Home
Search
Collections
Journals
About
Contact us
My IOPscience
Axial InAs/GaAs heterostructures on silicon in a nanowire geometry
This content has been downloaded from IOPscience. Please scroll down to see the full text. 2014 Nanotechnology 25 485602 (http://iopscience.iop.org/0957-4484/25/48/485602) View the table of contents for this issue, or go to the journal homepage for more
Download details: IP Address: 129.97.58.73 This content was downloaded on 16/12/2014 at 16:27
Please note that terms and conditions apply.
Nanotechnology Nanotechnology 25 (2014) 485602 (6pp)
doi:10.1088/0957-4484/25/48/485602
Axial InAs/GaAs heterostructures on silicon in a nanowire geometry C Somaschini1, A Biermanns2, S Bietti3, G Bussone2,4, A Trampert1, S Sanguinetti3, H Riechert1, U Pietsch2 and L Geelhaar1 1
Paul-Drude-Institut für Festkörperelektronik, Hausvogteiplatz 5–7, 10117 Berlin, Germany Universität Siegen, Festkörperphysik, 57068 Siegen, Germany 3 LNESS and Dipartimento di Scienza dei Materiali dellʼUniversità di Milano-Bicocca, via Cozzi 53, 20125 Milano, Italy 4 ESRF, 6 rue Jules Horowitz, BP220, F-38043 Grenoble Cedex, France 2
E-mail: [email protected] Received 7 March 2014, revised 8 October 2014 Accepted for publication 13 October 2014 Published 13 November 2014 Abstract
InAs segments were grown on top of GaAs islands, initially created by droplet epitaxy on silicon substrate. We systematically explored the growth-parameter space for the deposition of InAs, identifying the conditions for the selective growth on GaAs and for purely axial growth. The axial InAs segments were formed with their sidewalls rotated by 30° compared to the GaAs base islands underneath. Synchrotron X-ray diffraction experiments revealed that the InAs segments are grown relaxed on top of GaAs, with a predominantly zincblende crystal structure and stacking faults. Keywords: molecular beam epitaxy, nanowire, heterostructures, X-ray diffraction, droplet epitaxy
1. Introduction
fabricate axial heterostructures with a switch on the group-III sublattice by molecular beam epitaxy (MBE). Issues have been reported in the case of Au-catalyzed growth, where the complex energetics of the ternary Au-Ga-In system can give rise to peculiar phenomena, such as the crawling of the Au droplet on a side facet, which prevents the axial growth [9] and in the case of the self-assisted mode, where the presence of a gallium droplet on top of the GaAs NWs strongly limits the compositional variation along the growth axis [10]. Finally, unintentional radial growth frequently dominates over axial growth [11, 12]. Here, we propose droplet epitaxy (DE) [13, 14] as a possible way to achieve the epitaxial growth of axial InAs/ GaAs heterostructures on silicon substrates in a NW geometry, without the need of any pre-patterning. By DE, GaAs islands of the desired size and number density can be formed on silicon [15, 16] and used as a template for the overgrowth of InAs. We study the influence of the growth-parameters on the deposition of InAs segments on top of the GaAs islands. In particular, we show that high substrate temperature favors the selective deposition of InAs on GaAs, while low substrate temperature should be used for purely axial growth;
Semiconductor nanostructures of well-defined sizes and shapes are key elements for future nanophotonic devices and integrated circuits [1, 2]. In this field, the achievements obtained over the last two decades has allowed remarkable progress towards the integration of III–V materials on silicon substrates [3, 4]. In particular, the severe constraints posed by the differences in lattice constant, polarity and thermal expansion coefficient can be, in principle, overcome when the interface area between dissimilar semiconductors is reduced to the nanometer scale. Recently, these advantages have been successfully exploited using nanowire (NW) geometry [5, 6], where coherent layers of arbitrary thickness can be grown in a lattice mismatched system, if the NW radius is below a certain critical value (of the order of tens of nm), which depends on the material combination [7]. The technologically relevant case of InAs/GaAs heterostructures has been deeply investigated in planar layers, but realizing its counterpart in such one-dimensional geometry still represents a challenge. Although successful growth by metal-organic vapour phase epitaxy (MOVPE) has been reported [8], it is very difficult to 0957-4484/14/485602+06$33.00
1
© 2014 IOP Publishing Ltd Printed in the UK
C Somaschini et al
Nanotechnology 25 (2014) 485602
furthermore, a low arsenic flux can be used to reduce the amount of parasitic InAs islands directly grown on the substrate. Moreover, when InAs grows axially on top of the GaAs bases its sidewalls are rotated by 30° compared to the underlying GaAs islands, showing an uncommon {112} facet orientation. Finally, the crystallographic structure of purely GaAs islands and InAs/GaAs axial heterostructures was studied by X-ray diffraction. Structurally sensitive Bragg reflections show that the GaAs islands exhibit extended segments of both the cubic zincblende and the hexagonal wurtzite structures with different in-plane lattice parameters. Compared to these, the InAs segments on top of GaAs are fully relaxed and grow mainly in the zincblende phase with a large number of stacking faults.
dimensional pixel detector, reciprocal space maps (RSMs) of different structurally sensitive Bragg reflections were recorded [18].
3. Results and discussion The deposition of Ga on the bare Si (111) surface lead to the formation of Ga droplets of nearly hemispherical shape and with controllable size and number density, as already reported in [15]. The subsequent irradiation with arsenic converted the metallic droplets into GaAs islands with well-defined vertical side-walls and an hexagonal cross-section [16], similar to GaAs NWs, grown by the self-assisted mode [19]. The mean radius and height of the base islands were around 75 ± 8 nm and 100 ± 15 nm, respectively and their number density around 3.0 × 107 cm−2 ± 5 %. In our method, these islands are meant to be used as a base for the overgrowth with InAs, as in a selective area growth (SAG) process, where the material nucleates only in predefined regions of the substrate. In order to investigate the effect of the growth-parameters on the InAs deposition in terms of selectivity and aiming at axial growth, we systematically varied substrate temperatures and arsenic fluxes. This growth study is summarized by the growth diagram shown in figure 1. After the deposition of InAs, we could typically detect two different families of islands on the surface, and it was possible to discriminate between them based on their size. Indeed, smaller islands, with different number density from sample to sample, appeared on the surface, together with larger islands. For all the samples, the number density of the larger islands was nearly constant and matched the that of the original GaAs bases. On the contrary, the radius and/or the height of the larger islands after the growth of InAs greatly exceeded the those of the initial GaAs islands. For example, the large islands on the sample shown in figure 1 (c) have a final radius of 162 ± 16 nm, a height of 109 ± 11 nm and an aspect ratio (height/diameter) of 0.34, while for the sample shown in figure 1(f) the final radius and height are 112 ± 17 nm and 280 ± 35 nm with an the aspect ratio of 1.25. Considering the aspect ratio of the original GaAs bases (0.67), we conclude that in the first case, the islands expanded laterally, while in the second case the expansion occurred mainly axially. These observations imply that, in general, InAs was grown both on the substrate surface, causing the formation of the parasitic, smaller islands and on the GaAs base islands, increasing their final radius and/or height significantly. A closer look at the SEM images shown in figure 1 reveals that the radius of the final InAs/GaAs islands is larger when a higher substrate temperature was used for the growth of the InAs segment. Moreover, the total island number density (including both small and large ones) strongly depends on the growth-parameters, especially on the substrate temperature. The absence of parasitic InAs islands at the highest temperature is caused by the higher adatom mobility together with a reduced sticking coefficient of indium. For a quantitative understanding of the influence of the growth-parameters on the InAs overgrowth, we statistically
2. Experimental details Every growth experiment was performed in a MBE system on Si (111) substrates. This is the most frequently used substrate orientation in the field, as it provides the opportunity to work with vertically grown NWs. Initially, the 2-inch Si wafer was cleaned by the standard piranha solution and a final dip in HF. After that, the substrate was loaded into the MBE chamber, and Ga was deposited at 400 °C in absence of As supply, in order to form Ga droplets, followed by an arsenization step under a beam-equivalent pressure (BEP) of around 1.2 × 10−5 mbar for 10 min, which transformed the Ga droplets into GaAs islands, as for a standard DE process [14]. The substrate was then oxidized under ambient conditions outside the growth chamber and cut in smaller pieces of similar sizes (around 2 × 2 cm) that were dipped in a solution of HCl:H 2 O (1:1) prior to reloading. We performed this treatment to remove the native oxide from the GaAs islands and, at the same time, leave a thin SiO x layer on the substrate, which can act as a mask, in order to reduce parasitic growth [17]. Finally, InAs was deposited by MBE at different substrate temperatures, between 400 °C and 500 °C and at different As BEP, in the range between 1.5 × 10−6 mbar and 1.2 × 10−5 mbar for 60 min, with a deposition rate equivalent to 200 nm/ h for planar growth on a (001) surface. Initially, a test experiment was carried out, annealing a sample with the base GaAs islands at 580 °C in arsenic atmosphere for 10 min, to completely exclude any possible change in the island shape. We did not observe any modification in the sample morphology after the annealing at high temperature and, since the InAs growth on top of these islands was carried out at much lower temperatures, we can assume that the GaAs bases are thermally stable in our range of investigation. The morphology of the samples was characterized after growth by scanning electron microscopy (SEM) and the microstructure was analyzed by transmission electron microscopy (TEM). The average structural composition and the degree of relaxation was determined by asymmetric X-ray diffraction measurements. These experiments were performed at the ID01 beamline at the ESRF synchrotron source using an 9-keV-Xray beam with a size of 200 x 200 μm2. Employing a two2
C Somaschini et al
Nanotechnology 25 (2014) 485602
Figure 1. SEM images (acquired under an angle of 25° to the substrate normal) of InAs/GaAs islands on Si (111) as a function of substrate temperature and As BEP, used during the growth of the InAs segment. In the insets, magnified SEM images of a single structure are depicted. Scale bars are 1 μm in the main images and 100 nm in the insets.
analyzed the samples based on the SEM characterization, and present the results in figure 2. Here, the radius of InAs/GaAs (large) islands and the number density of all the islands found on the substrate surface at the end of the process are plotted as a function of the substrate temperature employed during the InAs deposition. Different symbols are used to distinguish between the As BEP supplied in this step: squares in the case of 1.2 × 10−5 mbar, triangles for 4.5 × 10−6 mbar and diamonds in the case of 1.5 × 10−6 mbar. For the sake of clarity, the same color code is used for the frames in figure 1 and for the symbols in figure 2. For comparison, we indicate the radius and the number density of the original GaAs islands by the dotted line (mean value) and shaded area (including error bars) in the two graphs. As far as the radius of the heterostructures is concerned, figure 2 (a) clearly shows that at higher substrate temperature the final radius of InAs/GaAs islands is much larger than the one of the original GaAs bases. We conclude that, under these conditions, InAs tends to grow mainly around the original GaAs islands, therefore in a core-shell geometry. As the substrate temperature is lowered, the final radius linearly decreases to a value matching the one of the original bases. From this result, we can conclude that a low substrate temperature, around 400 °C, has to be used when aiming at purely axial growth of InAs on top of GaAs. No significant differences in the described trend are observed when different As fluxes were used, except a slight reduction in the InAs/ GaAs island radius for the lowest As BEP. On the other hand, from figure 2 (b) we see that the number density of all the islands present on the substrate surface at the end of the growth procedure strongly increases as the substrate temperature is lowered. As already pointed
out, at 500 °C no parasitic islands are formed on the substrate surface, so that the final and the original (shown by the dashed line) number densities coincide. When approaching the lowest substrate temperature, the total number density increases by nearly two orders of magnitude, meaning that many InAs islands were grown directly on the silicon substrate. Again, the use of different As flux at a constant substrate temperature does not cause a significant change, even though it is clear that lowering the As BEP reduces the amount of InAs islands directly grown on the substrate. Summarizing, we found that the selective deposition of InAs on GaAs islands grown by DE is favored at high substrate temperature, while, to realize a purely axial heterostructure, low substrate temperature should be employed. In the first case, the results are easily understandable in terms of a larger adatom mobility when a higher substrate temperature is used. In the second case, we believe that lower temperatures, more similar to the ones used for the growth of InAs NWs directly on silicon substrate [20], promote the preferential axial growth on top of the GaAs islands. Our growth study also indicates that the use of a low arsenic flux during the growth of InAs is desirable, since it reduces both the heterostructure radius and the amount of parasitic growth. From these results, it becomes clear that a compromise has to be found if aiming at the fabrication of a sample where a low amount of parasitic growth and a good degree of axial growth coexist. A good balance has been found in the sample shown in figure 1(f). This sample has been further analyzed, as we will discuss later. In all the samples where a significant degree of axial growth was obtained (aspect ratio > 0.67), we observed an intriguing feature, clearly visible in the top-view SEM image 3
C Somaschini et al
Nanotechnology 25 (2014) 485602
[112]
a)
[110]
b)
Figure 3. (a) Top-view SEM image of a single, axial InAs/GaAs
Figure 2. Radius of the InAs/GaAs heterostructures (a) and number
heterostructure from the sample shown in figure 1 (f) and (b) TEM image of another structure from the same sample. Scale bars are 100 nm.
density of all the islands obtained after the overgrowth with InAs (b), as a function of the substrate temperature. Symbols are used to distinguish between the As BEP supplied in this step: squares in the case of 1.2 × 10−5 mbar, triangles for 4.5 × 10−6 mbar and diamonds in the case of 1.5 × 10−6 mbar. The radius and number density of the original GaAs base islands are shown in both graphs as dotted lines for comparison, and the shaded area indicate the error bar.
figure 3 (b). In the lower region of the island, vertical sidefacets appear, while in its higher part we recognize several steps, corresponding to a progressive reduction of the radius, up to the flat facet, which is typically formed at the top of selfassisted InAs NWs [20]. The TEM characterization also evidenced that many stacking faults (SFs) are present in the island, with a higher density in the top region. Although SFs are present also in the lower segment of the island, their number density there is lower. This observation suggests that the InAs part of the heterostructure is more defective, as typical in InAs NWs grown in the vapor-solid mode [20, 25, 26]. As far as the crystal structure is concerned, as determined by the TEM investigation, in our axial InAs/GaAs heterostructures extended regions of the two polytypes, zincblende (ZB) and wurtzite (WZ), generally coexist in both materials, with a predominance of ZB in the upper region, where InAs is found, together with SFs, as already mentioned. That a core-shell structure is not clearly visible in figure 3 (b) does not mean that it does not exist in the island. Most likely, the island was not cut in a suitable position close to the center, and/or the shell was, in this case, too thin to be imaged under the employed TEM conditions.
shown in figure 3 (a). When InAs grows on top of the GaAs base islands, it is possible to distinguish two different facet orientations for the two materials. At the bottom, the GaAs islands formed on the silicon substrate by DE expose six facets belonging to the {110} family, as in the case of selfassisted GaAs NWs [21]. In contrast, the InAs segments grown on their tops possess facets belonging to the {112} family, rotated by 30° compared to the underlying GaAs islands. Moreover, the InAs segment grows in a staircase fashion with a decreasing radius as the growth proceeds, while maintaining the same facet orientation. Although no previous reports exist showing the same phenomenon on axial InAs/GaAs heterostructures, similar findings have already been described for InSb/InAs axial [22] and for InAs/GaAs [23] and GaAs/AlInP [24] core-shell NW heterostructures. The SEM observations are further confirmed by the TEM micrograph of another single heterostructure presented in 4
C Somaschini et al
Nanotechnology 25 (2014) 485602
cubic reflections (c) measure the diffracted intensity from domains of the ZB lattice that are twinned and non-twinned with respect to the substrate lattice, while the hexagonal reflection (h) is caused by the wurtzite parts of the structure. The relative intensity of the different reflections contains information about the average structural composition of both GaAs and InAs materials. For GaAs, the WZ type reflection (qz ≈ 47.7 nm−1) has a higher intensity than the two ZB reflections. Because the (422)c reflection of the underlying Si substrate is located in close vicinity in reciprocal space and causes significant diffuse scattering in its surroundings, the GaAs (422)c reflection appears significantly stronger than the (331)c reflection. After integration of the total intensity, subtraction of the background signal and normalization for the respective structural factors, we find that in total (70 ± 2)% of GaAs exhibit the WZ phase. The narrow width of the WZ reflection along qz indicates that this material is present in the form of extended segments, in agreement with the TEM observations shown above. In addition, it is evident that the GaAs WZ reflection in the figure 4 is displaced towards larger qx from the vertical line connecting the two ZB reflections, implying a smaller lattice parameter of WZ GaAs compared to ZB GaAs. From the position of the WZ reflection in the reciprocal space, we obtain lattice parameters of a = (3.986 ± 0.002) Å and c = (6.576 ± 0.002) Å for WZ GaAs, in agreement with values observed before on GaAs NWs grown by the self-assisted growth mechanism [18, 31]. For the InAs segments, a different behaviour is observed. Here, the WZ reflection (qz ≈ 44.8 nm−1) is significantly less intense than the ZB reflections (only one ZB reflection is shown). From the integrated intensity, we estimate that in this case, only (28 ± 2)% of the material is present in the WZ phase. In addition, both ZB and WZ reflections are considerably broader along qz than in the GaAs case, indicating that the mean heights of the defect-free InAs segments of the respective phase are smaller. This agrees with the observation of a large number of SFs in figure 3. In addition, the InAs parasitic islands can also contribute to the signal, as we did not probe a single heterostructure. Compared to the GaAs case, we do not find a significant difference in the in-plane lattice parameters of ZB and WZ InAs, which has been observed before using similar X-ray measurements [32]. We attribute this behavior to the large number of stacking faults and the absence of extended WZ InAs segments. Instead, short WZ segments embedded within a ZB-type structure are strained and adapted to the ZB in-plane lattice parameter. Within the sensitivity of the XRD measurements, we did not find any InxGa1 −x As-related peak, implying that no significant transition layer is formed at the interface between the two materials. Finally, equivalent InAs and GaAs reflections are aligned along radial directions in reciprocal space, visually shown by the inclined dashed-dotted line pointing from the origin of the coordinate frame towards the Si (422)c reflection used as a reference. The inclination of 19.47° with respect to qz measures the angle between the (422) and (111) lattice planes in the crystals of cubic symmetry. Thus, for cubic materials of different lattice parameters, equivalent
Figure 4. X-ray reciprocal space map of axial InAs/GaAs islands
showing the intensity distribution in an asymmetric diffraction geometry. Several sensitive Bragg reflections of GaAs and InAs are accessible, measuring the two possible orientations of the ZB lattice (squares and circles), as well as the hexagonal wurtzite structure (diamonds). While GaAs shows a large abundance of the WZ phase, the dominating InAs structure is ZB. The dash-dotted line indicates the radial direction, showing the relaxed state of InAs.
In epitaxial growth, the crystal phase and orientation are determined by the complex interplay of several contributions and, in nanostructures, the surface energies can play an important role because of the larger surface-to-volume ratio compared to the bulk. In our case, the InAs segment grown on top of the GaAs island shows an uncommon facet orientation. Indeed, InAs NWs with ZB polytype and facets belonging to the {110} family have already been reported [27, 28], as well as InAs NWs with WZ crystal phase, both showing {11 2 0}type [20] and, more frequently, {10 1 0}-type [22, 23, 29, 30] facets. As already described, our InAs segments possess a ZB structure with facets belonging to the {112} family, very rarely reported previously [30]. A possible explanation for this observation can be the low substrate temperature (around 400 °C) we normally employed to achieve a high degree of axial growth, which is believed to favor the formation of {112}-type facets [24, 30]. The average structural composition as well as the lattice parameters of the InAs/GaAs heterostructures were probed by X-ray diffraction experiments both on the initial GaAs islands and on the sample shown in figure 1(f). Figure 4 shows a reciprocal space map of the heterostructure obtained in an asymmetric coplanar diffraction geometry. The x and y axis represent the momentum transfer of the probing X-ray beam along the [11 2] (qx) and [111] direction (qz) of the underlying substrate, respectively. The vertical dashed lines represent truncation rods of GaAs and InAs along which structurally sensitive Bragg reflections of the two materials are observed, i.e. the (331)c, (422)c and (10 1 5)h reflections, respectively. For a fixed azimuthal orientation of the substrate, the two 5
C Somaschini et al
Nanotechnology 25 (2014) 485602
[7] Glas F 2006 Phys. Rev. B 74 2–5 [8] Messing M E, Wong-Leung J, Zanolli Z, Joyce H J, Tan H H, Gao Q, Wallenberg L R, Johansson J and Jagadish C 2011 Nano Lett. 11 3899–905 [9] Paladugu M, Zou J, Guo Y N, Auchterlonie G J, Joyce H J, Gao Q, Tan H H, Jagadish C and Kim Y 2007 Small 3 1873–7 [10] Heiss M, Gustafsson A, Conesa-Boj S, Peiró F, Morante J R, Abstreiter G, Arbiol J, Samuelson L and Fontcuberta i Morral A 2009 Nanotechnology 20 75603 [11] Heiss M, Ketterer B, Uccelli E, Morante J R, Arbiol J and Fontcuberta i Morral A 2011 Nanotechnology 22 195601 [12] Johansson J and Dick K A 2011 Cryst. Eng. Comm. 13 7175 [13] Koguchi N, Takahashi S and Chikyow T 1991 J. Cryst. Growth 111 688 [14] Watanabe K, Koguchi N and Gotoh Y 2000 Jpn. J. Appl. Phys. 39 L79 [15] Somaschini C, Bietti S, Sanguinetti S, Koguchi N, Montalenti F and Frigeri C 2010 Appl. Phys. Lett. 97 053101 [16] Somaschini C, Bietti S, Trampert A, Jahn U, Hauswald C, Riechert H, Sanguinetti S and Geelhaar L 2013 Nano Lett. 13 3607–13 [17] Bietti S, Somaschini C, Frigeri C, Fedorov A, Esposito L, Geelhaar L and Sanguinetti S J. Phys. D: Appl. Phys. 47 394002 [18] Biermanns A, Breuer S, Davydok A, Geelhaar L and Pietsch U 2011 Phys. Status Solidi RRL 5 156 [19] Jabeen F, Grillo V, Rubini S and Martelli F 2008 Nanotechnology 19 275711 [20] Dimakis E, Lähnemann J, Jahn U, Breuer S, Hilse M, Geelhaar L and Riechert H 2011 Cryst. Growth Des. 11 4001–8 [21] Fontcuberta i Morral A, Spirkoska D, Arbiol J, Heigoldt M, Morante J R and Abstreiter G 2008 Small 4 899–903 [22] Lugani L, Ercolani D, Rossi F, Salviati G, Beltram F and Sorba L 2010 Cryst. Growth Des. 10 4038–42 [23] Popovitz-Biro R, Kretinin A A, Huth P V, Shtrikman H and Von Huth P 2011 Cryst. Growth Des. 11 3858 [24] Sköld N, Wagner J B, Karlsson G, Hernán T, Seifert W, Pistol M E and Samuelson L 2006 Nano Lett. 6 2743–7 [25] Koblmüller G, Vizbaras K, Hertenberger S, Bolte S, Rudolph D, Becker J, Döblinger M, Amann M C, Finley J J and Abstreiter G 2012 Appl. Phys. Lett. 101 053103 [26] Dimakis E, Ramsteiner M, Huang C N, Trampert A, Davydok A, Biermanns A, Pietsch U, Riechert H and Geelhaar L 2013 Appl. Phys. Lett. 103 143121 [27] Wei W, Bao X Y, Soci C, Ding Y, Wang Z L and Wang D 2009 Nano Lett. 9 2926–34 [28] Shin J C, Kim K H, Yu K J, Hu H, Yin L, Ning C Z, Rogers J A, Zuo J M and Li X 2011 Nano Lett. 11 4831–8 [29] Tchernycheva M, Travers L, Patriarche G, Glas F, Harmand J C, Cirlin G E and Dubrovskii V G 2007 J. Appl. Phys. 102 094313 [30] Dick K A, Caroff P, Bolinsson J, Messing M E, Johansson J, Deppert K, Wallenberg L R and Samuelson L 2010 Semicond. Sci. Technol. 25 24009 [31] Biermanns A, Breuer S, Trampert A, Davydok A, Geelhaar L and Pietsch U 2012 Nanotechnology 23 305703 [32] Kriegner D et al 2011 Nano Lett. 11 1483–9 [33] Pietsch U, Holý V and Baumbach T 2004 High-Resolution XRay Scattering: From Thin Films to Lateral Nanostructures (New York: Springer)
reflections line up along this direction and a distortion, e.g. due to strain, leads to a shift of the Bragg peak away from this line [33]. The absence of such a signal shows that the InAs segments are grown relaxed on top of the GaAs islands, suggesting that dislocations are formed at the heterointerface to release the strain. This agrees with theoretical calculations, predicting a critical radius of around 20 nm for the dislocation-free growth of InAs on GaAs, exhibiting a 7.1 % lattice mismatch [7].
4. Conclusion In conclusion, we have realized InAs/GaAs heterostructures on silicon, using GaAs islands created by droplet epitaxy as templates for InAs overgrowth. We found that InAs selectively grows on GaAs at high substrate temperature (around 500 °C) and that an axial InAs/GaAs heterostructure can be fabricated at low substrate temperature (around 400 °C). The InAs segments grow axially exposing {112}-type facets which are rotated by 30° compared to the original GaAs bases. As measured by X-ray diffraction, in InAs/GaAs islands ZB and WZ polytypes coexist, with a predominance of WZ in GaAs and of ZB in the InAs segment. The latter is fully relaxed, due to the relatively wide interface between the two materials (larger than 100 nm). Our results represent a step towards the realization of arsenide NW heterostructures integrated on silicon with a switch in the group-III-sublattice.
Acknowledgments The authors would like to acknowledge the financial support from the Deutsche Forschungsgemeinschaft (DFG) under grants Ge2224/2 and Pi 217/38 and from the CARIPLO Foundation under project SOQQUADRO (2012-0362), Claudia Herrmann for the maintenance of the MBE system, Anne-Kathrin Bluhm for the SEM images, Doreen Steffen for the preparation of TEM specimens, T Schülli from beamline ID01, ESRF for support during the synchrotron experiment and Faebian Bastiman for the critical reading of the manuscript.
References [1] del Alamo J A 2011 Nature 479 317–23 [2] Chau R, Doyle B, Datta S, Kavalieros J and Zhang K 2007 Nature Mater. 6 810–2 [3] Moutanabbir O and Gösele U 2010 Annu. Rev. Mater. Res. 40 469 [4] Nah J, Fang H, Wang C, Takei K, Lee M H, Plis E, Krishna S and Javey A 2012 Nano Lett. 12 3592–5 [5] Gudiksen M S, Lauhon L J, Wang J, Smith D C and Lieber C M 2002 Nature 415 617–20 [6] Martensson T, Svensson C P T, Wacaser B A, Larsson M W, Seifert W, Deppert K, Gustafsson A, Wallenberg L R and Samuelson L 2004 Nano Lett. 4 1987–90
6
Search
Collections
Journals
About
Contact us
My IOPscience
Axial InAs/GaAs heterostructures on silicon in a nanowire geometry
This content has been downloaded from IOPscience. Please scroll down to see the full text. 2014 Nanotechnology 25 485602 (http://iopscience.iop.org/0957-4484/25/48/485602) View the table of contents for this issue, or go to the journal homepage for more
Download details: IP Address: 129.97.58.73 This content was downloaded on 16/12/2014 at 16:27
Please note that terms and conditions apply.
Nanotechnology Nanotechnology 25 (2014) 485602 (6pp)
doi:10.1088/0957-4484/25/48/485602
Axial InAs/GaAs heterostructures on silicon in a nanowire geometry C Somaschini1, A Biermanns2, S Bietti3, G Bussone2,4, A Trampert1, S Sanguinetti3, H Riechert1, U Pietsch2 and L Geelhaar1 1
Paul-Drude-Institut für Festkörperelektronik, Hausvogteiplatz 5–7, 10117 Berlin, Germany Universität Siegen, Festkörperphysik, 57068 Siegen, Germany 3 LNESS and Dipartimento di Scienza dei Materiali dellʼUniversità di Milano-Bicocca, via Cozzi 53, 20125 Milano, Italy 4 ESRF, 6 rue Jules Horowitz, BP220, F-38043 Grenoble Cedex, France 2
E-mail: [email protected] Received 7 March 2014, revised 8 October 2014 Accepted for publication 13 October 2014 Published 13 November 2014 Abstract
InAs segments were grown on top of GaAs islands, initially created by droplet epitaxy on silicon substrate. We systematically explored the growth-parameter space for the deposition of InAs, identifying the conditions for the selective growth on GaAs and for purely axial growth. The axial InAs segments were formed with their sidewalls rotated by 30° compared to the GaAs base islands underneath. Synchrotron X-ray diffraction experiments revealed that the InAs segments are grown relaxed on top of GaAs, with a predominantly zincblende crystal structure and stacking faults. Keywords: molecular beam epitaxy, nanowire, heterostructures, X-ray diffraction, droplet epitaxy
1. Introduction
fabricate axial heterostructures with a switch on the group-III sublattice by molecular beam epitaxy (MBE). Issues have been reported in the case of Au-catalyzed growth, where the complex energetics of the ternary Au-Ga-In system can give rise to peculiar phenomena, such as the crawling of the Au droplet on a side facet, which prevents the axial growth [9] and in the case of the self-assisted mode, where the presence of a gallium droplet on top of the GaAs NWs strongly limits the compositional variation along the growth axis [10]. Finally, unintentional radial growth frequently dominates over axial growth [11, 12]. Here, we propose droplet epitaxy (DE) [13, 14] as a possible way to achieve the epitaxial growth of axial InAs/ GaAs heterostructures on silicon substrates in a NW geometry, without the need of any pre-patterning. By DE, GaAs islands of the desired size and number density can be formed on silicon [15, 16] and used as a template for the overgrowth of InAs. We study the influence of the growth-parameters on the deposition of InAs segments on top of the GaAs islands. In particular, we show that high substrate temperature favors the selective deposition of InAs on GaAs, while low substrate temperature should be used for purely axial growth;
Semiconductor nanostructures of well-defined sizes and shapes are key elements for future nanophotonic devices and integrated circuits [1, 2]. In this field, the achievements obtained over the last two decades has allowed remarkable progress towards the integration of III–V materials on silicon substrates [3, 4]. In particular, the severe constraints posed by the differences in lattice constant, polarity and thermal expansion coefficient can be, in principle, overcome when the interface area between dissimilar semiconductors is reduced to the nanometer scale. Recently, these advantages have been successfully exploited using nanowire (NW) geometry [5, 6], where coherent layers of arbitrary thickness can be grown in a lattice mismatched system, if the NW radius is below a certain critical value (of the order of tens of nm), which depends on the material combination [7]. The technologically relevant case of InAs/GaAs heterostructures has been deeply investigated in planar layers, but realizing its counterpart in such one-dimensional geometry still represents a challenge. Although successful growth by metal-organic vapour phase epitaxy (MOVPE) has been reported [8], it is very difficult to 0957-4484/14/485602+06$33.00
1
© 2014 IOP Publishing Ltd Printed in the UK
C Somaschini et al
Nanotechnology 25 (2014) 485602
furthermore, a low arsenic flux can be used to reduce the amount of parasitic InAs islands directly grown on the substrate. Moreover, when InAs grows axially on top of the GaAs bases its sidewalls are rotated by 30° compared to the underlying GaAs islands, showing an uncommon {112} facet orientation. Finally, the crystallographic structure of purely GaAs islands and InAs/GaAs axial heterostructures was studied by X-ray diffraction. Structurally sensitive Bragg reflections show that the GaAs islands exhibit extended segments of both the cubic zincblende and the hexagonal wurtzite structures with different in-plane lattice parameters. Compared to these, the InAs segments on top of GaAs are fully relaxed and grow mainly in the zincblende phase with a large number of stacking faults.
dimensional pixel detector, reciprocal space maps (RSMs) of different structurally sensitive Bragg reflections were recorded [18].
3. Results and discussion The deposition of Ga on the bare Si (111) surface lead to the formation of Ga droplets of nearly hemispherical shape and with controllable size and number density, as already reported in [15]. The subsequent irradiation with arsenic converted the metallic droplets into GaAs islands with well-defined vertical side-walls and an hexagonal cross-section [16], similar to GaAs NWs, grown by the self-assisted mode [19]. The mean radius and height of the base islands were around 75 ± 8 nm and 100 ± 15 nm, respectively and their number density around 3.0 × 107 cm−2 ± 5 %. In our method, these islands are meant to be used as a base for the overgrowth with InAs, as in a selective area growth (SAG) process, where the material nucleates only in predefined regions of the substrate. In order to investigate the effect of the growth-parameters on the InAs deposition in terms of selectivity and aiming at axial growth, we systematically varied substrate temperatures and arsenic fluxes. This growth study is summarized by the growth diagram shown in figure 1. After the deposition of InAs, we could typically detect two different families of islands on the surface, and it was possible to discriminate between them based on their size. Indeed, smaller islands, with different number density from sample to sample, appeared on the surface, together with larger islands. For all the samples, the number density of the larger islands was nearly constant and matched the that of the original GaAs bases. On the contrary, the radius and/or the height of the larger islands after the growth of InAs greatly exceeded the those of the initial GaAs islands. For example, the large islands on the sample shown in figure 1 (c) have a final radius of 162 ± 16 nm, a height of 109 ± 11 nm and an aspect ratio (height/diameter) of 0.34, while for the sample shown in figure 1(f) the final radius and height are 112 ± 17 nm and 280 ± 35 nm with an the aspect ratio of 1.25. Considering the aspect ratio of the original GaAs bases (0.67), we conclude that in the first case, the islands expanded laterally, while in the second case the expansion occurred mainly axially. These observations imply that, in general, InAs was grown both on the substrate surface, causing the formation of the parasitic, smaller islands and on the GaAs base islands, increasing their final radius and/or height significantly. A closer look at the SEM images shown in figure 1 reveals that the radius of the final InAs/GaAs islands is larger when a higher substrate temperature was used for the growth of the InAs segment. Moreover, the total island number density (including both small and large ones) strongly depends on the growth-parameters, especially on the substrate temperature. The absence of parasitic InAs islands at the highest temperature is caused by the higher adatom mobility together with a reduced sticking coefficient of indium. For a quantitative understanding of the influence of the growth-parameters on the InAs overgrowth, we statistically
2. Experimental details Every growth experiment was performed in a MBE system on Si (111) substrates. This is the most frequently used substrate orientation in the field, as it provides the opportunity to work with vertically grown NWs. Initially, the 2-inch Si wafer was cleaned by the standard piranha solution and a final dip in HF. After that, the substrate was loaded into the MBE chamber, and Ga was deposited at 400 °C in absence of As supply, in order to form Ga droplets, followed by an arsenization step under a beam-equivalent pressure (BEP) of around 1.2 × 10−5 mbar for 10 min, which transformed the Ga droplets into GaAs islands, as for a standard DE process [14]. The substrate was then oxidized under ambient conditions outside the growth chamber and cut in smaller pieces of similar sizes (around 2 × 2 cm) that were dipped in a solution of HCl:H 2 O (1:1) prior to reloading. We performed this treatment to remove the native oxide from the GaAs islands and, at the same time, leave a thin SiO x layer on the substrate, which can act as a mask, in order to reduce parasitic growth [17]. Finally, InAs was deposited by MBE at different substrate temperatures, between 400 °C and 500 °C and at different As BEP, in the range between 1.5 × 10−6 mbar and 1.2 × 10−5 mbar for 60 min, with a deposition rate equivalent to 200 nm/ h for planar growth on a (001) surface. Initially, a test experiment was carried out, annealing a sample with the base GaAs islands at 580 °C in arsenic atmosphere for 10 min, to completely exclude any possible change in the island shape. We did not observe any modification in the sample morphology after the annealing at high temperature and, since the InAs growth on top of these islands was carried out at much lower temperatures, we can assume that the GaAs bases are thermally stable in our range of investigation. The morphology of the samples was characterized after growth by scanning electron microscopy (SEM) and the microstructure was analyzed by transmission electron microscopy (TEM). The average structural composition and the degree of relaxation was determined by asymmetric X-ray diffraction measurements. These experiments were performed at the ID01 beamline at the ESRF synchrotron source using an 9-keV-Xray beam with a size of 200 x 200 μm2. Employing a two2
C Somaschini et al
Nanotechnology 25 (2014) 485602
Figure 1. SEM images (acquired under an angle of 25° to the substrate normal) of InAs/GaAs islands on Si (111) as a function of substrate temperature and As BEP, used during the growth of the InAs segment. In the insets, magnified SEM images of a single structure are depicted. Scale bars are 1 μm in the main images and 100 nm in the insets.
analyzed the samples based on the SEM characterization, and present the results in figure 2. Here, the radius of InAs/GaAs (large) islands and the number density of all the islands found on the substrate surface at the end of the process are plotted as a function of the substrate temperature employed during the InAs deposition. Different symbols are used to distinguish between the As BEP supplied in this step: squares in the case of 1.2 × 10−5 mbar, triangles for 4.5 × 10−6 mbar and diamonds in the case of 1.5 × 10−6 mbar. For the sake of clarity, the same color code is used for the frames in figure 1 and for the symbols in figure 2. For comparison, we indicate the radius and the number density of the original GaAs islands by the dotted line (mean value) and shaded area (including error bars) in the two graphs. As far as the radius of the heterostructures is concerned, figure 2 (a) clearly shows that at higher substrate temperature the final radius of InAs/GaAs islands is much larger than the one of the original GaAs bases. We conclude that, under these conditions, InAs tends to grow mainly around the original GaAs islands, therefore in a core-shell geometry. As the substrate temperature is lowered, the final radius linearly decreases to a value matching the one of the original bases. From this result, we can conclude that a low substrate temperature, around 400 °C, has to be used when aiming at purely axial growth of InAs on top of GaAs. No significant differences in the described trend are observed when different As fluxes were used, except a slight reduction in the InAs/ GaAs island radius for the lowest As BEP. On the other hand, from figure 2 (b) we see that the number density of all the islands present on the substrate surface at the end of the growth procedure strongly increases as the substrate temperature is lowered. As already pointed
out, at 500 °C no parasitic islands are formed on the substrate surface, so that the final and the original (shown by the dashed line) number densities coincide. When approaching the lowest substrate temperature, the total number density increases by nearly two orders of magnitude, meaning that many InAs islands were grown directly on the silicon substrate. Again, the use of different As flux at a constant substrate temperature does not cause a significant change, even though it is clear that lowering the As BEP reduces the amount of InAs islands directly grown on the substrate. Summarizing, we found that the selective deposition of InAs on GaAs islands grown by DE is favored at high substrate temperature, while, to realize a purely axial heterostructure, low substrate temperature should be employed. In the first case, the results are easily understandable in terms of a larger adatom mobility when a higher substrate temperature is used. In the second case, we believe that lower temperatures, more similar to the ones used for the growth of InAs NWs directly on silicon substrate [20], promote the preferential axial growth on top of the GaAs islands. Our growth study also indicates that the use of a low arsenic flux during the growth of InAs is desirable, since it reduces both the heterostructure radius and the amount of parasitic growth. From these results, it becomes clear that a compromise has to be found if aiming at the fabrication of a sample where a low amount of parasitic growth and a good degree of axial growth coexist. A good balance has been found in the sample shown in figure 1(f). This sample has been further analyzed, as we will discuss later. In all the samples where a significant degree of axial growth was obtained (aspect ratio > 0.67), we observed an intriguing feature, clearly visible in the top-view SEM image 3
C Somaschini et al
Nanotechnology 25 (2014) 485602
[112]
a)
[110]
b)
Figure 3. (a) Top-view SEM image of a single, axial InAs/GaAs
Figure 2. Radius of the InAs/GaAs heterostructures (a) and number
heterostructure from the sample shown in figure 1 (f) and (b) TEM image of another structure from the same sample. Scale bars are 100 nm.
density of all the islands obtained after the overgrowth with InAs (b), as a function of the substrate temperature. Symbols are used to distinguish between the As BEP supplied in this step: squares in the case of 1.2 × 10−5 mbar, triangles for 4.5 × 10−6 mbar and diamonds in the case of 1.5 × 10−6 mbar. The radius and number density of the original GaAs base islands are shown in both graphs as dotted lines for comparison, and the shaded area indicate the error bar.
figure 3 (b). In the lower region of the island, vertical sidefacets appear, while in its higher part we recognize several steps, corresponding to a progressive reduction of the radius, up to the flat facet, which is typically formed at the top of selfassisted InAs NWs [20]. The TEM characterization also evidenced that many stacking faults (SFs) are present in the island, with a higher density in the top region. Although SFs are present also in the lower segment of the island, their number density there is lower. This observation suggests that the InAs part of the heterostructure is more defective, as typical in InAs NWs grown in the vapor-solid mode [20, 25, 26]. As far as the crystal structure is concerned, as determined by the TEM investigation, in our axial InAs/GaAs heterostructures extended regions of the two polytypes, zincblende (ZB) and wurtzite (WZ), generally coexist in both materials, with a predominance of ZB in the upper region, where InAs is found, together with SFs, as already mentioned. That a core-shell structure is not clearly visible in figure 3 (b) does not mean that it does not exist in the island. Most likely, the island was not cut in a suitable position close to the center, and/or the shell was, in this case, too thin to be imaged under the employed TEM conditions.
shown in figure 3 (a). When InAs grows on top of the GaAs base islands, it is possible to distinguish two different facet orientations for the two materials. At the bottom, the GaAs islands formed on the silicon substrate by DE expose six facets belonging to the {110} family, as in the case of selfassisted GaAs NWs [21]. In contrast, the InAs segments grown on their tops possess facets belonging to the {112} family, rotated by 30° compared to the underlying GaAs islands. Moreover, the InAs segment grows in a staircase fashion with a decreasing radius as the growth proceeds, while maintaining the same facet orientation. Although no previous reports exist showing the same phenomenon on axial InAs/GaAs heterostructures, similar findings have already been described for InSb/InAs axial [22] and for InAs/GaAs [23] and GaAs/AlInP [24] core-shell NW heterostructures. The SEM observations are further confirmed by the TEM micrograph of another single heterostructure presented in 4
C Somaschini et al
Nanotechnology 25 (2014) 485602
cubic reflections (c) measure the diffracted intensity from domains of the ZB lattice that are twinned and non-twinned with respect to the substrate lattice, while the hexagonal reflection (h) is caused by the wurtzite parts of the structure. The relative intensity of the different reflections contains information about the average structural composition of both GaAs and InAs materials. For GaAs, the WZ type reflection (qz ≈ 47.7 nm−1) has a higher intensity than the two ZB reflections. Because the (422)c reflection of the underlying Si substrate is located in close vicinity in reciprocal space and causes significant diffuse scattering in its surroundings, the GaAs (422)c reflection appears significantly stronger than the (331)c reflection. After integration of the total intensity, subtraction of the background signal and normalization for the respective structural factors, we find that in total (70 ± 2)% of GaAs exhibit the WZ phase. The narrow width of the WZ reflection along qz indicates that this material is present in the form of extended segments, in agreement with the TEM observations shown above. In addition, it is evident that the GaAs WZ reflection in the figure 4 is displaced towards larger qx from the vertical line connecting the two ZB reflections, implying a smaller lattice parameter of WZ GaAs compared to ZB GaAs. From the position of the WZ reflection in the reciprocal space, we obtain lattice parameters of a = (3.986 ± 0.002) Å and c = (6.576 ± 0.002) Å for WZ GaAs, in agreement with values observed before on GaAs NWs grown by the self-assisted growth mechanism [18, 31]. For the InAs segments, a different behaviour is observed. Here, the WZ reflection (qz ≈ 44.8 nm−1) is significantly less intense than the ZB reflections (only one ZB reflection is shown). From the integrated intensity, we estimate that in this case, only (28 ± 2)% of the material is present in the WZ phase. In addition, both ZB and WZ reflections are considerably broader along qz than in the GaAs case, indicating that the mean heights of the defect-free InAs segments of the respective phase are smaller. This agrees with the observation of a large number of SFs in figure 3. In addition, the InAs parasitic islands can also contribute to the signal, as we did not probe a single heterostructure. Compared to the GaAs case, we do not find a significant difference in the in-plane lattice parameters of ZB and WZ InAs, which has been observed before using similar X-ray measurements [32]. We attribute this behavior to the large number of stacking faults and the absence of extended WZ InAs segments. Instead, short WZ segments embedded within a ZB-type structure are strained and adapted to the ZB in-plane lattice parameter. Within the sensitivity of the XRD measurements, we did not find any InxGa1 −x As-related peak, implying that no significant transition layer is formed at the interface between the two materials. Finally, equivalent InAs and GaAs reflections are aligned along radial directions in reciprocal space, visually shown by the inclined dashed-dotted line pointing from the origin of the coordinate frame towards the Si (422)c reflection used as a reference. The inclination of 19.47° with respect to qz measures the angle between the (422) and (111) lattice planes in the crystals of cubic symmetry. Thus, for cubic materials of different lattice parameters, equivalent
Figure 4. X-ray reciprocal space map of axial InAs/GaAs islands
showing the intensity distribution in an asymmetric diffraction geometry. Several sensitive Bragg reflections of GaAs and InAs are accessible, measuring the two possible orientations of the ZB lattice (squares and circles), as well as the hexagonal wurtzite structure (diamonds). While GaAs shows a large abundance of the WZ phase, the dominating InAs structure is ZB. The dash-dotted line indicates the radial direction, showing the relaxed state of InAs.
In epitaxial growth, the crystal phase and orientation are determined by the complex interplay of several contributions and, in nanostructures, the surface energies can play an important role because of the larger surface-to-volume ratio compared to the bulk. In our case, the InAs segment grown on top of the GaAs island shows an uncommon facet orientation. Indeed, InAs NWs with ZB polytype and facets belonging to the {110} family have already been reported [27, 28], as well as InAs NWs with WZ crystal phase, both showing {11 2 0}type [20] and, more frequently, {10 1 0}-type [22, 23, 29, 30] facets. As already described, our InAs segments possess a ZB structure with facets belonging to the {112} family, very rarely reported previously [30]. A possible explanation for this observation can be the low substrate temperature (around 400 °C) we normally employed to achieve a high degree of axial growth, which is believed to favor the formation of {112}-type facets [24, 30]. The average structural composition as well as the lattice parameters of the InAs/GaAs heterostructures were probed by X-ray diffraction experiments both on the initial GaAs islands and on the sample shown in figure 1(f). Figure 4 shows a reciprocal space map of the heterostructure obtained in an asymmetric coplanar diffraction geometry. The x and y axis represent the momentum transfer of the probing X-ray beam along the [11 2] (qx) and [111] direction (qz) of the underlying substrate, respectively. The vertical dashed lines represent truncation rods of GaAs and InAs along which structurally sensitive Bragg reflections of the two materials are observed, i.e. the (331)c, (422)c and (10 1 5)h reflections, respectively. For a fixed azimuthal orientation of the substrate, the two 5
C Somaschini et al
Nanotechnology 25 (2014) 485602
[7] Glas F 2006 Phys. Rev. B 74 2–5 [8] Messing M E, Wong-Leung J, Zanolli Z, Joyce H J, Tan H H, Gao Q, Wallenberg L R, Johansson J and Jagadish C 2011 Nano Lett. 11 3899–905 [9] Paladugu M, Zou J, Guo Y N, Auchterlonie G J, Joyce H J, Gao Q, Tan H H, Jagadish C and Kim Y 2007 Small 3 1873–7 [10] Heiss M, Gustafsson A, Conesa-Boj S, Peiró F, Morante J R, Abstreiter G, Arbiol J, Samuelson L and Fontcuberta i Morral A 2009 Nanotechnology 20 75603 [11] Heiss M, Ketterer B, Uccelli E, Morante J R, Arbiol J and Fontcuberta i Morral A 2011 Nanotechnology 22 195601 [12] Johansson J and Dick K A 2011 Cryst. Eng. Comm. 13 7175 [13] Koguchi N, Takahashi S and Chikyow T 1991 J. Cryst. Growth 111 688 [14] Watanabe K, Koguchi N and Gotoh Y 2000 Jpn. J. Appl. Phys. 39 L79 [15] Somaschini C, Bietti S, Sanguinetti S, Koguchi N, Montalenti F and Frigeri C 2010 Appl. Phys. Lett. 97 053101 [16] Somaschini C, Bietti S, Trampert A, Jahn U, Hauswald C, Riechert H, Sanguinetti S and Geelhaar L 2013 Nano Lett. 13 3607–13 [17] Bietti S, Somaschini C, Frigeri C, Fedorov A, Esposito L, Geelhaar L and Sanguinetti S J. Phys. D: Appl. Phys. 47 394002 [18] Biermanns A, Breuer S, Davydok A, Geelhaar L and Pietsch U 2011 Phys. Status Solidi RRL 5 156 [19] Jabeen F, Grillo V, Rubini S and Martelli F 2008 Nanotechnology 19 275711 [20] Dimakis E, Lähnemann J, Jahn U, Breuer S, Hilse M, Geelhaar L and Riechert H 2011 Cryst. Growth Des. 11 4001–8 [21] Fontcuberta i Morral A, Spirkoska D, Arbiol J, Heigoldt M, Morante J R and Abstreiter G 2008 Small 4 899–903 [22] Lugani L, Ercolani D, Rossi F, Salviati G, Beltram F and Sorba L 2010 Cryst. Growth Des. 10 4038–42 [23] Popovitz-Biro R, Kretinin A A, Huth P V, Shtrikman H and Von Huth P 2011 Cryst. Growth Des. 11 3858 [24] Sköld N, Wagner J B, Karlsson G, Hernán T, Seifert W, Pistol M E and Samuelson L 2006 Nano Lett. 6 2743–7 [25] Koblmüller G, Vizbaras K, Hertenberger S, Bolte S, Rudolph D, Becker J, Döblinger M, Amann M C, Finley J J and Abstreiter G 2012 Appl. Phys. Lett. 101 053103 [26] Dimakis E, Ramsteiner M, Huang C N, Trampert A, Davydok A, Biermanns A, Pietsch U, Riechert H and Geelhaar L 2013 Appl. Phys. Lett. 103 143121 [27] Wei W, Bao X Y, Soci C, Ding Y, Wang Z L and Wang D 2009 Nano Lett. 9 2926–34 [28] Shin J C, Kim K H, Yu K J, Hu H, Yin L, Ning C Z, Rogers J A, Zuo J M and Li X 2011 Nano Lett. 11 4831–8 [29] Tchernycheva M, Travers L, Patriarche G, Glas F, Harmand J C, Cirlin G E and Dubrovskii V G 2007 J. Appl. Phys. 102 094313 [30] Dick K A, Caroff P, Bolinsson J, Messing M E, Johansson J, Deppert K, Wallenberg L R and Samuelson L 2010 Semicond. Sci. Technol. 25 24009 [31] Biermanns A, Breuer S, Trampert A, Davydok A, Geelhaar L and Pietsch U 2012 Nanotechnology 23 305703 [32] Kriegner D et al 2011 Nano Lett. 11 1483–9 [33] Pietsch U, Holý V and Baumbach T 2004 High-Resolution XRay Scattering: From Thin Films to Lateral Nanostructures (New York: Springer)
reflections line up along this direction and a distortion, e.g. due to strain, leads to a shift of the Bragg peak away from this line [33]. The absence of such a signal shows that the InAs segments are grown relaxed on top of the GaAs islands, suggesting that dislocations are formed at the heterointerface to release the strain. This agrees with theoretical calculations, predicting a critical radius of around 20 nm for the dislocation-free growth of InAs on GaAs, exhibiting a 7.1 % lattice mismatch [7].
4. Conclusion In conclusion, we have realized InAs/GaAs heterostructures on silicon, using GaAs islands created by droplet epitaxy as templates for InAs overgrowth. We found that InAs selectively grows on GaAs at high substrate temperature (around 500 °C) and that an axial InAs/GaAs heterostructure can be fabricated at low substrate temperature (around 400 °C). The InAs segments grow axially exposing {112}-type facets which are rotated by 30° compared to the original GaAs bases. As measured by X-ray diffraction, in InAs/GaAs islands ZB and WZ polytypes coexist, with a predominance of WZ in GaAs and of ZB in the InAs segment. The latter is fully relaxed, due to the relatively wide interface between the two materials (larger than 100 nm). Our results represent a step towards the realization of arsenide NW heterostructures integrated on silicon with a switch in the group-III-sublattice.
Acknowledgments The authors would like to acknowledge the financial support from the Deutsche Forschungsgemeinschaft (DFG) under grants Ge2224/2 and Pi 217/38 and from the CARIPLO Foundation under project SOQQUADRO (2012-0362), Claudia Herrmann for the maintenance of the MBE system, Anne-Kathrin Bluhm for the SEM images, Doreen Steffen for the preparation of TEM specimens, T Schülli from beamline ID01, ESRF for support during the synchrotron experiment and Faebian Bastiman for the critical reading of the manuscript.
References [1] del Alamo J A 2011 Nature 479 317–23 [2] Chau R, Doyle B, Datta S, Kavalieros J and Zhang K 2007 Nature Mater. 6 810–2 [3] Moutanabbir O and Gösele U 2010 Annu. Rev. Mater. Res. 40 469 [4] Nah J, Fang H, Wang C, Takei K, Lee M H, Plis E, Krishna S and Javey A 2012 Nano Lett. 12 3592–5 [5] Gudiksen M S, Lauhon L J, Wang J, Smith D C and Lieber C M 2002 Nature 415 617–20 [6] Martensson T, Svensson C P T, Wacaser B A, Larsson M W, Seifert W, Deppert K, Gustafsson A, Wallenberg L R and Samuelson L 2004 Nano Lett. 4 1987–90
6
