Controlling the Abruptness of Axial Heterojunctions in III–V Nanowires

Letter pubs.acs.org/NanoLett

Controlling the Abruptness of Axial Heterojunctions in III−V Nanowires: Beyond the Reservoir Effect Kimberly A. Dick,*,†,‡ Jessica Bolinsson,† B. Mattias Borg,† and Jonas Johansson† †

Solid State Physics, Lund University, Box 118, SE-221 00 Lund, Sweden Polymer & Materials Chemistry, Lund University, Box 124, SE-221 00 Lund, Sweden

‡

S Supporting Information *

ABSTRACT: Heterostructure nanowires have many potential applications due to the avoidance of interface defects by lateral strain relaxation. However, most heterostructure semiconductor nanowires suffer from persistent interface compositional grading, normally attributed to the dissolution of growth species in the common alloy seed particles. Although progress has been made for some material systems, most binary material combinations remain problematic due to the interaction of growth species in the alloy. In this work we investigate the formation of interfaces in InAs−GaAs heterostructures experimentally and theoretically and demonstrate a technique to attain substantially sharper interfaces. We show that by pulsing the Ga source during heterojunction formation, In is pushed out before GaAs growth initiates, greatly reducing In carry-over. This procedure will be directly applicable to any nanowire system with finite nonideal solubility of growth species in the alloy seed particle and greatly improve the applicability of these structures in future devices. KEYWORDS: Nanowire, heterostructure, III−V semiconductor, MOVPE, XEDS

S

class, partly due to the wide variety of material combinations that can be attained, but also due to important material properties such as narrow direct band gaps (important for optical devices) and high carrier mobilities (important for electronic devices). However, in III−V systems, solubility of the growth species in the alloy particle is still a major limitation: in particular, the group III species Al, Ga, and In typically have much higher solubilities in Au (and many other seed metals) than the group V species and form numerous intermediate compounds across the entire composition range for most temperatures of interest. It is thus anticipated that sharp interfaces should be much more difficult to attain in group III heterostructures than in group IV heterostructures.12 While atomically sharp group V heterointerfaces have been reported,4,15−18 reports of group III heterointerfaces typically indicate grading.19−23 Additionally, it is reported that III−V heterointerfaces are generally sharper in one direction: For Ga−In heterostructures it is observed that the Ga−In switch is much sharper than the In−Ga switch.19−21,24,25 This is qualitatively explained by the higher affinity of Au for In than for Ga, so Ga is replaced with In much more rapidly than vice versa.19,20,26 This poses a major challenge for potential heterostructure nanowire applications, many of which depend on the ability to form segments of one material between segments of another (for example as tunnel

emiconductor nanowires are widely considered one of the most promising technologies for future device applications. One of the greatest advantages of these structures is the ability to form axial heterostructures using lattice-mismatched materials,1,2 where the formation of interfacial defects can be avoided due to lateral strain relaxation.3 Axial superlattice structures composed of InAs−InP,4 GaAs−GaP,5 and Si− SiGe,6 all with lattice mismatch around 3%, were demonstrated nearly 10 years ago, and prototypical nanowire-based tunneling device structures have followed.7−9 Nevertheless, the use of metal seed particles to grow nanowires via vapor−liquid−solid (VLS) and related processes introduces complications to heterostructure growth not present in thin-film systems. One major complication is the difficulty to form sharp interfaces, attributed to the solubility of one or more growth species in the metal seed particle. This effect is particularly problematic in Si−Ge heterostructures grown using Au seeds, as both Si and Ge have substantial solubility in liquid Au alloys.6,10−12 This so-called reservoir effect has been investigated in detail experimentally and theoretically10,12 and has significantly limited the applications of Si−Ge heterostructure nanowires. However, major breakthroughs have recently been demonstrated: by modifying the composition and/or phase of the Au seed alloy (with the addition of another component), very sharp interfaces have been demonstrated in Si−Ge heterostructure nanowires.13,14 Although recent breakthroughs have greatly improved the prospects of group IV heterostructures, III−V material combinations still represent an important and interesting © 2012 American Chemical Society

Received: March 28, 2012 Revised: May 8, 2012 Published: May 29, 2012 3200

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7.7 × 10−4, respectively. GaAs nanowire growth temperatures of 380, 420, 435, 460, and 470 °C were investigated, with TMGa at a molar fraction of 2.4 × 10−5. Any deviations from this procedure and conditions are noted in the text.

barriers or quantum dots);7−9 for this reason sharp switching in both directions is desired. An even more serious problem is that most nanowire heterostructures tend to kink in one interface direction unless special growth steps are taken to prevent it.27 For the combination of GaAs with InAs (and similarly for GaP−InP) in nanowires, all reports indicate a very high tendency for kinking when switching from GaAs to InAs.19,25,27−29 It has been shown that the probability of kinking can be reduced below 50% if the diameter is increased toward 100 nm;25 however, this improvement comes at the cost of a high probability of interface dislocations. Thus, even for applications in which only one heterointerface is desired, the InAs−GaAs direction is preferred. A major breakthrough of the type recently achieved for Si− Ge heterostructures is clearly needed for III−V heterostructures in order to achieve interfaces which are both sharp and straight. However, the larger number of components in III−V systems makes the problem more challenging. Binary heterostructures already contain at least three (and sometimes four) growth species in combination with the alloy particle, and the chance that any one of them will have high solubility in the alloy is considerably greater. The tendency of group III species to form intermediate compounds with a great variety of otherwise suitable seed particle metals also means that solubility cannot generally be reduced simply by lowering the temperature, as it can for Si−Ge. To solve the problem of interface grading in binary heterostructures of this type, we need to develop new switching procedures to reduce the accumulation of growth species inside the alloy particle. The aim of this study is to demonstrate a technique for reducing the concentration of growth species in the alloy seed particle when forming nanowire heterointerfaces in order to substantially improve interface sharpness. We first investigate the role of excess growth species stored in the alloy particle and its influence on the interface sharpness in InAs−GaAs axial heterostructures seeded by Au particles, and then we develop a theoretical model which describes the results. Following this we use the understanding of how the growth species interact in the alloy particle to develop improved switching procedures to create sharper interfaces at nanowire heterojunctions. We select this particular material combination as a model system for understanding interface sharpness in group III heterostructures, since it is well-reported in the literature and interesting for future device applications. However, the results and conclusions should be directly transferrable to other material combinations. Heterostructure nanowires were grown by metal−organic vapor phase epitaxy at 10 kPa with H2 carrier gas flow of 13 L/ min. Au aerosol seed particles with various diameters from 20 to 90 nm were deposited onto InAs (111)B substrates with a density of 0.2 μm−2 for each diameter, yielding nanowire diameters ranging from approximately 15 to 85 nm (measured at the tip particle−wire interface) with a total density of 1 μm−2. Most often particles with all diameters were deposited onto the same sample piece, but samples with single diameters were also used for comparison. In all cases, samples were heated directly to growth temperature in AsH3, growth was initiated by introducing trimethylindium (TMIn), and heterointerfaces were formed by turning off TMIn, flushing up to 30 s in AsH3, and then turning on trimethylgallium (TMGa). Growth was stopped by turning off the appropriate group III precursor, and samples were cooled in either AsH3 or H2 only. InAs growth temperatures ranging from 380 to 460 °C were used, with TMIn and AsH3 molar fractions of 6.9 × 10−6 and

Figure 1. Interface gradient length for InAs−GaAs heterostructure nanowires as measured by X-ray energy dispersive spectroscopy, XEDS. (a) High angle annular dark field scanning transmission electron microscopy (HAADF-STEM) image of a heterostructure nanowire grown at 435 °C. The growth direction is from left to right. The black line indicates the position of the XEDS line scan. (b) XEDS line scan for the nanowire shown in (a), showing the decrease of In and the increase of Ga along the interface region of the nanowire. (c) Gradient lengths in individual heterostructure nanowires with different diameters and GaAs growth temperatures, extracted from XEDS line scans such as in (b); see Supporting Information File S1 for details. The solid line is a fit of our mass transport model for the gradient length to the experimental data.

We used X-ray energy dispersive spectroscopy (XEDS) in scanning transmission electron microscopy (STEM; JEOL 3000F operated at 300 kV, point resolution 0.12 nm, probe size 0.6 nm) to determine the variation in composition along InAs− GaAs nanowire heterointerfaces for different GaAs growth temperatures and diameters; full details of the procedure for extracting interface gradient lengths from in total 30 individual nanowires are discussed in Supporting Information File S1. As can be seen in Figure 1, there is a clear linear relationship between diameter and length of the interface gradient for all investigated temperatures. A dependence of gradient length on GaAs growth temperature is however not so clear (Figure 1c). Indeed, most of the data for the various temperatures fall along approximately the same line. 3201

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We next tested very short (10 s) growths of GaAs on InAs nanowires, with the intention of stopping the GaAs growth before the end of the gradient region. These experiments were conducted at a single growth temperature, 450 °C, and a single Au aerosol particle diameter of 40 nm. Different interruption procedures were tested between the InAs and GaAs growth, including AsH3 flush at growth temperature (for up to 10 min), flushing at elevated temperatures, and cooling to room temperature (in some cases even removing to air) before heating to GaAs growth temperature. All samples were cooled in hydrogen after GaAs growth. Additionally, different V/III ratios for InAs growth were tested (in the same range as for pure InAs; see above). For each set of conditions 10 nanowires were studied by TEM and XEDS. No effect of either switching procedure or InAs growth conditions was found for this series. For all investigated conditions, a GaxIn1−xAs segment of 10−15 nm length was obtained (see Figure 2). Remarkably, we also

The direct diameter dependence shown in Figure 1c gives important insight into the source of residual In in the graded region. We note that In is incorporated into a nanowire via the nanowire−particle interface, which scales with the square of the nanowire radius, R2. The alloy particle volume scales with R3, however, and so an incorporation of In into the interface from the particle would lead to an R dependence (R3/R2), as observed. By contrast, if the primary source of In was residual material on the side facets (which scale with R), an R−1 dependence would be expected, while if the source was residual material on the alloy particle surface (which scales with R2), no diameter dependence would be expected. After concluding that the primary source of background In in the graded interface is material remaining in the Au seed particle, we next consider how much In is stored and how it can be reduced. The postgrowth In content in the particle was measured following growth of pure InAs nanowires at temperatures of 380, 400, 420, 435, 440, 450, 460, 470, and 480 °C, using TMIn and AsH3 molar fractions in the ranges of (0.75−4.6) × 10−6 and (1.4−11) × 10−4, respectively, yielding V/III ratios from 20 to 1500. Two cooling procedures were tested: maintaining the AsH3 supply during cooling down to 300 °C, and shutting off the AsH3 together with the TMGa at growth temperature, to cool only in H2. Cooling in AsH3 is intended to give a composition close to the equilibrium content, while cooling in H2 gives a first estimate of the in situ composition. In the former case, the In composition of the seed particle was measured to be 30 ± 2 atomic %, for all investigated growth conditions, while in the latter case we measured In content of 38.5 ± 2 atomic %, again regardless of growth conditions. In total, 15 samples with different temperatures and V/III ratios were investigated for each cooling condition, and particle composition was measured for and averaged over 10 nanowires for each of these samples. The essentially constant In compositions measured for very different growth conditions is remarkable and suggests that the in situ condition is affected very little by accessible growth parameters. This is consistent with the observation (not shown) that the gradient length in InAs−GaAs nanowires is not significantly affected by the InAs growth temperature. However, the content of the seed particles after growth of InAs−GaAs nanowires (cooled in AsH3) is entirely different. We do not for any growth conditions observe measurable Ga in the seed particle even after 1.5 μm of GaAs growth, which is consistent with most other reports of GaAs, InGaAs, or GaP in a growth system exposed to In.19,20,24,25,30 However, we observe In content in the range of 15−20 atomic % (composition was measured for each of the 30 nanowires for which gradient lengths were also extracted). This did not vary significantly with GaAs nanowire length ranging from 300 to 1500 nm, and there was no obvious trend with diameter; however, we observe significant wire-to-wire variation (see Supporting Information File S2). The conclusion is that the In in the particle is clearly reduced during the GaAs growth step (calculations correlating this directly with the gradient length are shown in Supporting Information File S3). The addition of Ga reduces the In content from the very stable equilibrium values measured above. Similarly, In in the Au particle reduces the Ga steady-state composition.30 We therefore speculate that the activities of In and Ga in the gold particle are not independent but in fact directly connected. In other words, the ideal solution assumptions of the classic reservoir model12 need to be modified.

Figure 2. Early stage growth of a nominally GaAs heterosegment on an InAs nanowire. (a) HAADF-STEM image showing the formation of a short segment of GaxIn1−‑xAs on an InAs nanowire after 10 s growth time. A short InAs segment forms after the GaxIn1−xAs segment, even though the In source is not turned on. The scale bar is 25 nm. (b) Intensity profile from the nanowire averaged over area marked in (a); the horizontal axis is intensity counts on an arbitrary scale. Since the thickness is essentially constant, the intensity is primarily a function of mass, giving a qualitative picture of the composition change.

observed that the ternary GaxIn1−xAs segment was followed by a pure InAs segment of typically 5−8 nm, containing no measurable Ga. We speculate that this segment forms after the Ga and As sources are (simultaneously) switched off. The measured seed particle composition was 28 ± 4 atomic % In for all samples in the series, regardless of the steps taken before (nominal) GaAs growth. This composition is considerably lower than for pure InAs cooled in H2, confirming that the introduction of Ga immediately reduces the solubility of In in Au, causing In to leave the particle. This process continues even after the Ga source has been switched off and nominal GaAs growth ceases. It is clear that interface grading in InAs−GaAs heterostructures arises from the removal of excess material stored in the alloy particle during growth of the second segment, entirely in line with the classical reservoir model. However, it also clear 3202

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with J0A and the composition gradient is given by χA ≈ JA/J0A = exp(−t/τ). However, in order to compare with experiments, it is more practical to express χA as a function of length, l. We note that the axial growth rate for late stage growth, ν = l/t, can be approximated by ν = 2Fλ/R, where F is the deposition rate and λ is the effective diffusion length along the nanowire sidewall. This results in a composition gradient of

that there is significant interaction between the species within the alloy, such that the introduction of a third species entirely changes the equilibrium between the first two. This suggests that the Au−Ga−In alloy cannot be treated as an ideal solution and that the reservoir model must be extended to allow for nonideal activity coefficients for the two group III species. To address this, we develop a phenomenological model to describe the interface compositional gradients in binary heterostructures. When growing nanowire heterostructures consisting of a segment of material B on top of a segment of A, it is customary to switch off A and then wait until one believes that the concentration of A, xA, has reached the solubility limit, xeq A, before switching on B. We identify four different scenarios related to this process: (i) The presence of B does not influence the solubility of A in the metal particle. In this case the reservoir effect 12 holds, and it should be possible to achieve heterostructures without compositional grading, provided that eq xA has reached xeq A before the concentration of B reaches xB and B starts to grow. (ii) A small concentration of B, xB ≪ xeq B, decreases the solubility of A, xeq A , such that the particle may empty of A before xB has surpassed xeq B and B starts to grow. Sharp interfaces would be favored by a small decrease in the solubility of A upon adding B and fast emptying kinetics of A compared to the dissolution kinetics of B. (iii) B substantially decreases the solubility of A and a substantial amount of B, xB, incorporates before A reaches its new equilibrium concentration, xeq,B A . In this case B will start to grow as the particle empties of A, since xB will surpass xeq B before xA has decreased eq,B from xeq A to xA . Sharp interfaces will be difficult to obtain. (iv) If B increases the solubility of A in the metal particle, the interface will not be graded unless A is switched on again Case i is most likely relevant for Si−Ge nanowire heterostructures.12 Case iv we regard as technologically very promising if such a system exists. Cases ii and iii are related and may be relevant for group III heterostructures. As noted above, when InAs is grown on GaAs, the heterointerface is typically abrupt.15−18 We attribute this to case ii. On the other hand, when GaAs is grown on InAs, heterointerfaces reported in the literature are graded,19−23 consistent with the present study. We attribute this to case iii, and below we propose a mass transport model that describes the interface grading for this case. More details of the following derivation are given in Supporting Information File S4. Equating the flux of A out of the metal particle, JA, as given by Fick’s first law with the flux given by the time derivative of A in the metal particle, we can derive the time dependence of the concentration of A in the metal particle xA(t ) = xAeq + (xA0 − xAeq)e−t / τ

χA (l) = e−l / L

where L (characteristic gradient length) is given by

L = vτ =

2R2 3DA θA

4FλR 3DA θA

(4)

From eq 4 we see that the characteristic interface gradient length is proportional to the materials deposition rate, F, and the radius, R. There is no clear temperature trend since L is proportional to λ/D A, where both λ and DA change exponentially with temperature. We have fit XEDS line-scan measurements of selected heterointerface gradients to eq 3, with measured values for growth rate and diameter and DA as the only fit parameter (see Figure 3). Lengths and diameters of ensembles of InAs−GaAs

Figure 3. In composition along a heterostructure nanowire following initiation of GaAs growth at 470 °C. The Ga source is turned on at position zero. The black squares represent the In composition as measured by XEDS, and the solid curve is a fit of eq 3 to the experimental data (see also Supporting Information File S4).

nanowires were determined from scanning electron microscopy (SEM) images using a dedicated software;31 details are shown in Supporting Information File S5. Lengths of GaAs segments for selected nanowires were determined using high angle annular dark field scanning transmission electron microscopy (HAADF-STEM); since InAs growth conditions were the same for all samples in this sample set, it was possible to extract approximate growth rates for each material for all conditions. For each temperature about 30 nanowires were measured by SEM and 10 nanowires by TEM to determine growth rates. From fits such as the one shown in Figure 3, it is possible to extract values of DA; here we obtain a value of 2 × 10−17 m2 s−1. This value is reasonable for a solid phase diffusion coefficient, but roughly 8 orders of magnitude lower than typical liquidphase diffusion coefficients. This suggests that the alloy is in a solid phase during the growth of the graded section; that is, growth proceeds via the vapor−solid−solid (VSS) mechanism.32 Growth from a solid seed particle is not unreasonable

(1)

0 where xeq A is the solubility concentration of A, xA is the starting concentration, and τ is a characteristic time given by

τ=

(3)

(2)

In eq 2, R is the radius of the metal particle, DA is the diffusivity of A through the metal particle, and θA is Darken’s thermodynamic factor. Using eq 1, we will express the time dependence of the composition gradient in the nanowire, χA, which is given by χA = JA/J, where J is the time-dependent flux of material from the metal particle into the nanowire going from the steady-state rates J0A to J0B. If J is approximately constant, we can replace it 3203

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for these materials at these growth temperatures and has been proposed for GaAs, InAs, and InGaAs in other works at similar and higher temperatures.20,32 The diffusion coefficient is extracted for the gradient itself, and therefore we cannot directly infer the growth mechanism of the InAs and GaAs segments. However, since the growth rates which serve as input parameters to our model are extracted for the entire nanowire based on its total growth time, it may be logical that the growth mechanism is the same for the entire growth. As outlined above, the In stored in the particle is responsible for interface gradients in InAs−GaAs heterostructure nanowires, and this excess material is removed only once Ga is introduced into the system. With this in mind and from analysis of the implications of the scenarios mentioned above, we have designed and evaluated a novel procedure for obtaining sharper axial heterojunctions in nanowires. If scenario iii above applies, there should be a finite time between switching on the Ga source and the initiation of GaAs growth. However, in the situation of a more or less immediate reduction of xeq In upon the introduction of Ga to the system, then we would expect In to be “pushed out” of the particle as soon as Ga is introduced, even before the level of Ga has reached a critical supersaturation. When GaAs growth is terminated after a very short period of time (as in the situation displayed in Figure 2 above), it is indeed interesting to note that after a GaInAs graded region a short InAs segment is formed. This observation strongly indicates that even low concentrations of Ga in the seed particle influence the equilibrium concentration of In in the seed particle. In order to reach sharper interfaces in the InAs−GaAs nanowire system, we therefore tested a technique whereby small amounts of Ga were introduced into the seed particle using short pulses of the TMGa precursor source (1 s “on” followed by 30 s “off”). Similar source pulsing approaches or supply interruption methods (SIM)33 have also been used to improve and engineer nanowire crystal structure.30,33,34 Our aim in using this approach was to evaluate whether the In content in the particle could be reduced by modifying its equilibrium concentration without supplying sufficient Ga for xeq Ga to be reached (that is, without initiating GaAs growth). Once the In content within the seed particle has been sufficiently reduced, the Ga supply can be turned on continuously to initiate growth of the GaAs segment. Growth temperatures of 380 and 460 °C were tested, and samples were cooled in AsH3. For both temperatures we observe substantially shorter gradient lengths compared to nanowires grown using the same parameters but without pulsing (Figure 4a). The gradient lengths can be tuned by optimization of the procedure, in this case by increasing the number of pulses (see Figure 4a); further optimization should be feasible by for example tuning TMGa and AsH3 flow during the pulsing. Figure 4b and Figure 4c show sample XEDS linescan and HAADF-STEM images, respectively, for a nanowire with 30 nm diameter grown using 20 pulses. Although the gradient length determined from the XEDS line scan is about 7 nm, the HAADF-STEM image shows apparently atomically sharp contrast change. Further, it is clear from intercepts of the linear trendlines in Figure 4 that the minimum resolvable gradient length for both the nonpulsed and the pulsed heterojunction is about 5 nm. The spatial resolution of the XEDS is limited by beam broadening due to scattering in the sample, and can be estimated for this material, average thickness and accelerating voltage to be around 5 nm.35 It

Figure 4. Reduction of interface gradient length in InAs−GaAs heterostructure nanowires using a pulsed interface switching technique. To achieve sharper gradients by reducing the steady-state In content of the Au alloy particles, the Ga source was pulsed before initiating GaAs growth. (a) Gradient lengths extracted from XEDS line scans (see Supporting Information File S1) for different diameters. Two different GaAs growth temperatures were tested, and as shown in the graph, substantially shorter interface gradients were the result for both temperatures. There is also a significant improvement with increased number of Ga source pulses, indicating the process can potentially be optimized to yield even sharper gradients. Calculated linear trendlines for both continuous and pulsed growth nanowires (shown for 460 °C) have intercepts of about 5 nm, confirming that this is the minimum resolvable distance by XEDS analysis under these conditions. (b) XEDS line scan for a nanowire with 30 nm tip diameter grown at 460 °C using a 20-pulse procedure. A gradient length of 7−8 nm can be extracted from this plot. (c) High-resolution HAADF-STEM image in the ⟨110⟩ zone axis of the interface of the same nanowire as in (b), showing contrast change on the order of a few atomic planes. The scale bar is 5 nm.

may be reasonable to assume then that all nanowire gradients are in fact sharper than the XEDS estimate by about this amount. This limits the applicability of the technique for quantification of the very sharpest heterojunctions attained by the supply interruption method. To further investigate the interface sharpness of nanowires with pulsed heterojunction switching, we performed highresolution microscopic analysis in both TEM and STEM mode. Figure 5a shows a high-resolution TEM image of an InAs− 3204

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GaAs nanowire grown using the pulsed switching (5 pulses). The uniform strain field surrounding the interface is a first indication that the lattice parameters change over a relatively short distance, without the creation of dislocations. Compositional mapping of the InAs and GaAs regions was performed by Fourier analysis of the image in Figure 5a, by calculating the associated fast Fourier transform (FFT) for the entire image, masking specific reflections associated with InAs and GaAs lattice parameters, and then calculating inverse FFTs from the masked regions. Overlaying these inverse FFTs thus gave a map of the regions associated with specific lattice parameters in the image. Further details are shown in Supporting Information File S6. The resulting Fourier map, shown in Figure 5b, confirms that pure InAs and GaAs are present very close to the interface. We have further examined this structure by high-resolution STEM, as shown in Figure 5c; the two segments of different material are clearly visible from their atomic contrast. It is evident that the structure changes abruptly, although there is some atomic mixing near the interface which is not observed for the sharper interface in Figure 4c. From these images and additional analysis of the lattice spacings in the associated fast Fourier transforms (see Supporting Information File S6), an interface gradient of about 10 nm can be estimated for this 38 nm diameter nanowire. This is significantly sharper than previously reported InAs−GaAs heterointerfaces as well as those grown in this study without pulsed switching (∼35 nm gradient for this diameter). XEDS line-scan analysis of this particular nanowire yielded a gradient length of 15 nm (see Figure 5d), further confirming that the gradients as determined by XEDS are limited to a minimum resolvable distance of about 5 nm and giving weight to the conclusion that the sharpest interfaces in this study (for a smaller diameter) are approaching atomic-level sharpness. In summary, we have shown substantial improvement in InAs−GaAs heterostructure nanowire interface sharpness by introducing a novel technique to remove excess growth species from the alloy during heterojunction formation. Interface gradients in InAs−GaAs heterostructure nanowires are attributed to the significant solubility of the group III species in the Au alloy seed particle, very similar to what has been demonstrated for Si−Ge heterostructure nanowires. However, the components in this study (In, Ga, and Au) do not form an ideal solution, but instead the equilibrium concentrations of In and Ga are related to each other. Consequently, excess In is released from the seed particle when the Ga source is introduced due to a reduction in its equilibrium concentration, leading to an interface composition gradient. We demonstrate that pulsing the Ga source before switching to GaAs introduces Ga levels in the alloy which are sufficient to lower the In solubility (and thus drive out In) before GaAs growth is initiated. We show that interface sharpness can be improved by a factor of 3−4 to levels approaching a few atomic layers for optimized conditions. This technique should be applicable to other heterostructure combinations in nanowires, and thereby similar improvements of the interface sharpness should be achievable for other material systems.

Figure 5. High-resolution evaluation of the InAs−GaAs heterointerface for nanowires grown using pulsed interface switching. (a) Highresolution TEM image in the ⟨110⟩ zone axis of the interface region of an InAs−GaAs nanowire with 38 nm tip diameter, grown with a fivepulse Ga source switching procedure at 380 °C. The scale bar is 20 nm. The uniform strain field surrounding the interface indicates that the change in lattice parameter occurs over a relatively short distance. (b) Compositional map obtained by Fourier mapping (see Supporting Information File S6) using the 111 spots in the fast Fourier transform (FFT), with red representing the lattice parameter of InAs and green the lattice parameter of GaAs. The color change is complete over a distance of less than 10 nm, again indicating a quite sharp interface. A uniform strain field is visible here as well, indicating variation in the lattice constant due to interface strain. Further analysis of the InAs− GaAs heterojunction, including more Fourier maps and FFTs, are available in Supporting Information File S6. (c) High-resolution HAADF-STEM image in the ⟨110⟩ zone axis of the interface of the same nanowire, with 5 nm scale bar. It is clear that the contrast changes from brighter (InAs) to darker (GaAs) spots over a distance of not more than 10 nm, although there is evident atomic mixing in the interface region. (d) XEDS line scan from the same nanowire. Analysis of this line scan indicates a gradient length of about 15 nm, indicating that the XEDS resolution limits the determination of the interface gradient in this range. XEDS point analysis confirms detection of Ga at about the same position as the contrast changes in (c).

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ASSOCIATED CONTENT

S Supporting Information *

Interface gradient extraction methods, composition of Au alloy particles, correlation of gradient length with seed particle composition, details of mass transport model, extraction of 3205

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nanowire lengths, high-resolution imaging and compositional mapping of interfaces. 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]. Author Contributions

The manuscript was written through contributions of all authors. K.A.D. performed TEM and XEDS analysis and directed the project, J.B. and B.M.B. performed growth experiments and SEM analysis, and J.J. performed modeling analysis. All authors discussed the results and implications and contributed to writing the paper and gave approval to the final version. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the Swedish Research Council (VR), the Swedish Foundation for Strategic Research (SSF), the Knut and Alice Wallenberg Foundation (KAW), and the Nanometer Structure Consortium at Lund University (nmC@ LU).

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REFERENCES

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dx.doi.org/10.1021/nl301185x | Nano Lett. 2012, 12, 3200−3206