GaN Multiple-Quantum-Wells on GaN

Oct 3, 2011 - The horizontal error bar is estimated from TEM and SIMS measurements. Single-wire horizontal devices have been fabricated to demonstrate...
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M-Plane Core Shell InGaN/GaN Multiple-Quantum-Wells on GaN Wires for Electroluminescent Devices Robert Koester,† Jun-Seok Hwang,‡ Damien Salomon,†,§ Xiaojun Chen,† Catherine Bougerol,‡ Jean-Paul Barnes,§ Daniel Le Si Dang,‡ Lorenzo Rigutti,|| Andres de Luna Bugallo,|| Gwenole Jacopin,|| Maria Tchernycheva,|| Christophe Durand,† and Jo€el Eymery†,* †

CEA-CNRS-UJF group, Nanophysique et Semi-conducteurs, SP2M, UMR-E CEA/UJF-Grenoble 1, INAC, Grenoble, F-38054, France CEA-CNRS-UJF group, Nanophysique des semi-conducteurs, Institut Neel-CNRS/Universite J. Fourier, 25 avenue des Martyrs, BP 166 38042 Grenoble cedex 9, France § CEA-Leti, Minatec-Campus, 17 rue des Martyrs, 38054 Grenoble Cedex 9, France Institut d'Electronique Fondamentale UMR CNRS 8622, Universite Paris Sud 11, 91405 Orsay Cedex, France

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‡

bS Supporting Information ABSTRACT: Nonpolar InGaN/GaN multiple quantum wells (MQWs) grown on the {1100} sidewalls of c-axis GaN wires have been grown by organometallic vapor phase epitaxy on c-sapphire substrates. The structural properties of single wires are studied in detail by scanning transmission electron microscopy and in a more original way by secondary ion mass spectroscopy to quantify defects, thickness (1 8 nm) and In-composition in the wells (∼16%). The core shell MQW light emission characteristics (390 420 nm at 5 K) were investigated by cathodo- and photoluminescence demonstrating the absence of the quantum Stark effect as expected due to the nonpolar orientation. Finally, these radial nonpolar quantum wells were used in room-temperature single-wire electroluminescent devices emitting at 392 nm by exploiting sidewall emission. KEYWORDS: Wires, nitrides, nonpolar, multiple quantum wells

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he past decade has seen tremendous progress in the control of nanowire (NWs) properties for their use as functional elements in nano- and opto-electronics.1 3 For nitride materials, which are usually used in UV-blue and white light-emitting diodes (LEDs), bidimensional (2D) growth problems encountered by the lattice mismatch, or the strain due to different thermal expansion coefficients are substantially reduced in the wire geometry. For example, GaN NWs can be grown on very large area substrates like silicon without any substrate bending and detrimental dislocations.4 NW LEDs also take advantage of 1D shape-related features such as enhancement of light extraction based on light guiding and polarization5 and the NW length may be used to integrate radial or longitudinal heterostructures that benefit from free surface strain relaxation. Molecular beam epitaxy (MBE) has demonstrated first the growth of catalyst-free GaN nanowires6 followed by the addition of longitudinal InGaN/ GaN heterostructures exhibiting promising optical properties7 for light-emitting devices.8 It has been shown that the piezoelectric field can be strongly reduced in c-axis longitudinal superlattices through efficient strain relaxation.9 It leads to a mitigation of the quantum-confined Stark effect (QCSE) by reducing the piezoelectric polarization along the polar c-axis growth direction. The fabrication of such c-longitudinal InGaN/GaN multiquantum-wells (MQWs) in NWs has been also demonstrated by metal organic hydride vapor phase epitaxy (MOHVPE) without catalyst to obtain NW-based LED devices,10 and more recently high-quality homogeneous c-axis GaN wires have been obtained by metal organic r 2011 American Chemical Society

vapor phase epitaxy (MOVPE), which is the standard industrial technique for nitride device fabrication.11,12 To take advantage of sidewall emission along the wire length, InGaN/GaN core shell nanowire heterostructures that enable LED and optically pumped lasers have been first demonstrated with Ni catalyst along the [1120]-axis (a-axis) having triangular cross-section with two equivalent semipolar {1101} planes and one polar {0001} plane.13 15 The electroluminescence of InGaN/GaN MQW emission along the semipolar (1011) and (1012) side facets of c-axis pyramidal-shaped wires16 as well as the cathodoluminescence of top polar (0001) facets have also been shown.12 The demonstration of purely nonpolar sidewall emission has not been reported yet, although this type of structure, free of electrostatic fields, improves the optical quantum efficiency for planar epitaxy. 17 19 In this work, we report on the epitaxial growth, structural, and optical characterization of InGaN/GaN nonpolar radial MQWs formed on sidewall {1100}-facets (m-plane) of c-axis GaN wires. The wire growth is performed using catalyst-free MOVPE on sapphire substrates. The thickness of QWs as a function of position along the wire, as well as structural defects are assessed by transmission electron microscopy. Time of flight-secondary ion mass spectrometry (ToF-SIMS) analysis performed on single Received: August 3, 2011 Revised: September 27, 2011 Published: October 03, 2011 4839

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Figure 1. (a) Overview SEM image of GaN wires surrounded by InGaN/GaN multiquantum wells. (b) Diagram of the core shell MQW structure. (c) Higher-magnification SEM image showing the morphology of the MQW grown on the m-plane facets of the wire.

wires provides an accurate estimation of In composition of the core shell MQWs. The optical properties of single wires are probed by cathodo- and microphotoluminescence measurements evidencing an intense emission from InGaN/GaN MQWs tunable in the 390 420 nm range and the absence of quantum confined Stark effect in agreement with the nonpolar QW orientation. Finally, we processed single wires containing an n-type core, an unintentionally doped MQW structure and a p-type external GaN barrier shell to fabricate nanoscale LEDs. These devices exhibit a strong room-temperature electroluminescence and constitute the first demonstration of catalyst-free single wire light-emitting diodes based on purely nonpolar InGaN/GaN MQWs in MOCVD wires. Self-assembled c-oriented GaN wires having lateral {1100} m-plane facets are grown by MOVPE on c-sapphire substrates using a thin in situ SiNx thin film predeposition.11 Grazing incidence X-ray diffraction and electron microscopy have shown that these wires have the conventional epitaxial relationships with sapphire corresponding to a 30° rotation of the two unit cells along the c-axis: [1100]GaN//[1210]Sapp and [0001]GaN//[0001]Sapp.11 The stem of the wires is grown at 1000 °C using trimethylgallium (TMG) and ammonia precursors with a low V/III ratio as well as silane addition to achieve n+-doping.11 The silane addition is switched off after about 25 μm length to grow an unintentionally doped GaN part (about 10 μm long) at the top of the wires. Finally, the top of the GaN wire sidewalls are coated with five unintentionally doped radial InGaN/GaN quantum wells. The MOCVD deposition of InGaN layers is performed under nitrogen at 400 mBar using triethylgallium (TEG) and trimethylindium (TMI) precursors. After the growth of the GaN core at high temperature (around 1000 °C), the temperature is decreased to about 720 760 °C to grow InGaN and then rapidly increased to 830 870 °C to deposit the GaN barrier using also the TEG

Figure 2. (a,b) High-angle annular dark-field (HAADF) STEM images of 20 80 s InGaN well growth time wires prepared by FIB. The wire cross-section is orthogonal to the length direction and shows an irregular hexagonal shape with m-plane facets (i.e., {1100} planes). These images give InGaN well-GaN barrier thicknesses about (a) 1.0 7.5 and (b) 5.2 11.6 nm. (c) HAADF STEM of a cross-section of a wire (sample b) parallel to the growth direction taken along the [1120] growth axis. (d) Variation of the GaN barrier and InGaN well thickness along a 1.5 μm long region of sample b. (e) Low-magnification dark-field TEM image taken with g = [1100] showing in bright defects originating from the first QW and propagating across the QWs. (f) High-resolution TEM image taken along the [1120] zone axis showing that the defects evidenced in panels c and e correspond to stacking faults (SFs).

precursor. This temperature is chosen as a compromise between achieving a high quality GaN barrier and preventing indium desorption/interdiffusion in the well. The QW growth time at 750 °C has been set to 20, 40, 80, and 120 s to vary the QW thickness (tQW= 1.25, 2.5, 5, 7.5 nm nominal values) at constant nominal barrier thickness of about tbarr= 10 nm (285 s growth time). The precise thickness values were then determined by electron microscopy measurements and checked by ToF-SIMS depth profiling. Figure 1a shows the SEM image of an ensemble of wires covered by radial MQWs, which grow on the nonpolar m-plane facets (see schematic in Figure 1b). As seen in Figure 1c, the sidewall m-plane facets of the wire are smooth at the upper part and their roughness gradually increases toward the bottom where small grains are observed.20 4840

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Figure 3. (a) Schematics of the ToF-SIMS experiment consisting in the measurement of a wire dispersed on a planar InGaN/GaN multiple quantum well reference. (b) Longitudinal scan along the (0004) Bragg peak of the planar MQW in epitaxy on c-sapphire substrate (open circles for measured data). The X-ray wavelength is 0.0688 nm and the best fit of the data is indicated by a blue line. (c) Visualization of the In+ ion depth mapping extracted from different parts of the wire and from the planar MQW substrate. Because of etching time, only the upper side of the wire is seen. Horizontal and vertical scales are indicated on the figure. (d) Estimation of the indium concentration as a function of the depth along the wire profile indicated by an arrow in (c).

In order to investigate the MQW heterostructures by highresolution scanning transmission electron microscopy (HRSTEM), wire cross sections orthogonal and parallel to the length direction have been prepared with a dual-beam focused ion beam setup and samples have been handled by micromanipulators. Microscopy experiments were performed using either a spherical aberration probe-corrected FEI Titan operated at 300 kV or a JEOL 4000Ex operated at 400 keV. Two transversal crosssectional scanning transmission electron microscopy images in high-angle angular dark-field (HAADF) mode of wires corresponding to 20 and 80 s QW growth times are shown in Figure 2a,b. The core/shell geometry is clearly observed and demonstrates a good structural quality of the InGaN/GaN superlattice grown on the m-plane sidewall facets. In the HAADF image, the dark (bright) contrast corresponds to the GaN core (InGaN shell). The thickness measurements for these two samples (20 and 80 s QW growth times) give, respectively, 1.0 (5.2) nm for the well and 7.5 (11.6) nm for the barrier with an error bar of about 0.2 nm. GaN/InGaN interfaces are sharp, although the interface transition from GaN to InGaN is more diffuse due to the temperature ramps as explained in the growth section. Figure 2c shows a HAADF-STEM image for the 80 s sample taken along the [1120] zone-axis of the MQWs. The well and barrier thicknesses measured at different position along about 1.5 μm of wire length are reported in Figure 2d. It indicates a thickness gradient of about 14% in this region both for the

barrier and well. Figure 2c shows also bright stripes with quite high density. The dark-field TEM image of the same sample taken with g = [1100] (see Figure 1e) shows that these defects originate from the first GaN/InGaN interface. The high-resolution image taken along the [1120] zone axis in a region including such a defect (see Figure 1f) evidence stacking faults (SFs). This observation is similar to what is commonly observed in the literature for planar growth21,22 where SFs also originate at the first InGaN well and are associated to Frank dislocations with 1/6 [2203] Burgers vector, which relaxes misfit strain along the c-direction. The 1/2 [0001] lateral shift in the (0001) planes adjacent to the SFs produces a cubic-like insertion inside the hexagonal phase. The formation mechanism of these SFs, still unresolved even in planar materials,22 is certainly driven by the strain accumulation in the heterostructures. Core shell InGaN/GaN MQWs have been studied by dualbeam ToF-SIMS from ION-TOF GmbH (Munster, Germany). A bismuth liquid metal ion gun is used to obtain ToF spectra and a 500 eV oxygen beam is used for depth profiling. The depth resolution and lateral resolution are estimated to be 2 and 200 nm, respectively. The detection limit for the In+ ions measured in this study is around 1016 atoms/cm3 (see details in ref 23). In this paper, to avoid dealing with matrix effects (i.e., non linear variation of ionization yields with concentration),23 the indium concentration in the core shell structure is obtained by comparison with a reference sample of similar concentration analyzed 4841

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Nano Letters in identical conditions. Wires are dispersed on a planar reference sample having five quantum wells grown on a GaN layer, itself deposited on c-plane sapphire. The depth profiling is performed on an area containing both the wire and the planar MQWs (see Figure 3a,c). The visualization of the structure is achieved by combining the 3D position of the scanned Bi ion beam with the time-of-flight data recorded at each pixel. This 4D data cube (x, y, z, δt) allows depth profiles to be extracted from different parts of the wire or from the reference substrate. The planar MQW sample has been studied previously by X-ray diffraction at the European Synchrotron Radiation Facility (ESRF). The (0002) and (0004) symmetrical Bragg reflections have been measured within the standard θ-2θ geometry at a wavelength of 0.0688 nm and fitted24 to get the In-concentration as well as an estimation of the thicknesses. The analysis of the (0004) reflection shown in Figure 3b corresponds to an In-content of 18.5 ( 0.5% in the 2.75 ( 0.1 nm well and a barrier thickness of about 8.3 ( 0.1 nm (the error bars being determined from the (0002) and (0004) scan simulations). A cross-section of the 3D In ion mapping is shown in Figure 3c for the 40 s wire, the top of the figure corresponds to the wire and the bottom to the 2D layer. The z-axis is related to the depth and a gradient in the MQW thickness is measured along the c-axis of the wire. Several In+-ion profiles at different lateral positions have been performed. The depth profile at the middle of the length of the MQW (see the arrow in Figure 3c) is given in Figure 3d. The In+-ion number has been transformed into atomic concentration as a function of depth using to the known concentration and thickness of the reference planar layer. The In concentration in the QWs estimated by the time-of-flight secondary ion mass spectrometry is about 16% ((2%). The damping of the concentration profile is attributed to the loss of depth resolution induced by sputtering of a nonplanar object (see the depth resolution of the 2D MQW, which is well preserved over the 5 layers), and also to the fact that deeper QWs are no longer parallel to the sputtered surface. It can be checked that the integral of the In-concentration oscillation (i.e., the average composition in the well) is conserved when small enough region of interests are taken (i.e., smaller than about 800 nm). By measuring the profile at several positions along the full length of the MQW, a 14% variation of the period is determined in complete agreement with the gradient measured by HR-TEM. The larger thickness at the top of the wires results in two contributions. First to the antitapering of the wire stem for which silane addition does not completely prevent the lateral extension of the wires and second from the core shell coating without silane. The materials incorporation mechanisms of these two steps probably involve different precursor collection in the gas phase at the top c-plane, edges, and sidewalls. As shown later, the In-concentration of the MQW is chosen to be low enough to prevent emission heterogeneities due to chemical demixing and strong variability along the MQW length. This MQW gradient must nevertheless be taken into account for device optimization: the core shell covering length must be adjusted as a function of the targeted emission bandwidth. The local measurements of the light emission of single wires (20s, 750 °C wells) are studied by CL in Figure 4. At the top of the wire (see position P1 in Figure 4a,c), the CL intensity spectrum exhibits an intense peak around 396 nm with a full width at half-maximum (fwhm) of about 20 nm. As shown by the side and top views (respectively Figure 4e,g) of the CL-mappings, this emission comes from the part of the wire covered by the radial InGaN/GaN MQWs (for the spectra all along the wire see

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Figure 4. Low-temperature (5 K) cathodo-luminescence (CL) measurements of single core shell 5 InGaN/GaN MQWs wires grown by MOVPE. Wells are grown at 750 °C for 20 s (see text). CL spectra taken (a) at top of the wire covered by the MQW and (b) at the stem of the MQW (see position P1 and P2 in the SEM image (c)). (d f) Side-view CL mapping at 350, 396, 560 nm corresponding respectively to the n-doped part of the wire stem, MQW, and undoped part of the GaN under the MQW. (g) Top-view CL-mapping of a wire measured at the MQW emission wavelength (396 nm).

Figure S1, Supporting Information). No optical signature coming from the axial MQWs is observed at the top of these wires within our experimental resolution. A small near band edge (NBE) emission of unintentionally doped GaN is measured at 358 nm (Figure 4a). This peak is shifted to smaller wavelengths (350 nm, see position P2 in Figure 4b,d) in the n-doped stem of the wire due to Si doping and band-filling (Burnstein-Moss) effects, the freeelectron concentration being estimated to be about 8  1019 cm 3 from homoepitaxial GaN layers measurements.25 The top of the wire also exhibits a weak “yellow band” (YB) emission centered at 560 nm (see Figure 4a,f). The interpretation of this contribution is still controversial in the literature; in our samples, the YB is only observed in the unintentionally doped part, whereas it is absent in the heavily n-doped part. The same behavior has been measured in similar bare n-u GaN wires,11 so that the influence of MQW 4842

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Figure 5. Low-temperature (5 K) μ-photoluminescence (μ-PL) measurements of single core shell InGaN/GaN MQWs wires grown by MOVPE. (a) Spectra obtained for samples with five MQWs grown at different growth temperatures for 80 s well growth duration. (b) Position of the μ-PL emission as the function of the well growth temperature. (c) Spectra for several well growth times (20 to 120 s) grown at 750 °C. (d) Analysis of the μ-PL emission wavelength. Symbols correspond to different wire measurements. The well thickness measured for the 20 and 80 s samples by TEM (see Figure 2) is used to determine the duration-thickness relationship. The horizontal error bar is estimated from TEM and SIMS measurements.

growth can be excluded. It suggests that heavy silicon incorporation limits the occurrence of radiative defect centers.26 The CL-analysis has been extended to μ-PL measurements at low temperature to study in more detail the MQW emission properties. Figure 5a shows the evolution of the PL MQW emission wavelength at 5 K as a function of the InGaN well growth temperature between 720 and 760 °C for nominal 5 nm InGaN/ 10 nm GaN MQWs (80 s well growth time), that is, for a relatively large well thickness, where confinement can be neglected. It is well-known in planar MQWs that indium incorporation is very sensitive to the growth temperature and that a decrease of temperature increases the In-incorporation inside the well. As shown in Figure 5b, the emission wavelength of the wires depends linearly on the InGaN growth temperature with a slope of about 0.75 nm/K. Supposing no change of well thickness and comparing with homogeneous InGaN layers materials,27 a variation of the In incorporation (from 16 to 24%) in the MQWs may explain the low-temperature emission red-shift from 390 to 430 nm for a growth temperature varying from 760 to 720 °C. In parallel, the maximum of intensity of the peaks is decreased by more than 2 orders of magnitude and the peaks are broadened with an increase of the fwhm from 100 to 175 meV. As already reported in the literature, the increase of the fwhm of the MQW peak emission as a function of the indium mole fraction in the QW may be due to an increase of structural and chemical defects (such as clustering).28 In a second series of experiments, the InGaN well growth temperature has been set to 750 °C and the barrier thickness to about 10 nm to study the shift of the PL MQW emission wavelength at 5 K as a function of the InGaN well growth time: 20, 40, 80, 120 s (Figure 5c). The μ-PL emission wavelength is strongly decreased at short QW growth time (392 nm for 20 s), whereas it reaches the asymptotic value of

about 412 nm for longer QW growth times (120 s). To confirm this behavior and estimate the wavelength distribution for the different wires, several μ-PL measurements have been performed and reported in Figure 5d. The abscissa thickness-axis takes into account the TEM measurements and the error bars correspond to the profile gradients along the wire length. The In-concentration for the QWs presented in Figures 5c can be roughly estimated from the InGaN PL emission energies. The MQW emission related to the asymptotic value of Figure 5d, which excludes quantum confinement effect, occurs at about 412 nm at 5 K. It roughly corresponds to an In-concentration of about 20% considering X-ray calibration measurements performed at room temperature on 2D c-axis InGaN alloy layers.27,29 The discrepancy with Tof-SIMS measurements can be explained by local InGaN alloy thickness and composition fluctuations leading to localization phenomena, which lower the effective PL transition energy. The fwhm of the PL-peaks is almost constant at a value of about 160 meV, which is consistent with a constant In-concentration in the QWs. Consequently, the blue shift of the emission wavelength should be directly attributed to a decrease of the QW thickness with decreasing growth time and therefore to quantum confinement.30 32 The transition energy E in a QW is given in first approximation by E ≈ Eg + Eel + Eh1 eFd where Eg is the band gap of the strained well, Ee1 (Eh1) are the first confinement energy levels of the electrons (holes) and the last term represents the influence of the quantum-confined Stark effect (QCSE) as a function of the piezoelectric field F and the QW thickness d. This contribution is particularly important for wells grown along the c-axis and becomes dominant for thick QWs, which results in a linear decrease of the QW transition energy with QW thickness. Without piezoelectric field, the QCSE term tends to zero and the transition energy only depends on the electron and hole energies, 4843

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Figure 6. (a) SEM view of an electroluminescent single wire device with (b) its related optical microscopy image showing light emission. (c) Current/voltage characteristic at 300 K. (d) Electroluminescence spectrum at 300 K for 15 μA current injection.

which are almost constant for thick QWs due to the absence of quantum confinement. The PL measurements of the core/shell MQW emission shown in Figure 5c,d are therefore consistent with a very small (or zero) piezoelectric field resulting from the m-plane growth on the GaN wire sidewalls. The absence of QCSE has also been confirmed by power-dependent CL measurements. The electron hole pair concentration in the QWs increasing with excitation power should screen the piezoelectric field and a blue shift of the emission peak position should be observed. The CL spectra for excitation powers varied over more than 3 decades (from 0.019 to 60 kW/cm2) have been measured (Figure S2, Supporting Information) and no blue shift of the MQW emission peaks has been observed, which excludes the presence of a significant piezoelectric field in the QWs that separates the carriers in usual c-plane structures. The main advantage to grow MQWs on equivalent crystallographic m-plane sidewalls is to prevent emission heterogeneities that occur in triangular crosssection wires having two semipolar {1101} planes and one polar {0001} plane.33 Single-wire horizontal devices have been fabricated to demonstrate the electroluminescence properties of such radial

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InGaN/GaN MQW heterostructures grown on the wire facets (see also in ref 34 an example of photodetectors based on similar MQW heterostructures). The InGaN QWs are grown at T = 750 °C for t = 20 s corresponding to 16% In content and a 1 nm thickness of QWs. A radial p-GaN shell grown at 920 °C is added on the MQWs using biscyclopentadienylmagnesium as a precursor source followed with an annealing step performed at 750 °C under N2 during 20 min to activate Mg-dopants. Wires are then dispersed on Si/SiO2 substrate after an ultrasonic bath in ethanol. To make the electrical contact of these thick wires (diameter in the range of 400 600 nm), they are first encapsulated in a 500 nm thick spin-on glass layer (H-silsesquioxane (HSQ ) transformed into silicon oxide by annealing). Dry etching is then performed to uncover the wire surfaces and metal contact areas are defined by e-beam lithography. It is worth mentioning that the n-type GaN core could be contacted without intermediate etching steps, as the MQW system is located only in the upper part of the wire. The contact to the n- (p-) GaN is obtained by the deposition of Ti/Al/Ti/Au (Ni/Au) metals respectively (Figure 6a). Current voltage (I V) characteristics measured at room temperature and in ambient air show typical n-p diode behavior with sharp current increase at around 9 V in forward bias (Figure 6c). As shown in Figure 6d, the single-wire devices produce a strong electroluminescence at room temperature for positive bias (>9.75 V) at 392 nm with a fwhm of 25 nm This electroluminescence comes from the upper part of the single wires close to the p-contact in the region of MQW deposition (Figure 6b). Two smaller contributions measured respectively at 445 and 491 nm may be attributed to the emissions of localized states in radial wire and to longitudinal MQWs at the top or the wires. The relative intensities of these contributions depend heavily on the excitation conditions, which can be advantageously used in devices to tune the wavelength35 and to obtain blue emission.4 The MOVPE growths of nonpolar InGaN/GaN multiple quantum wells on the {1100} sidewalls of c-GaN wires have been demonstrated on c-sapphire substrates and carefully characterized to analyze their optical properties in terms of piezoelectric field and thickness dependence. STEM experiments were performed to measure the thickness of the wells (1 8 nm) and barriers (7.5 11 nm), and stacking faults have been observed at the first interface of the MQW. The In-composition (∼16%) of single wire wells grown at 750 °C has been estimated from ToF-SIMS measurements with a planar MQW reference. This technique allows also the MQW thickness gradient along the wire to be estimated at ∼14%, which is in agreement with local STEM observations. The core shell MQW sidewall emission (390 420 nm at 5 K) has been investigated by cathodo- and photoluminescence and showed nonmeasurable quantum Stark effect in agreement with the nonpolar orientation. One application of this radial coating that takes advantage of sidewall emission along the wire length has been illustrated by a roomtemperature single-wire electroluminescent devices emitting at 392 nm, which is promising for high-efficiency nanowire-based LED devices.4

’ ASSOCIATED CONTENT

bS

Supporting Information. Cathodoluminescence spectra along the length of a wire (Figure S1). Power dependent cathodoluminescence measurements (Figure S2). This material is available free of charge via the Internet at http://pubs.acs.org.

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’ AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected].

’ ACKNOWLEDGMENT The authors thank J. Dussaud and M. Lafossas for the MOCVD technical support, M. Terrier for the FIB sample preparation, M. den Hertog for her help in STEM experiments, E. Bellet-Amalric for her participation in X-ray simulations, D. Sam Giao and B. Gayral for photoluminescence measurements and M. Azize for his participation in first MOVPE growths. The French CRG ESRF beamline is also thanked for its contribution to this study. Research was partially supported by the French government ANR-05-NANO-016 Filemon35, ANR-08-NANO031 Bonafo, ANR RTB, and Carnot Eclairage projects. ’ REFERENCES (1) Li, Y.; Qian, F.; Xiang, J.; Lieber, C. M. Mater. Today 2006, 9, 18–27. (2) Thelander, C.; Agarwal, P.; Brongersma, S.; Eymery, J.; Feiner, L. F.; Forchel, A.; Scheffler, M.; Riess, W.; Ohlsson, B. J.; Goesele, U.; Samuelson, L. Mater. Today 2006, 9, 28–35. (3) Pauzauskie, P. J.; Yang, P. Mater. Today 2006, 9, 36–45. (4) Bavencove, A.-L.; Salomon, D.; Lafossas, M.; Martin, B.; Dussaigne, A.; Levy, F.; Andre, B.; Ferret, P.; Durand, C.; Eymery, J.; Le Si Dang, D.; Gilet, P. Electron. Lett. 2011, 47, 765–767. (5) Yan, R. X.; Gargas, D.; Yang, P. D. Nat. Photonics 2009, 3, 569–576. (6) Yoshizawa, M.; Kikuchi, A.; Mori, M.; Fujita, N.; Kishino, K. Jpn. J. Appl. Phys. 1997, 36, L459–L462. Calleja, E.; Sanchez-Garcia, M. A.; Sanchez, F. J.; Calle, F.; Naranjo, F. B.; Munoz, E. Phys. Rev. B 2000, 62, 16826–16834. (7) Kikuchi, A.; Kawai, M.; Tada, M.; Kishino, K. Jpn. J. Appl. Phys. 2004, 43, L1524–L1526. (8) Guo, W.; Zhang, M.; Banerjee, A.; Bhattacharya, P. Nano Lett. 2010, 10, 3355–3359. (9) Wang, C.-Y.; Chen, L.-Y.; Chen, C.-P.; Cheng, Y.-W.; Ke, M.-Y.; Hsieh, M.-Y.; Wu, H.-M.; Peng, L.-H.; Huang, J. J. Opt. Express 2008, 16, 10549–10556. (10) Kim, H.-M.; Cho, Y.-H.; Lee, H.; Kim, S.; Ryu, S. R; Kim, D. Y.; Kang, T. W.; Chyng, K. S. Nano Lett. 2004, 6, 1059–1062. (11) Koester, R.; Hwang, J. S.; Durand, C.; Le Si Dang, D.; Eymery, J. Nanotechnology 2010, 21, 015602. (12) Bergbauer, W.; Strassburg, M.; K€olper, Ch; Linder, N.; Roder, C.; L€ahnemann, J.; Trampert, A.; F€undling, S; Li, S. F.; Wehmann, H.-H.; Waag, A. Nanotechnology 2010, 21, 305201. (13) Qian, F.; Li, Y.; Gradecak, S.; Wang, D.; Barrelet, C. J.; Lieber, C. M. Nano Lett. 2004, 4, 1975–1979. (14) Qian, F.; Gradecak, S.; Li, Y.; Wen, C.-Y.; Lieber, C. M. Nano Lett. 2005, 5, 2287–2291. (15) Qian, F.; Li, Y.; Gradecak, S.; Park, H.-G.; Dong, Y.; Ding, Y.; Wang, Z. L.; Lieber, C. M. Nat. Mater. 2008, 7, 701–706. (16) Frajtag, P.; Hosalli, A. M.; Bradshaw, G. K.; Nepal, N.; ElMasry, N. A.; Bedair, S. M. Appl. Phys. Lett. 2011, 98, 143104. (17) Waltereit, P.; Brandt, O.; Trampert, A.; Grahn, H. T.; Menniger, J.; Ramsteiner, M.; Reiche, M; Ploog, K. H. Nature 2000, 406, 865–868. (18) Okamoto, K.; Ohta, H.; Chichibu, S. F.; Ichihara, J.; Takasu, H. Jpn. J. Appl. Phys. 2007, 9, L187–L189. € ur, U € .; Baski, (19) Ni, X.; Lee, J.; Wu, M.; Li, X.; Shimada, R.; Ozg€ A. A.; Morkoc-, H.; Paskova, T.; Mulholland, G.; Evans, K. R. Appl. Phys. Lett. 2009, 95, 101106. (20) In these growths, wires are partially covered by the MQW heterostructure. The explanation of the covering mechanism is beyond the scope of this article. A patent is pending and this description will lead to a specific publication.

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