Letter pubs.acs.org/NanoLett
Enhanced Luminescence Properties of InAs−InAsP Core−Shell Nanowires Julian Treu,† Michael Bormann,† Hannes Schmeiduch,† Markus Döblinger,‡ Stefanie Morkötter,† Sonja Matich,† Peter Wiecha,† Kai Saller,† Benedikt Mayer,† Max Bichler,† Markus-Christian Amann,† Jonathan J. Finley,† Gerhard Abstreiter,†,§ and Gregor Koblmüller*,† †
Walter Schottky Institut, Physik Department, and Center of Nanotechnology and Nanomaterials, Technische Universität München, Am Coulombwall 4, Garching, 85748, Germany ‡ Department of Chemistry, Ludwig-Maximilians-Universität München, Munich 81377, Germany § Institute for Advanced Study, Technische Universität München, Garching 85748, Germany S Supporting Information *
ABSTRACT: Utilizing narrow band gap nanowire (NW) materials to extend nanophotonic applications to the mid-infrared spectral region (>2−3 μm) is highly attractive, however, progress has been seriously hampered due to their poor radiative efficiencies arising from nonradiative surface and Auger recombination. Here, we demonstrate up to ∼102 times enhancements of the emission intensities from InAs NWs by growing an InAsP shell to produce core−shell NWs. By systematically varying the thickness and phosphorus (P)-content of the InAsP shell, we demonstrate the ability to further tune the emission energy via large strain-induced peak shifts that already exceed >100 meV at comparatively low fractional P-contents. Increasing the P-content is found to give rise to additional line width broadening due to asymmetric shell growth generated by a unique transition from {110}- to {112}-sidewall growth as confirmed by cross-sectional scanning transmission electron microscopy. The results also elucidate the detrimental effects of plastic strain relaxation on the emission characteristics, particularly in core−shell structures with very high P-content and shell thickness. Overall, our findings highlight that enhanced mid-infrared emission efficiencies with effective carrier confinement and suppression of nonradiative recombination are highly sensitive to the quality of the InAs−InAsP core−shell interface. KEYWORDS: InAs−InAsP core−shell nanowires, PL spectroscopy, molecular beam epitaxy, metalorganic vapor phase epitaxy, transmission electron microscopy
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probability for charge carrier scattering at the NW surface reduced14 but the usually high nonradiative surface recombination rates at surface states can be suppressed in NWs fabricated from commonly used compound semiconductors.13,15,16 This was recently shown to significantly enhance the photoluminescence (PL) efficiency of, for example, GaAs NWs passivated by a conformal AlGaAs,13,16−18 GaInP,19 or GaAsP5 shell. Consequently, this has also fueled important
adial core−shell nanowires (NW) have been extensively studied over recent years1−3 due to their great potential in advancing fundamental research of one-dimensional (1D) nanostructures while improving several semiconductor-based device technologies, including NW-based lasers,4,5 lightemitting diodes,6,7 photodetectors and solar cells,8−10 and high-performance, highly downscaled transistors.11,12 The major advantages provided by the radial core−shell geometry are an effective carrier confinement along with unique waveguiding capabilities, and a viable surface passivation scheme that is made particularly important by the large surface-to-volume ratios in NWs.1,3,13 Thus, not only is the © 2013 American Chemical Society
Received: September 6, 2013 Revised: November 15, 2013 Published: November 25, 2013 6070
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molecular beam epitaxy (MBE)/metal−organic vapor phase epitaxy (MOVPE) growth procedure. Using this procedure, the InAs NW cores were grown by MBE, whereas the shell overgrowth was performed via MOVPE with all growth details listed in the Supporting Information. Two respective series of samples were grown: Series (A) under conditions yielding fixed shell composition (i.e., phosphorus content x(P) ∼ 0.09) but variable shell thickness, and series (B) with a fixed nominal thickness, while the P-content was varied across the entire compositional range, even up to x(P) = 1.0 (binary InP). Figure 1 shows two exemplary scanning electron microscopy (SEM) images of sample series (A), that is, (a) the “core-only”
advances in NW−based optical devices, such as efficient GaAs− AlGaAs and GaAs−GaAsP core−shell NW light-emitting diodes and lasers5,7,20 emitting in the near-infrared (IR) spectral range. Utilizing this concept for NW-based light sources with emission wavelengths (λ) extending further into the mid-IR spectral region is of the utmost importance for future IRphotonic applications. This wavelength range can be reached using lower−energy gap InGaAs and InAs NW materials with λ ∼2−3 μm. However, progress has been hampered by very poor optical efficiency of these materials in a NW geometry due to very strong nonradiative surface and Auger recombination combined with the well-known difficulties of performing spectroscopy in the IR spectral region (poor detector sensitivity and strong atmospheric absorption).21 Consequently, very few optical emission studies have been reported to date on InAs and high In-content InGaAs NWs, and those that have appeared show relatively poor emission efficiencies in the ∼2−3.5 μm spectral range. This is especially true for lattice temperatures above ∼150 K due to nonradiative surface and Auger recombination.22−25 As a result, passivation by a suitable conformal shell structure is much needed to yield improvements in PL efficiency of In(Ga)As-based NWs. In this respect, InP and InAsP shells can be considered to be the most suitable surface passivating structures since they can be grown in situ and epitaxially around the In(Ga)As NW core generating a desired type-I band alignment.11 Furthermore, any modification of the composition (i.e., As/P ratio) provides the flexibility to tune the strain, band-offsets, and successively the carrier confinement accordingly.26−28 However, as the In(Ga)As−In(As)P core−shell NW system is prone to lattice mismatches of up to ∼3.1%, it requires very careful selection of shell thickness and composition in order to maintain pseudomorphic and defect−free structures. Moreover, the long carrier lifetimes and comparatively low surface recombination velocities of In(As)P make such core−shell NW combinations also promising for applications in the field of photovoltaics.29 Given that InAs-based NWs are simultaneously very promising candidates for novel high-speed electronic devices, passivation schemes by InP and InAsP shells have already been employed to explore the effects on improved electrical transport characteristics.11,30−32 In contrast, the influence of in situ grown InAsP shells on the optical emission characteristics of InAs NWs has received little attention despite the relevance for future NW-based photonics in the mid-IR spectral region. In this Letter, we report the strong enhancement of the luminescence efficiency of InAs NWs grown using a catalystfree procedure by introducing an InAsP shell that acts as a surface passivation layer. Depending on the thickness and phosphorus (P) content of the InAsP shell, the PL efficiency of the InAs core is found to be up to ∼102 times higher than unpassivated InAs NWs facilitating optical studies up to room temperature (300 K). The thickness and P-content of the shell was found to have a strong influence on the radiative properties and strain-related PL peak energy variations, resulting in blue shifts as large as >100 meV for P-contents exceeding x(P) ∼ 0.2. These results demonstrate the effective suppression of surface states and enhanced carrier confinement in InAs NWs utilizing high-quality heterointerfaces with the surrounding InAsP shell. All InAs−InAsP core−shell NWs investigated in this study were grown in a position-controlled and completely catalystfree growth mode on prepatterned Si(111) using a hybrid
Figure 1. SEM images of (a) core-only InAs NW reference and (b) InAs−InAsP core−shell NWs from sample series A (with shell thickness of ∼8 nm, overgrowth time of 19.1 min). Note the different length scales laterally and vertically arise from the absence of SEM tilt correction for better visibility. (c) Schematic illustration of the actual InAs−InAsP core−shell NW structure that includes also an axial InAsP top segment.
InAs NWs as reference, and (b) InAs−InAsP core−shell NWs with the thickest grown InAsP shell. Note that this series contained a total of four samples with different shell thicknesses (overgrowth times) with all other respective SEM images shown in the Supporting Information S1. We find that all samples reveal a very high growth yield of >90% (i.e., number of NWs/defined nucleation site) demonstrating the stable growth conditions on the prepatterned Si(111) substrate.33 The reference sample in Figure 1a shows a very homogeneous size distribution and perfect vertical directionality along the [111] direction of the core-only InAs NWs with a hexagonal shape terminated by {110} sidewall facets.34 Note that the nomenclature for the crystallographic indices is chosen in zinc blende (ZB) structure notation, adapted from the underlying [111] orientation of the Si substrate. As expected, with increasing InAsP growth time the NWs become wider yielding radial shell growth along the NW sidewalls (see Figure 1b and Supporting Information Figure S1). However, the NWs also tend to increase in length, indicating that the InAsP overgrowth also produces an axially grown region along the [111] direction on top of the NW core (as illustrated in Figure 1c). Concurrently, increased overgrowth times also yield a larger variation in length and diameter distribution together with a slight tilt (1 μm), such that not all photocarriers excited within the segment diffuse into the NW core and recombine there. We have confirmed this also by measurements performed on core− shell NWs with much shorter (longer) InAsP segment lengths, which resulted in a strong reduction (increase) in the segmentrelated emission intensity. The data shows a clear ∼20-fold increase in the PL peak intensity of the InAs core luminescence upon capping with InAsP. This demonstrates directly the very effective reduction of nonradiative surface recombination by the passivation of surface states, where the InAsP shell acts as an energy barrier confining carriers to the NW core and preventing their diffusion to the NW surface. Interestingly, this effect occurs already at relatively low barrier heights (calculated band offsets of ∼45 meV for the respective P content of x(P) = 0.09) and a very small shell thickness (∼3 nm). Furthermore, the strong
given transmission conditions. The HR-TEM image further verifies that no plastic relaxation has occurred between the NW core and shell region based on the absence of Moiré fringes and the fact that no splitting in the electron diffraction pattern is visible. This is also in full accord with our expectations considering the very low, that is, ∼ 0.3% lattice mismatch between the InAs core and InAsP shell layer with a P-content of x(P) = 0.09. Detailed insights into the effect of the InAsP shell on the optical properties of the NWs were obtained from photoluminescence (PL) spectroscopy measurements. Here, we used a home-built PL setup specifically designed for the infrared (IR) region, using a He/Ne laser (633 nm) for excitation and a Horiba Jobin-Yvon HR640 monochromator equipped with a liquid nitrogen cooled InSb photodetector. Further experimental details can be found elsewhere.25 The excitation laser spot size focused onto the sample had a diameter of ∼5.5 μm corresponding to the excitation of an ensemble of ∼350 NWs. It is important to note that due to the site selective growth method used the NW density and, hence, the excitation volume was similar for all investigated samples, allowing direct comparison of the absolute emission intensities between different samples. Figure 3a compares the measured PL spectra of the core-only InAs NWs (reference sample) and an InAs− InAs0.91P0.09 core−shell NW sample (sample with intermediate overgrowth time and average shell thickness of ∼3 nm), as recorded at 8 K and an excitation power density of 140 W/ 6073
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Figure 4. (a) Representative SEM image of the InAs−InAsP core−shell NW sample with P content of x(P) = 0.37 (nominally x(P) = 0.6). The morphology shows a bimodal distribution of thin and thick NWs, consistent with core−shell structures with either a single embedded NW core or, multiple fully embedded NW cores as shown in the respective STEM-HAADF micrographs of (b,c), respectively. Note the unique change from {110}- to {112}-oriented sidewall facets from InAs core to InAsP shell.
spatially separated transitions due to surface Fermi level pinning do not contribute much to the PL signal of the core luminescence, while the line width is rather limited by crystalline defects.25 As summarized in Figure 3(d), not only does the peak energy increase for larger shell thicknesses but also the PL peak intensity. Note, however, that the maximum PL peak intensity is reached at a shell thickness of ∼3 nm, while it decreases again beyond. We ascribe this primarily to the increased laser absorption by the InAsP top segment (inhibited excitation of the NW core) with increasing segment length. Further insights into the effect of lattice-mismatch strain on the NW core luminescence are provided from compositiontuned InAsP shell overgrowths (series B). For this series, we successively changed the nominal P content from InAsP with low x(P) up to pure InP (x(P) = 1), corresponding to much higher lattice mismatches (up to 3.1%) as compared to series A. The entire set of growth parameters, together with the resulting NW dimensions as obtained from SEM measurements are listed in Table 2 in the Supporting Information. While both the core-only InAs NW reference and the core−shell NW sample with the lowest nominal P-content (x(P) = 0.2) showed very uniform high-periodicity NW morphology (see Figure S2 in Supporting Information), NWs overgrown with higher Pcontent shells resulted in rather inhomogeneous morphology as seen in Figure 4a and Supporting Information Figure S2. Despite the underlying high homogeneity of the NW core growth, the overgrowth produced a distribution of core−shell NWs with either a single core or, surprisingly, also with multiple completely embedded cores (see Figure 4b,c). This complicates a reliable determination of shell thickness and other geometrical parameters. Consequently, two data sets are presented in Table 2 with separate values for single core−shell and multiple core−shell structures (Supporting Information). The reason for the overgrowth around multiple cores could be exactly the same stress-induced mechanism that triggered tilt and convergence of NWs for extended overgrowths as noted in the discussion of Figure 1. Indeed, SEM images of the nearsubstrate surface region (not shown) reveal that the bottom parts of individual NWs are still straight and well-aligned while obviously the NWs have bent slightly in the upper regions. The STEM-HAADF images of Figure 4b,c (taken from the sample with nominal P content of x(P) = 0.6) illustrate that the growth of the InAsP layer takes place with a 30° rotation with respect to the {110} sidewall facets of the NW core. This observation indicates that the extended side facets of the shell structure can be assigned to the {112} family of facets, which
increase in the intensity of the core luminescence allows us to obtain even emission at 300 K (see Figure 3b), which so far has been difficult to observe in InAs-based NWs.22−25 Note that the temperature dependence of the core luminescence does not adhere to the commonly observed continuous redshift with temperature (Varshni behavior). This is most likely due to localized trap states near the band edge (mainly crystalline defects), as recently also observed in WZ-type In-rich InGaAs NWs with stacking defects.25 Another interesting feature is the characteristic blueshift in the InAs emission (∼22 meV) when the NW core is capped with an InAsP shell. To account for this blueshift we considered several possibilities. First, the blueshift could be due to differences in the band-filling characteristics of the core−shell NW sample with respect to the unpassivated InAs NW reference. Excitation power-dependent measurements (not shown), however, rule these out since both samples show similar band filling, that is, identical blueshifts and peak broadening over the entire excitation power range. Second, for unpassivated InAs NWs the surface Fermi level pinning and associated band bending at the surface may result in spatially separated electrons near the surface and holes confined rather in the center of the NW. In analogy, capping by In(As)P is expected to decrease the band bending near the heterointerface,30 such that the electron−hole pair transitions may occur at slightly higher energy due to their reduced spatial separation.39 Considering, however, that the NW core diameter is relatively wide and the band offsets with the low P-content InAsP shell small, we expect the contribution from spatially separated photocarriers to be rather small. We therefore conclude that strain is the most likely origin for the observed blueshift. Recent measurements of the strain state in InAs−InAsP core−shell NWs have shown that the NW core is compressively strained in all dimensions,29 which induces a widening of the InAs band gap and hence explains a shift of the PL peak emission to higher energies. Such strain-induced peak shifts of the NW core emission have also been observed in other core−shell NW systems, such as, for example, GaAs− GaInP and GaAs−AlGaAs core−shell NWs.40,41 Figure 3c shows normalized PL spectra including all other core−shell NW samples (from series A), illustrating that with increasing shell thickness and hence increased strain the peak emission shifts to higher energies. Note that the PL peak shape and widths are very similar for all samples, indicating that passivation by the InAsP shell does not reduce the line width. This is also an indication that the previously considered 6074
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Figure 5. (a) PL spectrum of InAs−InAs0.63P0.37 core−shell NWs in comparison with that of core-only InAs NWs (reference) recorded at 8 K and excitation power density of 140 W/mm2. A ∼110-fold increase in PL efficiency together with a strong blueshift of the core luminescence is observed. (b) Normalized PL spectra of all InAs−InAsP core−shell NW samples from series B (variable P content series) measured under the same conditions. The nominal P content as well as the actual P content (in brackets) is given. (c) Plot of PL peak energy and peak intensity of the core luminescence as a function of actual P content. (d) TEM micrographs recorded along the [110] zone axis for a NW with a P content in the shell of x(P) = 0.37 (upper image), showing no indication of relaxation at the core−shell interface, and a NW with a pure InP shell (lower image) where distinct Moiré fringes evidence at least a partial relaxation of the shell.
the preferential {112} facets for higher P content. To fully understand the intricate dependencies of {110}- versus {112}faceted growth, entire parameter space maps as a function of all relevant growth parameters and P-content need to be generated. To determine the composition within these overgrown structures, EDXS spectral maps were also recorded from the same cross-sectional foils of Figure 4b, yielding a P content of x(P) = 0.37 ± 0.06 that is almost a factor of 2 lower than the nominal value. Moreover, it is important to note that the heterointerfaces between the InAs cores and the InAsP shell layers do not evidence signatures of plastic relaxation, confirming that these structures are still pseudomorphically strained. Figure 5a compares PL spectra from the same InAs−InAsP core−shell NW sample as of Figure 4 and the core−only InAs NW reference sample (recorded again at 8 K and an excitation power density of 140 W/mm2). As compared to Figure 3, the larger P content of x = 0.37 resulted here in an even stronger PL efficiency (110-fold enhancement) and much larger blueshift of the core luminescence. Note also that the InAsPassociated peak at 0.73 eV corresponds closely with the Pcontent derived from EDXS measurements. Normalized PL spectra of the core luminescence of all InAs−InAsP core−shell NW samples from series B (variable P content series) are further shown in Figure 5b. Comparison of the PL linewidths of the samples with higher P content (x(P) = 0.37 (nominally 0.6) and x(P) = 1) give peaks that are approximately twice as broad.
was confirmed by the orientation of NWs relative to the substrate cleavage planes as well as by TEM. A closer look, however, reveals that the shell facets are not entirely composed of {112} facets but also of very narrow {110} facets (widths 100 meV already at P contents as low as ∼0.3. This clearly demonstrates the very effective suppression of surface states and carrier confinement provided by pseudomorphically strained, high-quality InAs− InAsP heterointerfaces as confirmed by TEM. Increased Pcontent, however, yielded larger PL line width broadening associated with the formation of asymmetric shells due to a transition from {110}- to {112}-faceted sidewall growth as shown by cross-sectional STEM-HAADF imaging. Finally, very high P-content InAs−InAsP core−shell structures showed a drastic reduction in both PL efficiency and blueshift due to the onset of plastic relaxation, highlighting the pronounced sensitivity of the emission characteristics to the quality of the core−shell NW heterointerface.
This can be related to the inhomogeneous shell thicknesses arising from the {112}-faceted growth and the corresponding fluctuations in strain and composition, as averaged by the ensemble measurements. Nevertheless, as expected for increasing P content and increasing compressive strain in the InAs core, the PL peak shifts gradually to higher energies with shifts that are as high as 115 meV (for the sample with x(P) = 0.37, see also Figure 5c). In strong contrast, for the sample with binary InP overgrowth the blueshift is significantly reduced, most likely due to the onset of plastic strain relaxation and associated dislocation formation, but also likely due to a change in the shape of the wave function and thus the interband matrix element that ultimately influences the effective transition energy. To verify the effects of strain on the shape of the wave function and associated transition energies (i.e., PL peak shift), we performed a two-dimensional (2D) simulation of the hexagonal core−shell NW and calculated the bandstructure (together with the 10 lowest wave functions and eigenvalues in the conduction and valence band, see Supporting Information for details). Good qualitative agreement between the calculated effective transition energies and the as-measured PL peak energies is observed. Furthermore, to account for the anticipated strain relaxation occurring in the core−shell NWs with the highest P content (binary InP overgrowth), we have performed additional TEM measurements along the [110] zone axis of individual NWs as shown in Figure 5d. In the upper image, a NW with a P content in the shell of x(P) = 0.37 is shown, revealing no signatures of strain relaxation related defects at the core−shell interface. In contrast, the NW with a pure InP shell (lower image) shows distinct Moiré fringes, evidencing a partial relaxation of the shell. This is in good agreement with previous microstructural data by Keplinger et al.,29 who mapped for the same core−shell material system the transitions between pseudomorphic and plastically relaxed growth as a function of P-content and shell thickness. To facilitate a direct comparison of the relevant PL data, Figure 5c summarizes the measured PL peak positions and peak intensities of the core luminescence as a function of the P content x(P). Obviously, increasing P content not only results in larger blueshifts but also a strong increase in PL efficiency, provided that the structures are not yet plastically relaxed. Thus, a maximum of ∼102-fold increase in PL intensity was found for the core−shell NW sample with x(P) = 0.37, proving the more effective carrier confinement and hence superior passivation as compared to InAsP overgrowth with lower P content. As expected, for the binary InAs−InP core−shell NW sample the PL efficiency is substantially reduced by approximately 1 order of magnitude. This can be ascribed again to the degraded microstructure due to plastic relaxation but partly also to higher laser absorption in the extended InP top segment with less effective excitation of the core. This means that symmetric shell growth across all sidewall facets with smaller shell thicknesses will be necessary to allow for defect-free, pseudomorphic InP overgrowth with unperturbed emission properties. In summary, we demonstrated the importance of in situ passivation in radial InAs−InAsP core−shell NWs on their optical emission efficiency and associated recombination properties in photoluminescence (PL) experiments. Most strikingly, the radiative emission intensities were ∼101−102 times higher compared to unpassivated InAs NWs with emission well up to 300 K, as identified by systematic variation
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ASSOCIATED CONTENT
S Supporting Information *
Experimental details on growth procedures and morphological investigations, as well as calculations of the strain-induced transition energies. 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]. Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS The authors thank S. Birner for valuable discussions regarding nextnano simulations, A. W. Holleitner and P. Weiser for SEM support, as well as G. Scarpa and P. Lugli for support with NIL. This work was funded by the Marie Curie FP7 Reintegration Grant, the EU FP7 project SOLID, the DFG excellence program Nanosystems Initiative Munich, and the collaborative research center SFB 631. Further support was provided by the Technische Universität München, Institute for Advanced Study, funded by the German Excellence Initiative.
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REFERENCES
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dx.doi.org/10.1021/nl403341x | Nano Lett. 2013, 13, 6070−6077
