q.; Tan, H. H.; Fu, L.; Parkinson, P.; Breuer, S.; WongLeung, J.; Jagadish, C. Conf. Optoelectron. Microelectron. Mater. Devices 2012, 45−46. (24) Bolinsson, J.; Caroff, P.; Mandl, B.; Dick, K. A. Nanotechnology 2011, 22 (26), 265606. (25) Ghalamestani, S. G.; Heurlin, M.; Wernersson, L.-E.; Lehmann, S.; Dick, K. A. Nanotechnology 2012, 23 (28), 285601. (26) Ikejiri, K.; Kitauchi, Y.; Tomioka, K.; Motohisa, J.; Fukui, T. Nano Lett. 2011, 11 (10), 4314−4318.
AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. Notes
The authors declare no competing financial interest.
■
ACKNOWLEDGMENTS The Australian Research Council is acknowledged for financial support. Access to facilities used in this work is made possible through the Australian National Fabrication Facility and Australian Microscopy and Microanalysis Research Facility. We are grateful to Professor Leigh Smith from University of E
dx.doi.org/10.1021/nl5021409 | Nano Lett. XXXX, XXX, XXX−XXX
Nano Letters
Letter
(27) Ikejiri, K.; Kitauchi, Y.; Tomioka, K.; Motohisa, J.; Fukui, T. Nano Lett. 2012, 12 (1), 524−525. (28) Caroff, P.; Bolinsson, J.; Johansson, J. IEEE J. Sel. Top. Quantum Electron. 2011, 17 (4), 829−846. (29) Lehmann, S.; Wallentin, J.; Jacobsson, D.; Deppert, K.; Dick, K. A. Nano Lett. 2013, 13 (9), 4099−4105. (30) Shapiro, J. N.; Lin, A.; Ratsch, C.; Huffaker, D. L. Nanotechnology 2013, 24 (47), 475601. (31) Zimmler, M. A.; Capasso, F.; Müller, S.; Ronning, C. Semicond. Sci. Technol. 2010, 25 (2), 024001. (32) Tuin, G.; Borgström, M.; Trägårdh, J.; Ek, M.; Wallenberg, L. R.; Samuelson, L.; Pistol, M.-E. Nano Res. 2011, 4 (2), 159−163. (33) De, A.; Pryor, C. E. Phys. Rev. B 2010, 81 (15), 155210. (34) Ahrenkiel, R. K. Carrier Lifetime in III−V Semiconductors. In Semiconductors and Semimetals; Academic Press: New York, 1993; Vol. 39, Chapter 2, pp 39−150. (35) Jiang, N.; Parkinson, P.; Gao, Q.; Breuer, S.; Tan, H. H.; WongLeung, J.; Jagadish, C. Appl. Phys. Lett. 2012, 101 (2), 023111. (36) Yoo, Y.-S.; Roh, T.-M.; Na, J.-H.; Son, S. J.; Cho, Y.-H. Appl. Phys. Lett. 2013, 102 (21), 211107. (37) Ren, F.; Ng, K. W.; Li, K.; Sun, H.; Chang-Hasnain, C. J. Appl. Phys. Lett. 2013, 102 (1), 012115. (38) Li, K.; Sun, H.; Ren, F.; Ng, K. W.; Tran, T.-T. D.; Chen, R.; Chang-Hasnain, C. J. Nano Lett. 2013, 14 (1), 183−190. (39) Fang, Z. M.; Ma, K. Y.; Cohen, R. M.; Stringfellow, G. B. Appl. Phys. Lett. 1991, 59 (12), 1446−1448. (40) Mayer, B.; Rudolph, D.; Schnell, J.; Morkötter, S.; Winnerl, J.; Treu, J.; Müller, K.; Bracher, G.; Abstreiter, G.; Koblmüller, G.; Finley, J. J. Nat. Commun. 2013, 4, 2931. (41) Wang, Z.; Tian, B.; Paladugu, M.; Pantouvaki, M.; Le Thomas, N.; Merckling, C.; Guo, W.; Dekoster, J.; Van Campenhout, J.; Absil, P.; Van Thourhout, D. Nano Lett. 2013, 13 (11), 5063−5069.
F
dx.doi.org/10.1021/nl5021409 | Nano Lett. XXXX, XXX, XXX−XXX
