Structural Properties of Wurtzite InP–InGaAs Nanowire Core–Shell

Letter pubs.acs.org/NanoLett

Structural Properties of Wurtzite InP−InGaAs Nanowire Core−Shell Heterostructures Magnus Heurlin,*,† Tomaš Stankevič,‡ Simas Mickevičius,‡ Sofie Yngman,† David Lindgren,† Anders Mikkelsen,† Robert Feidenhans’l,‡ Magnus T. Borgström,† and Lars Samuelson† †

The Nanometer Structure Consortium, Department of Physics, Lund University, Box 118, 221 00 Lund, Sweden Niels Bohr Institute, University of Copenhagen, Universitetsparken 5, 2100 Copenhagen, Denmark

‡

S Supporting Information *

ABSTRACT: We report on growth and characterization of wurtzite InP−In1−xGaxAs core−shell nanowire heterostructures. A range of nanowire structures with different Ga concentration in the shell was characterized with transmission electron microscopy and X-ray diffraction. We found that the main part of the nanowires has a pure wurtzite crystal structure, with occasional stacking faults occurring only at the top and bottom. This allowed us to determine the structural properties of wurtzite In1−xGaxAs. The InP−In1−xGaxAs core−shell nanowires show a triangular and hexagonal facet structure of {1100} and {101̅0} planes. X-ray diffraction measurements showed that the core and the shell are pseudomorphic along the caxis, and the strained axial lattice constant is closer to the relaxed In1−xGaxAs shell. Microphotoluminescence measurements of the nanowires show emission in the infrared regime, which makes them suitable for applications in optical communication. KEYWORDS: Nanowire, core−shell, wurtzite, X-ray diffraction, InP, InGaAs

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of InP−In1−xGaxAs in the core−shell nanowire geometry where the nanowire crystal structure was intentionally tuned to WZ for a high morphological quality of the heterointerface. To study the material properties of WZ In1−xGaxAs, we designed WZ InP−In1−xGaxAs core−shell nanowires where we varied the composition of the In1−xGaxAs shell by tuning the growth parameters in metal organic vapor phase epitaxy (MOVPE). The WZ InP nanowire cores were grown in the particle assisted growth mode using sulfur as a dopant,9,16 after which In1−xGaxAs shells, adopting the crystal structure of the core, were grown. The nanowires were synthesized from nano imprint lithography (NIL) defined Au particles on (1̅1̅1̅)B InP substrates. The NIL defined pattern had a period of 400 nm with aligned columns of Au particles where each column is slightly offset from a rectangular arrangement. The imprint process, developed by Obducat AB, includes an intermediate polymer stamp (IPS) and soft press technology,17 which were used to pattern 2″ wafers. From these, smaller samples were cleaved, typically 5 mm × 5 mm in size, for use in growth experiments carried out in a horizontal flow MOVPE reactor. A total flow of 6 l/min at 100 mbar pressure with H2 as the carrier gas was used, and the precursors used for growth were trimethylindium (TMIn), trimethylgallium (TMGa), AsH3, PH3, H2S, and HCl. HCl is used in the growth process to avoid tapering of the core nanowires and thus improve their material

nGaAs is in combination with InP, an important material system for applications such as bipolar transistors,1 solar cells,2 and optical communication.3 The emission wavelength from In1−xGaxAs, which is typically combined with InGaAsP to realize strained quantum well stacks,3 is tunable in the 1.3−1.6 μm region, which is suitable for optical communication due to a low dispersion and low attenuation in silica glass optical fibers. Semiconducting nanowires provide an interesting approach toward engineering new types of heterostructures where, for example, efficient strain relaxation via the free surface in axially defined nanowire structures reduces the need for lattice matched material combinations.4 Also, thin core−shell nanowires show enhanced strain accommodation properties since both the core and the shell can be strained.5 However, in a core−shell nanowire geometry with larger core diameter, the heterointerface requirement for lattice matching becomes similar to those for thin films,6 which can lead to formation of dislocations.7 Thus, lattice matched, or close to lattice matched, material combinations will become important to realize high quality nano structures with low defect density. In addition to the new device geometries offered by nanowires, they also introduce a new degree of freedom compared to planar structures since the crystal structure can be tuned to either wurtzite (WZ) or zinc blende (ZB) by using, for example, dopant flows,8,9 tuning the group V flow,10,11 or adjusting the seed particle size.12 The crystal structure has proven to be an important factor when forming core−shell nanowire heterostructures since it can affect both defect formation13,14 and strain relaxation in the shell.15 In this Letter, we have investigated the strain, composition, and luminescence © XXXX American Chemical Society

Received: December 22, 2014 Revised: February 12, 2015

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Figure 1. SEM images of three samples with varying values of xTMGa corresponding to (a) 0.40, (b) 0.55, and (c) 0.70. The nanowire morphology is similar for xTMGa below 0.65, but at xTMGa = 0.70, substantial accumulation of growth material close to the top can be observed. In addition, nanowire bending is present for xTMGa = 0.70. The scale bar in panel c is valid for all three images, which were acquired with a 30° tilt angle with respect to the substrate normal.

Figure 2. TEM images in (a) bright field and (b) dark field of an InP−InGaAs core−shell nanowire with xTMGa = 0.50. Panel c shows the corresponding diffraction pattern of the nanowire. The 1̅1̅1 beam used in panel b has been marked. In panel d, a high-resolution image of the center of the nanowire is shown, which displays the pure WZ structure existing in the main part of the nanowire.

structure, high-resolution TEM and dark field imaging were combined. The EDX measurements were made by evaluating three nanowires for each In1−xGaxAs composition. Each nanowire was in turn evaluated by three spectra acquired close to the nanowire base, middle, and top. The positions close to the base and top were within 400 nm from each end of the broken off nanowire. Reciprocal space mapping (RSM) was done in the grazing incidence geometry using the 12.4 keV synchrotron radiation at the I811 beamline19 at MAXII synchrotron in Sweden. A fivecircle diffractometer was used to perform rocking scans of the Bragg peaks by measuring the scattered intensity with a 2D position sensitive detector. In-plane and out-of-plane reciprocal space maps were produced showing Bragg peaks of the InP cores and In1−xGaxAs shells. To precisely determine the lattice constants, RSMs were made using the ZB InP substrate peak as reference. At least two symmetric peaks were measured to eliminate possible effects of tilting. The optical properties were investigated with a μ-PL setup at a temperature of 4 K. The laser used for excitation consisted of a frequency doubled Nd:YAG emitting at 532 nm. The measurements were performed on single nanowires that were mechanically removed from the growth substrate and transferred to an Au coated Si wafer before measurement.

quality.18 The InP−In1−xGaxAs core−shell nanowire structure was fabricated in two separate steps. First, the core InP nanowire was grown using molar fractions χTMIn = 8.2 × 10−6, χPH3 = 18.5 × 10−3, χH2S = 1.7 × 10−5, and χHCl = 1.0 × 10−5 at a temperature of 395 °C, after which the nanowire samples were removed from the MOVPE reactor and wet etched to remove the Au−In alloy particle.14 In the second MOVPE step, an InP−In1−xGaxAs shell was grown on the InP nanowire cores at a temperature of 600 °C using molar fractions χTMIn = 4.9 × 10−6 and χPH3 = 37.0 × 10−3 for the InP layer and χTMGa = 2.6 × 10−6 − 11.3 × 10−6, χAsH3 = 2.2 × 10−3, and χTMIn = 4.9 × 10−6 for the In1−xGaxAs layer. The growth time was 1 min for the InP shell and 20 min for the In1−xGaxAs shell. Since the Au particle was removed, axial growth is limited, and most of the In1−xGaxAs material deposits on the nanowire side facets. The resulting nanowire core−shell structure consisted of an InP core approximately 120 nm in diameter and an In1−xGaxAs shell with a thickness that varied between 10 and 40 nm. The nanowire samples were studied by means of scanning electron microscopy (SEM), transmission electron microscopy (TEM), X-ray diffraction (XRD), and microphotoluminescence (μ-PL). For the TEM investigation, a JEOL 3000F 300 kV microscope equipped with scanning-TEM capability and an energy dispersive X-ray (EDX) detector for measuring the chemical composition were used. To determine the crystal B

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Nano Letters A total of eight samples were grown with varying xTMGa = χTMGa/(χTMIn + χTMGa) from xTMGa = 0.35 to xTMGa = 0.70 in steps of 0.05. SEM images of the samples with xTMGa = 0.40, 0.55, and 0.70 are shown in Figure 1 (see Supporting Information Figure S1 for images of all compositions). The nanowire morphology of samples grown with xTMGa from 0.35− 0.65 is similar, but at low values of xTMGa, some nanowires are observed to bend. This indicates the presence of strain between the core and shell in combination with a nonuniform shell deposition.20 If the shell growth is inhomogeneous and strain is present between the core and shell, the nanowires bend due to the increased compressive/tensile force on one side compared to the other. Since the amount of strain depends on the composition of the shell, this effect should be largest when there is a large lattice mismatch between core and shell. For high Ga content structures, that is, xTMGa = 0.70, material also accumulates on the (11̅ 1̅ )̅ top facet of the nanowires, as seen in Figure 1, panel c. In addition, the shell deposition is more inhomogeneous, which causes stronger bending of the nanowires as compared to lower values of xTMGa. The nanowire cross-section varies between a triangular and hexagonal shape without any observable trend between different compositions of the shell. Some of the triangular shaped nanowires are rotated by 60 degrees with respect to each other, which is most likely caused by the formation of twin planes in the early stages of the nanowire core growth.21 The core−shell nanowires were investigated using TEM to determine any local variations in their structural properties. Figure 2, panels a and b show bright and dark field imaging of a nanowire with xTMGa = 0.50. By selecting electrons that diffract through the 1̅1̅1 beam (Figure 2c), we can identify the parts of the nanowire that contain ZB segments, which show up as bright areas in Figure 2, panel b. We find that only two small regions contain any ZB contribution, namely the top and bottom parts of the nanowire. In the bottom part, typically 200 nm long, these originate from the original InP nanowire core where there seems to be a transient before the crystal structure turns into pure WZ. EDX investigations show that the top part, typically 20 nm long, consists of In1−xGaxAs and is thus grown during the shell deposition step. The remainder of the nanowire consists of a pure WZ crystal without stacking defects, as shown in Figure 2, panel d. Complementary to the TEM characterization, the XRD measurement will give information on the average structural properties of nanowires on a 0.5 mm × 5 mm area given by the X-ray beam footprint. Thus, the number of nanowires characterized is on the order of 107. In these measurements, the nanowires are still standing on the substrate, which allows us to assess the state of strain without concern for externally imposed mechanical stress. Radial and axial strain in both core and shell can be found from the in-plane and out-of-plane RSMs. For precise determination of the strain, three WZ Bragg peaks were used: (101̅1), (1̅011), and (202̅1) as well as ZB peaks (111̅), (11̅1), (131,) and (002) for reference. Additional information on the nanowire morphology can be found from the in-plane RSMs, since the shape of the Bragg peaks represents a Fourier transform of the shape of the object under investigation. From the (101̅1) WZ reflection shown in Figure 3, panels a and b, it can be seen that there are 12 distinct crystal truncation rods (CTRs), streaks emerging from the central peak, which are caused by faceting of the nanowires. Usually only six CTRs are visible for hexagonal nanowires,22 which correspond to the six facets.

Figure 3. In-plane RSMs of the (101̅0) WZ peak of the sample with (a) xTMGa = 0.65 and with (b) xTMGa = 0.40. Twelve streaks emerging from the central peak show the presence of the two types of facets: {101̅0} facets and {1100} facets for the core−shell nanowire structure. (c) Out-of plane RSMs of the (1̅01l) rod measured using the sample grown with xTMGa = 0.65. The central vertical feature is the CTR originating from stacking faults present in both the nanowires and in parasitic substrate growth. Three main structures present in the core are clearly visible: ZB, TW, and WZ. Additional GaAs peaks originate from parasitic planar growth, and the In1−xGaxAs shell is visible in form of extensions of the peaks toward larger values of momentum transfer. The dashed lines show directions along which relaxed ternary alloys are expected. Qx is perpendicular to the (101̅0) plane, and Qz is perpendicular to the (0001̅) plane; values are normalized to the reciprocal unit cell of ZB InP.

Presence of an additional family of the CTRs, rotated by 30 degrees, can be caused by the triangular morphology of the shell. Thus, two types of facets are present: {101̅0} and {1100}. The presence of the shell is also reflected in the asymmetry of the Bragg peak, since it is extended toward larger values of the momentum transfer Qx (Figure 3a). The shoulder is caused by the smaller a-axis lattice constant, which corresponds to In1−xGaxAs with a higher Ga content than necessary for a lattice-matched structure. Determination of the radial strain is not straightforward due to a small separation of the Bragg peaks. In a previous study,23 it was shown that in the case when the shell is fully strained around the core, the scattered intensity is distributed among several peaks along the CTRs, which makes them highly asymmetric. Here, in contrast, the In1−xGaxAs Bragg peak appears as a localized shoulder on one side of the InP peak. This means that the shell is relaxed in the radial direction. Strain along the c-axis and polytypism can be investigated from the out-of-plane RSMs, shown in Figure 3, panel c. The rod (1̅01l) is shown containing the Bragg peaks of the main structures present in the samples: ZB, twinned zinc blende (TW), and WZ. The units on the axes are normalized to the reciprocal lattice of the substrate ZB system. The central vertical feature is the CTR originating from stacking faults both in the nanowires and parasitic planar growth on the substrate. Three main InP peaks can be seen along each of the rods corresponding to ZB, TW, and WZ structural polytypes. Offset toward the origin from the main peaks are two GaAs peaks showing presence of a plastically relaxed GaAs crystal in the ZB and TW phases. The small size of the peaks and full relaxation suggest that they come from parasitic substrate growth between the nanowires. The actual In1−xGaxAs shell is visible as an extension of the WZ and TW peaks toward the larger values of momentum transfer. The absence of any In or C

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only be seen in samples with xTMGa = 0.65 and xTMGa = 0.7. The signal from the shell can be seen as a small extension of the Bragg peak for the samples with xTMGa = 0.35−0.6 due to the smaller lattice mismatch. By combining the knowledge obtained from the in-plane and out-of-plane RSMs we can state that for all nanowires, the axial stress, caused by lattice mismatch, is the dominating source of strain in the core and shell. In the following paragraph, we use this model to determine the composition and strain in both the cores and the shells. Strain in the core and shell as well as composition of the shell was evaluated using Vegard’s law and anisotropic elastic theory. The elastic constants of InAs and GaAs used for calculations were reported by Larsson et al. and Strauch et al., respectively,26,27 while the relaxed lattice constants of WZ InAs and GaAs were measured by Kriegner et al. and Biermanns et al., respectively.28,29 Relaxed lattice constants of InP nanowires were recently reported by Kriegner et al.30 Detailed derivations and the calculation results can be found in the Supporting Information. From the calculations, we can see that gallium concentration xGa = Ga/(Ga + In) in the In1−xGaxAs shell varies from 0.43−0.67 for the samples grown with xTMGa equal to 0.35 and 0.7, respectively, and average axial strain in the core and shell ranges from negative (−0.15%) to positive (0.56%), respectively. Results of the calculations are shown in Figure 5 where panels a and c show measured radial and axial lattice constants with respect to the calculated composition values xGa. Relaxed InP and relaxed In1−xGaxAs lattice constants calculated using Vegard’s law are shown with black and red solid lines, respectively. Because of the different a/c ratios for InP and InGaAs, the radial and axial lattice constants coincide at different values of x. This can be seen from the intersection of the solid lines in Figure 5. From the radial lattice constant a, we can see that the InP core exhibits compressive strain when xGa is larger than 0.5, which then saturates and decreases at xGa = 0.67. On the other hand, the radial lattice constant of the shell closely follows Vegard’s law with a small positive strain for xGa < 0.45 and negative strain for xGa > 0.45 due to Poisson’s effect. Geometrical phase analysis from TEM images of the sample with xGa = 0.67 also shows that the radial lattice constant changes abruptly at the core−shell interface, which indicates a sharp compositional transition (see Supporting Information Figure S2). The lattice constants along the c-axis (Figure 5c) are equal for the core and shell to an extent of ±1%. From panel c in Figure 5, it is shown that the c-axis lattice constant takes values in between the relaxed In1−xGaxAs and InP, closer to In1−xGaxAs, meaning that the nanowire crystals overall axial lattice constant is closer to the relaxed shell than the core. Larger relaxation occurs only when xGa > 0.6. This is possibly related to the fact that these samples show the largest amount of random tilting of the nanowires, as seen from the SEM images (Figure 1c) and spread of the Bragg peaks shown in Figure 4. Thus, bending of the nanowires may be considered as a mechanism for strain relaxation. From the EDX measurements (see Supporting Information Figure S3), we similarly observe that xGa could be tuned from 0.40 at xTMGa = 0.35 to 0.64 at xTMGa = 0.70. These measurements additionally provide high spatial resolution of individual nanowires. We could observe a small spread in the measured composition within both single nanowires and between different nanowires. The standard deviation between the nine measurements performed for each sample was between 0.02−0.04. No trend in xGa that would indicate

Ga rich In1−xGaxAs WZ peak indicates that no phase segregation occurs in this ternary system opposite to what has been found for AlInP24 and AlInAs25 shells. From Figure 3, it is clear that the In1−xGaxAs region does not follow the dashed relaxation line but stretches horizontally. The fact that the In1−xGaxAs peak is vertically aligned with the InP core peak can only be seen from the samples with the largest lattice mismatch, where the two peaks are separated. However, provided that there is no axial relaxation for the samples with the largest lattice mismatch, we can reasonably assume that also nanowires with smaller lattice mismatch will be fully strained. c-axis lattice constants of the InP core and the lattice matched In1−xGaxAs shell can be found from the relative vertical position Qz of the WZ peak with respect to the ZB reference peak. Average Qz was taken from the (1̅011) and (101̅1) Bragg peaks to eliminate possible effects of tilting. Gaussian profiles were fitted to determine the positions of the peaks, and the lattice constants were calculated taking the bulk ZB lattice constant aInP ZB = 5.8687 Å as a reference. In the radial direction, the lattice constants of the core and the shell are different, hence a sum of two Gaussians was fitted to each intensity profile along the Qx direction. RSMs of the (202̅1) Bragg peaks for the whole range of samples are shown in Figure 4. In all RSMs, the central InP WZ

Figure 4. Out-of-plane RSMs of the (202̅1) Bragg peaks for samples with different xTMGa values (given in the corner of each figure). The main InP peak is extended diagonally due to the mosaicity caused by random tilting of the nanowires. The shell appears as a small extension at smaller momentum transfer compared to the InP peak for the samples with low xTMGa. With a larger value of xTMGa, the lattice constant of the shell decreases; therefore, the In1−xGaxAs peak reappears on the right side of the InP peak and completely separates when xTMGa is 0.65 and 0.7. Qx is perpendicular to the (101̅0) plane, and Qz is perpendicular to the (0001̅) plane; values are normalized to the reciprocal unit cell of ZB InP.

peak is visible with a diagonal streak, which is due to the mosaic spread caused by random tilting of the nanowires. The spread is smallest for samples with xTMGa = 0.5 and xTMGa = 0.6 and largest for the sample with xTMGa = 0.7, which suggests that the tilting is caused by the strain in the nanowires, which SEM images showed could result in nanowire bending (Figure 1). The signal from the shell is most visible for samples with the largest lattice mismatch, and separate In1−xGaxAs peaks can D

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Figure 5. Radial and axial lattice constants of the InP core and In1−xGaxAs shell plotted with respect to the calculated Ga concentration (a and c) and Ga concentration measured with EDX (b and d). Black horizontal lines show the lattice constants of the relaxed InP in the WZ phase. Red diagonal lines represent the relaxed WZ lattice constants of the In1−xGaxAs shell calculated using Vegard’s law. a denotes the lattice parameter in the radial direction, that is, perpendicular to the nanowire growth direction, while c denotes the lattice parameter in the axial direction, that is, along the nanowire growth direction.

nm with a second smaller peak at 1315 nm. Additionally, we see a small peak emerging at low excitation intensities at approximately 1375 nm, which at higher intensities, merges with the peak at 1215 nm. We believe that the different peaks could originate from local composition variations in the nanowires, for example, corresponding to the In1−xGaxAs segment growing on the top (1̅1̅1̅)B facet and from the side facets of the NW. A span from 1315−1415 nm would, in bulk ZB, approximately correspond to a compositional variation of 0.06 in xGa.31 This order of variation is similar to that observed in the EDX measurements where a standard deviation of xGa between 0.02 and 0.04 was observed. The emission wavelengths are, however, in the region where present optical communication operates, which could potentially be a future application area. In conclusion, we have investigated the structural, compositional, and optical properties of WZ InP−In1−xGaxAs core− shell nanowires. We have shown that the nanowires are axially strained due to lattice mismatch between the core and the shell. The average axial lattice constant is in between that of the relaxed InP and relaxed In1−xGaxAs, which means that both core and shell are strained. Since the a/c ratio is different for InP and In1−xGaxAs, no completely lattice matched structure can be obtained. Instead, strain is present either along the aaxis, c-axis, or both a- and c-axes. Optical emission around wavelengths important for optical communication give these types of nanostructures a potential future application area as nanoscale light emitters in the infrared region.

different diffusion lengths of the In and Ga species on the nanowires surface could be found, that is, that the base would always have a higher xGa as compared to the top. Measured lattice constants by XRD are plotted in Figure 5, panels b and d with respect to the composition found from the EDX measurements. It is clear from Figure 5, panel b that the radial lattice constants of the shell are offset from the relaxed values by a constant offset for the entire range. This shows discrepancy between the XRD and EDX measurements. To determine the optical properties of WZ In1−xGaxAs, we fabricated an InP−In1−xGaxAs−InP core−shell nanowire heterostructure where the surface of the In1−xGaxAs is protected by a higher bandgap InP cladding layer. Figure 6 shows μ-PL data from an InP−In1−xGaxAs−InP sample where xTMGa = 0.6, and thus xGa = 0.52 according to our XRD measurements. The main emission peak appears around 1415

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ASSOCIATED CONTENT

S Supporting Information *

SEM images, EDX measurements, geometrical phase analysis, detailed derivations of strain equations, and calculation results for all samples. This material is available free of charge via the Internet at http://pubs.acs.org.

Figure 6. Intensity dependent μ-PL measurement of a single InP− In1−xGaxAs−InP nanowire with xTMGa = 0.6. A laser intensity of 100% corresponds to 10 W/cm2. E

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(13) Ghalamestani, S. G.; Heurlin, M.; Wernersson, L.-E.; Lehmann, S.; Dick, K. A. Nanotechnology 2012, 23 (28), 285601. (14) Kawaguchi, K.; Heurlin, M.; Lindgren, D.; Borgström, M. T.; Ek, M.; Samuelson, L. Appl. Phys. Lett. 2011, 99 (13), 131915. (15) Conesa-Boj, S.; Boioli, F.; Russo-Averchi, E.; Dunand, S.; Heiss, M.; Rüffer, D.; Wyrsch, N.; Ballif, C.; Miglio, L.; Morral, A. F. i. Nano Lett. 2014, 14 (4), 1859−1864. (16) van Weert, M. H. M.; Helman, A.; van den Einden, W.; Algra, R. E.; Verheijen, M. A.; Borgström, M. T.; Immink, G.; Kelly, J. J.; Kouwenhoven, L. P.; Bakkers, E. P. A. M. J. Am. Chem. Soc. 2009, 131 (13), 4578−4579. (17) Eriksson, T.; Yamada, S.; Venkatesh Krishnan, P.; Ramasamy, S.; Heidari, B. Microelectron. Eng. 2011, 88 (3), 293−299. (18) Borgström, M.; Wallentin, J.; Trägårdh, J.; Ramvall, P.; Ek, M.; Wallenberg, L. R.; Samuelson, L.; Deppert, K. Nano Res. 2010, 3 (4), 264−270. (19) Grehk, T. M.; Nilsson, P. O. Nucl. Instrum. Methods Phys. Res., Sect. A 2001, 467−468, 635−638. (20) Mohan, P.; Motohisa, J.; Fukui, T. Appl. Phys. Lett. 2006, 88 (1), 013110. (21) Zou, J.; Paladugu, M.; Wang, H.; Auchterlonie, G. J.; Guo, Y.-N.; Kim, Y.; Gao, Q.; Joyce, H. J.; Tan, H. H.; Jagadish, C. Small 2007, 3 (3), 389−393. (22) Mariager, S. O.; Sørensen, C. B.; Aagesen, M.; Nygård, J.; Feidenhans’l, R.; Willmott, P. R. Appl. Phys. Lett. 2007, 91 (8), 083106−083106. (23) Stankevič, T.; Mickevicius, S.; Nielsen, M. S.; Ciechonski, R.; Vescovi, G.; Bi, Z.; Kryliouk, O.; Gundlach, C.; Mikkelsen, A.; Samuelson, L.; Feidenhans’l, R. J. Appl. Crystallogr. 2015, 48, DOI: 10.1107/S1600576715000965. (24) Sköld, N.; Wagner, J. B.; Karlsson, G.; Hernán, T.; Seifert, W.; Pistol, M.-E.; Samuelson, L. Nano Lett. 2006, 6 (12), 2743−2747. (25) Tomioka, K.; Yoshimura, M.; Fukui, T. Nature 2012, 488 (7410), 189−192. (26) Larsson, M. W.; Wagner, J. B.; Wallin, M.; Håkansson, P.; Fröberg, L. E.; Samuelson, L.; Wallenberg, L. R. Nanotechnology 2007, 18 (1), 015504. (27) Strauch, D. New Data and Updates for IV-IV, III-V, II-VI, and IVII Compounds, Their Mixed Crystals, and Diluted Magnetic Semiconductors; Springer: Berlin, Heidelberg, 2011; Vol. 44D. (28) Kriegner, D.; Panse, C.; Mandl, B.; Dick, K. A.; Keplinger, M.; Persson, J. M.; Caroff, P.; Ercolani, D.; Sorba, L.; Bechstedt, F.; Stangl, J.; Bauer, G. Nano Lett. 2011, 11 (4), 1483−1489. (29) Biermanns, A.; Breuer, S.; Davydok, A.; Geelhaar, L.; Pietsch, U. Phys. Status Solidi RRL 2011, 5 (4), 156−158. (30) Kriegner, D.; Wintersberger, E.; Kawaguchi, K.; Wallentin, J.; Borgström, M. T.; Stangl, J. Nanotechnology 2011, 22 (42), 425704. (31) Goetz, K. H.; Bimberg, D.; Jürgensen, H.; Selders, J.; Solomonov, A. V.; Glinskii, G. F.; Razeghi, M. J. Appl. Phys. 1983, 54 (8), 4543−4552.

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Author Contributions

M.H. and T.S. contributed equally. The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. M.H. performed the nanowire growth, SEM and TEM characterization, and coordinated the writing of the manuscript. T.S., S.M., and S.Y. performed the X-ray diffraction measurements. T.S. and S.M. did the X-ray diffraction data analysis. D.L. performed the optical measurements. R.F., A.M., M.T.B., and L.S. supervised the project. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was partly performed within the Nanometer Structure Consortium at Lund University (nmC@LU) and was supported by the EU project NWs4Light under Grant No. 280773, the Danish National Research Council DANSCATT, the Swedish Research Council (VR), the Swedish Foundation for Strategic Research (SSF), the Swedish Energy Agency, and the Knut and Alice Wallenberg Foundation. Part of this research was carried out at beamline I811, MAX-lab synchrotron radiation source, Lund University, Sweden.

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ABBREVIATIONS NW, nanowire; ZB, zinc blende; WZ, wurtzite; TW, twinned zinc blende; RSM, reciprocal space map; SEM, scanning electron microscopy; TEM, transmission electron microscopy; EDX, energy-dispersive X-ray spectroscopy; μ-PL, microphotoluminescence; XRD, X-ray diffraction; CTR, crystal truncation rod

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REFERENCES

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DOI: 10.1021/nl5049127 Nano Lett. XXXX, XXX, XXX−XXX