NANO LETTERS
Phase Segregation in AlInP Shells on GaAs Nanowires
2006 Vol. 6, No. 12 2743-2747
Niklas Sko1 ld,† Jakob B. Wagner,‡ Gunnel Karlsson,‡ Tania Herna´n,† Werner Seifert,† Mats-Erik Pistol,† and Lars Samuelson*,† Solid State Physics/The Nanometer Structure Consortium, Lund UniVersity, Box 118, SE-221 00 Lund, Sweden, and Polymer and Materials Chemistry/The Nanometer Structure Consortium, Box 124, SE-221 00 Lund, Sweden Received July 21, 2006; Revised Manuscript Received October 17, 2006
ABSTRACT We have studied morphology and phase segregation of AlInP shells on GaAs nanowires. Photoluminescence measurements on single core− shell nanowires indicated variations in the shell composition, and phase segregation was confirmed by cross-sectional scanning transmission electron microscopy on 30 nm thin slices of the wires. It was discovered that Al-rich domains form in the 〈112〉 directions where two {110} shell facets meet during growth. We propose that the mechanism behind this phase segregation is a variation in the chemical potential along the circumference of the nanowire together with a difference in diffusion lengths for the different growth species. From the morphology of the core and the shell, we conclude that the side facet growth is temperature dependent forming {112} facets at low growth temperature and {110} facets at high growth temperature.
Recently, core-shell nanowires have attracted attention as a new approach for nanowire-based devices. Core-shell nanowire heterostructures have been used as light-emitting diodes,1 to form a Fabry-Perot optical cavity around a lasing nanowire,2 and for surface state passivation to enhance the nanowire emission efficiency.3,4 Shell structures have been designed to supply an electron- or hole-gas to the core which allows for high mobility transistors,5,6 and oxidized shells have been used as gate-dielectric for coaxially gated nanowire field-effect transistors.7 By choice of a ternary shell material, the band gap and lattice constant can be controlled; there have however been indications that such ternary shells are not necessarily homogeneous.3 Growth on nonplanar substrates such as the wire side facets is different from planar growth since the chemical potential, µ, varies over the surface. Such variations in µ are known to cause phase segregation for alloy growth8,9 and have been used to form vertical quantum-wells10 in V-grooves on planar substrates. The variation in µ along the circumference of the wire is therefore likely to create corresponding variations in the composition of a ternary shell. As a homogeneous epitaxial shell is important for most applications, studies of these phase segregation effects can lead to improvements of future core-shell nanowire devices. The nanowires were grown by low-pressure metal-organic vapor phase epitaxy (MOVPE) at 100 mbar pressure. GaAs * Corresponding author. E-mail:
[email protected]. † Solid State Physics. ‡ Polymer and Materials Chemistry. 10.1021/nl061692d CCC: $33.50 Published on Web 11/02/2006
© 2006 American Chemical Society
wire growth was seeded from size-selected Au aerosols, deposited on (111)B GaAs substrates. The samples were placed on a graphite susceptor in a laminar flow MOVPE reactor cell and annealed at 580 °C in ambient AsH3/H2 for 10 min prior to growth in order to desorb any surface oxide. After the sample was annealed, the temperature was ramped down to 450 °C and the nanowire growth commenced when trimethylgallium (TMG) was supplied to the reactor cell. The precursor molar fractions were as follows: AsH3, 5 × 10-4; TMG, 2 × 10-5. At this low temperature, growth on the side facets is surface reaction limited and the Au assisted growth dominates. The wires were grown for 5 min after which the TMG flow was switched off and growth was terminated. The temperature was subsequently ramped up to 610 °C to overcome the kinetic hindrance of the surface reactions. At this temperature AsH3 was switched off and PH3 followed by trimethylaluminum (TMA) was switched on in order to grow a thin AlP spacer to prevent interdiffusion over the core-shell interface. After 2 s, trimethylindium (TMI) was switched on as well, and the AlInP shell was grown for 1 min. The shell was grown lattice matched to the core, and the precursor molar fractions were as follows: PH3, 1.5 × 10-2; TMA, 1 × 10-5; TMI, 2 × 10-5 (all values in a H2 flow of 6 L/min). A more detailed discussion on the core-shell growth can be found in ref 3. Sample characterization was then performed by scanning electron microscopy (SEM), photoluminescence (PL), and cross-sectional scanning transmission electron microscopy (STEM). For the PL studies, wires were transferred to a
patterned Au surface where they could be studied individually. The measurements were performed at a temperature of 5 K using the 458 nm line from an Ar+ laser for excitation. The emission was collected by an optical microscope, dispersed through a spectrometer and detected by a liquid N2 cooled charge coupled device (CCD) camera. For the STEM studies the wires were embedded in an epoxy resin and 30 nm thin slabs were cut using an ultramicrotome with a diamond knife. The cut was not perpendicular to the wires but at a 2° angle. Each slab therefore consisted of wire slices from different heights on the wires, slices on one edge of the slab originated from the top of the wires while slices on the opposite edge originated from the base. More information on the ultramicrotomy can be found in ref 11. The slabs were then placed on a lacey carbon-filmed Cu grid for STEM studies using a JEOL 3000F TEM operated at 300 keV. SEM images showed that the wires grew perpendicular to the (111)B substrate surface. Previous studies have shown that a large number of stacking faults makes the crystalline structure of 〈111〉B wires alternate from zinc blende to wurtzite.12-14 In this paper we will, for simplicity, only use the zinc blende notation for the crystal orientation. In SEM top views, wires with shell appeared to have {110} side facets in contrast to wires without shell which had {112} facets; thus a 30° rotation of the side facets takes place as the shell is grown. This rotation could also be seen for wires with GaAs shells grown at high temperature and is therefore not material specific. The assignment of the wire side facets was done using the {110} cleavage planes of the substrate as a reference; the results were later confirmed with highresolution TEM. Figure 1a shows a SEM image of the GaAs-AlInP core-shell nanowires with the substrate tilted 45° toward the e-beam. Panels b, c, and d of Figure 1 show SEM images of three wires seen from above. Wire 1b is a GaAs wire grown Au assisted at 450 °C without any shell. Wire 1c is a GaAs-AlInP core-shell wire where the core was grown at 450 °C and the shell at 610 °C, and 1d is a GaAs-GaAs “core-shell” wire where both the core and shell material are GaAs. The circular feature in the center of the wires is the Au particle, and the hexagonal shape originates from the side facets. Because of tapering, the core-shell wires display two hexagons, an inner dark one representing the shell side facets at the top of the wire and an outer bright one representing the shell side facets at the base. The core cannot be seen in these images but should have the same morphology as the wire without shell in Figure 1b. For the wire without shell, the contrast only allows us to determine the side facets at the base. The images show that as the shell is grown at 610 °C, the side facets rotate from {112} (Figure 1b) to {110} (Figure 1c,d). In reflection high-energy electron diffraction (RHEED) studies of molecular-beam epitaxy on GaAs substrates, it has been observed by No¨tzel at al. that there is a reversible onset of the formation of {110} facets on a (112) surface in the temperature range of 550-590 °C.15 This corresponds well with our results. PL measurements on the core-shell nanowires showed a single peak from the core at 1.51 eV; no substantial strain 2744
Figure 1. (a) SEM image of GaAs-AlInP core-shell nanowires, substrate tilted 45° toward e-beam. The scale bar is 1 µm. (b-d) SEM top views. (b) GaAs wire grown Au assisted at 450 °C. (c) GaAs wire with AlInP shell grown at 610 °C. (d) GaAs-GaAs “core-shell” wire where both the core and shell materials are GaAs. The shell is grown at 610 °C. Scale and orientation are the same for (b-d); the scale bar is 100 nm.
related shift in the PL peak was found. This indicates lattice matching between core and shell or, in other words, an average shell composition of approximately Al0.5In0.5P. At the bottom of Figure 2 the PL spectrum of a typical wire can be seen. The shell enhances the emission efficiency of the GaAs core at least 2 orders of magnitude by removing the surface states. The shell luminescence was on the other hand very weak. When lattice matched to GaAs, AlInP has an indirect band gap and the charge carriers predominantly recombine in the core. In order to study the shell separately, Nano Lett., Vol. 6, No. 12, 2006
Figure 2. PL spectra of a GaAs-AlInP core-shell nanowire (bottom) and an AlInP nanotube (top) at 5 K. The inset shows a SEM top view of a nanotube. Scale bar is 100 nm.
the core was therefore removed using a selective wet etch (140:3:1 H2O/NH3/H2O2) forming AlInP nanotubes. At the top of Figure 2 the spectrum from a typical AlInP nanotube (the shell spectrum) can be seen. A SEM top view of a tube is seen in the inset. Multiple peaks between 1.9 and 2.1 eV in the shell spectrum indicate that the shell is not homogeneous. When lattice matched to GaAs, the shell is expected to have an indirect band gap of 2.34-2.36 eV.16 The peaks at lower energies therefore indicate that local band gap potential minima exist. These potential minima are due to phase segregation forming In-rich domains, possibly in combination with sublattice ordered domains. Sublattice
ordering alone can however not explain the entire shift as it is only expected to give a band reduction of approximately 200 meV.17 STEM images of the core-shell nanowire cross sections showed that the core is tapered and that side facet growth occurs even at 450 °C, which is consistent with other studies.18,19 Parts a, b, and c of Figure 3 show high-angle annular dark field scanning TEM (HAADF-STEM) images of cross sections from the top, the center (1.2 µm from the top), and the base (2.1 µm from the top) of the wire; the total wire length was 2.5 µm. The core diameters are 25, 45, and 50 nm, respectively. In Figure 3d a diffraction pattern from a cross section illustrates the orientation of the images. The images show that the wires form {110} facets at the Au interface (Figure 3a) which are then converted gradually (Figure 3b) into {112} facets (Figure 3c) due to side facet growth at low temperature (tapering). As the shell is grown at a higher temperature, it again forms {110} facets in agreement with the results of No¨tzel at al.15 Due to the different contrast of lighter and heavier elements, where lighter elements allow a higher transmission, the inhomogeneity of the shell can be studied in parts a-c of Figure 3. As Al is a lighter element than In, Al-rich domains are seen as dark areas in the dark field images. These Al-rich domains can be seen as a shadow around the core due to the AlP spacer but also as dark lines in the 〈112〉 directions where two {110} facets meet during growth. Close to the top of the wire, the Al-rich domains go all the way in to the core. Further down along the wire these domains generally do not start until at some distance away from the core although they sometimes display a weak neck going all the way in to the core (in Figure 3c three such necks are seen). It seems as if segregation does not begin until the shell has developed
Figure 3. (a-c) Cross-sectional HAADF-STEM images of GaAs-AlInP core-shell nanowires. (a) Approximately 50 nm from the top of the wire. (b) Approximately 1.2 µm from the top of the wire. (c) Approximately 2.1 µm from the top of the wire. Al-rich domains have darker contrast. Scale and orientation are the same for all three images as indicated in (a). The scale bar is 50 nm. (d) Diffraction pattern of a cross section. (e) Schematic of the partial phase segregation near the base of the wires. Al-rich domains are marked black. Nano Lett., Vol. 6, No. 12, 2006
2745
{110} facets. At the top of the wire the core already displays this orientation, and therefore the Al-rich domains go all the way in to the core (Figure 3a). Further down along the wire the core side facets are a mixture of {110} and {112} facets. It is not until after some shell growth that well-defined {110} facets are formed and the Al can start aggregating at the interface between these facets (Figure 3b,c). X-ray energy dispersive spectrometry (XEDS) point measurements gave a composition of Al0.6In0.4P at the interface between the {110} facets in the 〈112〉 directions while the rest of the shell had a composition of Al0.5In0.5P. These were the only phase segregated domains that could be confirmed within the accuracy of the XEDS ((2%) although the HAADF-STEM images showed weak contrast variations in the growth direction of the shell as well (the origin of these oscillations is not yet identified). The In-rich domains, indicated by the PL measurements, could not be resolved by the XEDS measurements of the nanowire cross sections and no indication of sublattice ordering was observed in diffraction patterns of core-shell nanowires in side view. The experiments showed that a GaAs nanowire forms {110} facets as it is grown from the Au particle. Lateral growth however forms {112} facets at low temperatures but {110} facets at higher temperatures. Due to tapering during the nanowire growth, the core will therefore display {112} facets at the base (which is the part seen in a top view projection) while the shell displays {110} facets. According to Wulff’s rule the equilibrium shape of the shell represents the low specific surface energy facets,20 or in other words µ{110} < µ{112} during shell growth. Variations in µ determine the driving force in epitaxy and gradients in µ induce surface diffusion fluxes toward regions with lower µ, in this case the {110} facets. The adatom flux jj is described by the Nernst-Einstein relation jj ) -
nDs ∇µ kBT
where n is the surface density of adatoms, Ds is the surface diffusion coefficient, kB is Boltzmann’s constant, and T is the temperature. Since the growth species have different diffusion coefficients, phase segregation will take place. Diffusion coefficients are determined by the bond strengths,21 and as the Al-P bonds are stronger than the In-P bonds,22 In adatoms have longer diffusion lengths than Al adatoms. In will therefore move away from high µ regions more easily and there will be an aggregation of Al at the interface between the {110} facets in the 〈112〉 directions. µ not only is orientation dependent but also is affected by the surface curvature, surface stress, and entropy of mixing.8,9 As the surface is transformed from a mixture of {110} and {112} facets to pure {110} facets, the profile becomes sharper (the transition from a dodecagon to a hexagon) causing a larger gradient in µ.23 Hence larger phase segregation takes place as the pure {110} facets have been developed. The increase in the µ gradient is then to some extent counterbalanced by the surface stress and the decrease in entropy that the phase segregation implies. Figure 3e shows a schematic of the 2746
phase segregation. The core is white and the shell is gray; Al-rich domains are marked black. The mechanism behind the phase segregation as described here is the same as has been suggested in ref 8 as a model for the formation of selfordered nanostructures during growth on nonplanar surfaces such as vertical quantum wells formed in V-grooves. In summary, we have studied the morphology and phase segregation of AlInP shells on GaAs nanowires using crosssectional STEM. We have shown that for a nanowire with a ternary shell, the shell composition varies along the circumference due to a corresponding variation in µ together with a difference in diffusion lengths for the different growth species. For an AlInP shell this phase segregation forms Alrich domains in the 〈112〉 directions where two {110} facets meet during growth. According to XEDS measurements the Al-rich domains had a composition of Al0.6In0.4P compared to the Al0.5In0.5P composition of the rest of the shell. PL measurements however indicated that more In-rich domains also exist. We have also shown that the morphology of the side facet growth is temperature dependent forming {112} facets at a low growth temperature and {110} facets at a high growth temperature. Acknowledgment. The authors thank Anders Gustafsson for valuable discussions. This work was carried out within the Nanometer Structure Consortium in Lund and was supported by the Swedish Foundation for Strategic Research (SSF), the Swedish Research Council (VR), NoE SANDiE (EU Grant No. 500101), and NODE (EU Grant No. 015783). References (1) Qian, F.; Li, Y.; Gradecak, S.; Wang, D. L.; Barrelet, C. J.; Lieber, C. M. Nano Lett. 2004, 4, 1975-1979. (2) Choi, H.-J.; Johnson, J. C.; He, R.; Lee, S.-K.; Kim, F.; Pauzauskie, P.; Goldberger, J.; Saykally, R. J.; Yang, P. J. Phys. Chem. B 2003, 107, 8721-8725. (3) Sko¨ld, N.; Karlsson, L. S.; Larsson, M. W.; Pistol, M.-E.; Seifert, W.; Tra¨ga˚rdh, J.; Samuelson, L. Nano Lett. 2005, 5, 1943-1947. (4) Fang, Y.-P.; Xu, A.-W.; Dong, W.-F. Small 2005, 1, 967-971. (5) Xiang, J.; Lu, W.; Hu, Y.; Wu, Y.; Yan, H.; Lieber, C. M. Nature 2006, 441, 489-493. (6) Li, Y.; Xiang, J.; Qian, F.; Gradecak, S.; Wu, Y.; Yan, H.; Blom, D. A.; Lieber, C. M. Nano Lett. 2006, 6, 1468-1473. (7) Lauhon, L. J.; Gudiksen, M. S.; Wang, D.; Lieber, C. M. Nature 2002, 420, 57-61. (8) Biasiol, G.; Kapon, E. Phys. ReV. Lett. 1998, 81, 2962-2965. (9) Biasiol, G.; Gustafsson, A.; Leifer, K.; Kapon, E. Phys. ReV. B 2002, 65, 205306 (10) Biasiol, G.; Reinhardt, F.; Gustafsson, A.; Martinet, E.; Kapon, E. Appl. Phys. Lett. 1996, 69, 2710-2712. (11) Bovin, J.-O.; Alfredsson, V.; Karlsson, G.; Carlsson, A.; Blum, Z.; Terasaki, O. Ultramicroscopy 1996, 62, 277-281. (12) Hiruma, K.; Yazawa, M.; Katsuyama, T.; Ogawa, K.; Haraguchi, K.; Koguchi, M.; Kakibayashi, H. J. Appl. Phys. 1995, 77, 447462. (13) Wacaser, B. A.; Deppert, K.; Karlsson, L. S.; Samuelson, L.; Seifert, W. J. Cryst. Growth 2006, 287, 504-508. (14) Akiyama, T.; Sano, K.; Nakamura, K.; Ito, T. Jpn. J. Appl. Phys. 2006, 45, L275-L278. (15) No¨tzel, R.; Da¨weritz, L.; Ploog, K. Phys. ReV. B 1992, 46, 47364743. (16) Vurgaftman, I.; Meyer, J. R.; Ram-Mohan, L. R. J. Appl. Phys. 2001, 89, 5815-5875. (17) Tang, X.; Zhao, J.; Chin, M. K.; Mei, T.; Yin, Z.; Deny, S.; Du, A. Y. Appl. Phys. Lett. 2005, 87, 181906. (18) Borgstro¨m, M.; Deppert, K.; Samuelson, L.; Seifert, W. J. Cryst. Growth 2004, 260, 18-22.
Nano Lett., Vol. 6, No. 12, 2006
(19) Verheijen, M. A.; Immink, G.; de Smet, T.; Borgstro¨m, M. T.; Bakkers, E. P. A. M. J. Am. Chem. Soc. 2006, 128, 1353-1359. (20) Wulff, G. Z. Kristallogr. 1901, 34, 449 (21) Kley, A.; Ruggerone, P.; Scheffler, M. Phys. ReV. Lett. 1997, 79, 5278-5281.
Nano Lett., Vol. 6, No. 12, 2006
(22) Phillips, J. C. Bonds and Bands in Semiconductors; Academic Press: New York and London, 1973. (23) Mullins, W. W. J. Appl. Phys. 1957, 28, 333-339.
NL061692D
2747
