Structural Phase Control in Self-Catalyzed Growth of GaAs Nanowires

Oct 8, 2010 - Peter Krogstrup,*,† Ronit Popovitz-Biro,§ Erik Johnson,† Morten Hannibal Madsen,†. Jesper Nygård,† and Hadas Shtrikman*,‡...
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Structural Phase Control in Self-Catalyzed Growth of GaAs Nanowires on Silicon (111) Peter Krogstrup,*,† Ronit Popovitz-Biro,§ Erik Johnson,† Morten Hannibal Madsen,† Jesper Nygård,† and Hadas Shtrikman*,‡ † ‡

Nano-Science Center, Niels Bohr Institute, University of Copenhagen, Denmark and § Electron Microscopy Unit and Braun Center for Submicron Research, Weizmann Institute of Science, Rehovot 76100, Israel ABSTRACT Au free GaAs nanowires with zinc blende structure, free of twin planes and with remarkable aspect ratios, have been grown on (111) Si substrates by molecular beam epitaxy. Nanowires with diameters down to 20 nm are obtained using a thin native oxide layer on the Si substrates. We discuss how the structural phase distribution along the wire length is controlled by the effective V/III ratio and temperature at the growth interface and explain how to obtain a pure twin plane free zinc blende structure. KEYWORDS Nanowire, molecular beam epitaxy, wurtzite, zinc blende, twin planes, GaAs

I

n recent years, III-V nanowires (NWs) grown on Si substrates have received increasing attention,1-15 due to the superior electronic properties of direct band gap III-V materials combined with the low cost and well-known properties of Si. Still, to achieve full benefit from these advantages better understanding of the growth mechanisms and thus better control of the NW properties, namely, morphology, structure, and purity, are essential. Au-assisted III-V NWs have previously been grown on Si substrates (MBE,5-8 MOCVD9-12), but the incorporation of the Au, which may create deep levels acting as recombination centers, degrading the electronic properties of the material, is still ambiguous.16,17 Formation of Au-free GaAs NWs grown by molecular beam epitaxy (MBE) has recently been demonstrated with a thin deposited layer of SiO2 on Si or GaAs substrates,2,4,13,15,18,19 where the obvious advantage is the Au free and particularly low C environment. Prior to actual NW growth, initial Ga adatoms are believed to accumulate at existing defects or sub nanometer pin holes in the amorphous oxide layer. The accumulation leads to larger nanosize Ga1-xAsx droplets,2,18,20,21 which then act as small liquid phase epitaxy systems promoting NW growth. In this letter, we present growth of Au-free GaAs NWs with a large aspect ratio, grown on untreated Si(111) substrates by the vapor-liquid-solid mechanism and we explain how to obtain a ZB structure free of twin planes (TP). All NWs presented here were grown by solid source MBE in two different systems, a Riber 32 system at the Weizmann Institute of Science, and a Varian GEN II system at the Niels Bohr Institute. All substrates used were undoped and epiready 2 in. Si wafers with a (111) surface orientation and were degassed in a buffer chamber prior to growth.

The diameter of the NWs is controlled by the size of the droplet and shows a remarkable uniformity (within a given region on the substrate, affected mostly by the temperature), varying by no more than 10%. This suggests that once the Ga droplets have reached a certain (temperature and partial pressure dependent) critical size, they become supersaturated enough to initiate epitaxial NW growth, most likely through oxide dissolution at the droplet/substrate interface. Below T ∼ 590 °C, no wires but only GaAs chunks were observed on the substrate, and for growths above T ∼ 670 °C, no GaAs nucleation (neither chunks nor wires) was observed, indicating that the amount of Ga sticking on the oxide is too low and reflecting a significant temperature dependence. Nonvertical wires are a typical problem when growing III-V NWs on Si, as shown in Figure 1a. This is not only due to the four possible nonpolar low energy directions on Si, namely, one vertical to the substrate and three nonvertical ones, but also due to the possibility of initiating nonvertical NW growth on the oxide layer without direct contact with the underlying Si.18 At low temperatures, 590-630 °C, the fraction of vertical wires was observed to decrease with increasing temperature, while at high growth temperatures, 640-660 °C, the fraction of vertical wires increased with temperature. While the fraction of vertical wires follows a nonmonotonic temperature dependence, vertical wires were generally favored at high V/III ratios for the whole temperature range (see Supporting Information, S1). This is an indication that the effective V/III ratio (defined as the ratio of V/III materials sorbed in the droplet) plays an important role on the initial conditions and most likely follows a nonmonotonic temperature dependence through the adatom/admolecule kinetics. A possible explanation could be as follows: in the high temperature range where the Ga accumulation (or the effective collection area of Ga adatoms, AIII,s) on the oxide decreases with increasing temperature, the As is mostly absorbed directly in the droplets. Thus, in this temper-

* To whom correspondence should be addressed. E-mail: (P.K.) [email protected]; (H.S.) [email protected]. Received for review: 07/1/2010 Published on Web: 10/08/2010 © 2010 American Chemical Society

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FIGURE 2. NWs with a diameter of ∼20 nm and length of ∼7 µm were produced on a thin oxide layer. Only the lowest sections of the NWs are shown. The scalebar is 200 nm.

on a hot plate at T ) 100 °C, which provided an oxide layer of approximately half the thickness of a fully grown native oxide layer.22 This resulted in wire diameters as small as 18-20 nm, as shown in Figure 2. Thus, the critical droplet size is smaller on a thinner native oxide, which seems reasonable since the droplet has to dissolve less material before reaching contact with the underlying crystalline Si. The structure of III-V NWs is typically composed of both cubic and hexagonal phases due to the small difference in the internal energy of the two phases. However, it is possible to favor one of the phases by tuning the growth conditions. Here we show and explain how to get pure ZB NWs without twin planes, which to the best of our knowledge has not been reported before for Au-free growth of III-V NWs. The cubic ZB structure has an ABCABC stacking sequence in the 〈111〉 growth direction, while the hexagonal WZ structure has an ABABAB stacking along the respective 〈0001〉 direction. Thus, a so-called rotational twin plane, ABCABC-A-CBACBA, will be regarded as being a single layer of WZ sandwiched in between two mirror symmetric ZB segments. Common for all growths is that ZB is the dominating structure with {11¯0} side faceting. In order to explain the mechanisms determining the NWs structure, we start with an analysis of the effect of two important parameters: the effective V/III ratio in the liquid (described by the As atomic fraction, xV), and temperature, T, at the growth interface. To estimate the temperature decrease at the growth interface in relation to the substrate surface, we directly used the heat transfer calculation approach of the GlasHarmand model,23 which is plotted in Figure 3a as a function of the NW length. The largest uncertainty associated with this approach is the value of the thermal conductivity in thin NWs. The thermal conductivity was predicted in ref 24 by considering a GaAs nanobeam with a cross section of 100 × 250 nm2. We have used their theoretically extrapolated values even though the value might change significantly for smaller wires, as discussed in ref 25, and it is therefore likely that the decrease

FIGURE 1. Two SEM images of the same growth with a V/III ratio of ∼160 but taken at different regions on the wafer. The difference observed at the two regions is solely due to the difference in temperature. However, the large fraction of nonvertical wires in (a) is explained primarily by the temperature dependent effective V/III ratio at the surface. In (b), the wires are ∼40 nm in diameter, which is due to a large Ga desorption from oxide surface, making the effective V/III ratio large. The temperature gradient across the wafer during growth stems from a nonhomogenous heat input and was measured with a pyrometer. The scale bar is 5 µm and applies for both images.

ature range the effective V/III ratio will increase with increasing temperature because AIII,s becomes smaller. This increases the liquid supersaturation and thus makes it less probable to alloy and form nonvertical facets at the liquid-substrate interface, implying a higher fraction of vertical wires as seen in Figure 1b. Conversely, in the low temperature regime, AIII,s is increasing with increasing temperature, because of an increase in the effective diffusion length of the Ga adatoms. Thus in this temperature range the effective V/III ratio will decrease with increasing temperature and therefore imply a decrease in the fraction of vertical wires (we will see later that the NW structure in the beginning of the NW growth is also controlled mainly by AIII,s). Thin wires are generally formed at high V/III ratios and high temperature. The thinnest NW diameters obtained on a fully oxidized native oxide layer were approximately 40 nm with a V/III ratio of ∼160, as shown in Figure 1b (and Supporting Information S2b). However, at the same growth conditions even thinner NWs were grown on a thinner, prepared oxide layer. This was achieved by heating the substrates prior to growth in a separate treatment chamber at T ) 1100 °C. During this procedure the oxide layer was totally desorbed (as confirmed by clear reflection high energy electron diffraction streaks). Subsequently, the substrate was unloaded and kept in a clean hood for an hour © 2010 American Chemical Society

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FIGURE 3. Key parameters for self-assisted growth of GaAs NWs, T and xV, are naturally changed at the growth interface during growth. (a) The Harmand-Glas NW heat transfer model predicts a not negligible temperature gradient for thin and long NWs. The longer the NW becomes during growth the lower the temperature at the growth interface. (b) A plot of how the bulk liquid supersaturation as a function of xV would look like (i.e., µsurface ≈ µsurface,eq). For small xV values higher temperatures imply a smaller undercooling effect on the driving force. The inset in (b) shows ∆µ as a function of xV up to 0.5. Because of the lower interaction energies of the As species, mixing of group III and V gives a maximum ∆µ at xV ≈ 0.3. Values for Li and µi are listed in the Supporting Information, S4.

∆µ = µl - µs ) (µvolume + µmix + µsurface) (µvolume,eq + µmix,eq + µsurface,eq)

where µvolume is given by the sum of the pure components, µi, times their respective effective molar fractions, xV, in the liquid (i.e.µvolume ) xIIIµIII + xVµV). The second term µmix takes account of the enthalpy and entropy of mixing group III and V elements. Using the Redlich-Kister formalism with two binary liquid interaction parameters,28-30 Li, we account for the asymmetry of the compositional effect of the internal energy µmix ) xIIIxV[L0 + L1(xIII - xV)] + RT(xIII ln(xIII) + xV ln(xV))

FIGURE 4. (a) Schematic illustration of the Ga-assisted growth model illustrating the different effective collection areas for adatoms contributing to the NW growth. The As atomic fraction xV in the liquid depends on these areas which changes as the NW grows longer, see text for explanation.

where R is the gas constant. The equilibrium atomic fraction xV,eq (or equilibrium liquid solubility) of As species in liquid Ga is found to follow xV,eq ) 6.752 × 10-7 exp(0.0141(T + 273)) in the temperature range T ) 400-800 °C.30 Because of the rather complex nature of segregation and diffusion at the interfaces as a function of xV and temperature, xV is for simplicity assumed to be constant along the liquid-vapor interface and therefore also at the triple phase line where the nucleation most likely occurs.31,32 ∆µ is plotted in Figure 3b as a function of xV where the change in liquid surface free energy and Gibbs-Thomson effect is neglected, µsurface ≈ µsurface,eq. The change in the total number of atoms in the droplet during growth can be expressed as

in T is more significant. Thus, it is possible that the temperature gradient plays an important role for thin and long wires. Nevertheless, the temperature plays an important role on the surface kinetics for all diameters, and therefore on the composition at the growth interface. The liquid supersaturation, ∆µ ) µ1 - µ1,eq (where µ1 and µ1,eq are the specific free energies of the supersaturated liquid and the liquid in equilibrium with the solid, respectively), is likely to play an important role on the structural phase. To obtain a simple and useful expression for ∆µ as a function of xV we assume that the change in Gibbs free energy and surface contributions of the solid phase from equilibrium to supersaturated conditions are negligible. Using this, the supersaturation at the growth interface is as follows26,27 © 2010 American Chemical Society

(

)

xV 1 - xV d + - Iinc n(t) ) IIII + IV - Adns dt τIII,l τV,l

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FIGURE 5. A TEM image of a NW (bottom image) after Ga flux and temperature optimization (step 2) with final parameters of fGa ) 9.25 × 10-8 Torr, fAs2 ) 1.3 × 10-5 Torr, T ) 660 °C. The images above are segments chosen along the length with the respective electron diffraction patterns.

where Ii and Ad(xins)/(τi,l) account for the liquid sorption current and liquid evaporation current of the ith element, respectively. Ad is the liquid-vapor surface area, ns is the atomic surface density and τi,l is the characteristic lifetime between evaporation events at the liquid-vapor surface. Iinc is the incorporation current of atoms in the solid. The NW diameter d(t) is determined by the size of the Ga droplet, which contrary to Au-catalyzed growth can change significantly during growth, with the possibility of obtaining reverse tapering33 of the NWs at large Ga beam flux, fGa. However, in the high-temperature range (where the radial growth is negligible), the wires are nontapered for © 2010 American Chemical Society

a wide range of V/III ratios. Therefore, as a rough approximation we will assume that on a time scale larger than the time taking to form a monolayer, the apparent contact angle and therefore the number of particles is approximately constant, (d/dt)n(t) ≈ 0. It is known that the vapor pressure of Ga is orders of magnitude lower than As, which evaporates from GaAs surfaces in the high temperature range, leaving liquid Ga droplets on the surface.34 Hence, it is reasonable to assume (xV)/(τV,l) . (1 - xV)/(τIII,l) and IIII ≈ (1/2)Iinc at relatively high supersaturation conditions, and eq 1 can be rewritten as follows 4478

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xV ≈

(IV - IIII)τV,l Adns

increasing supersaturation. Furthermore, assuming total liquid coverage of the growth interface, it also suggests that for extremely thin NWs the nucleation area size effect may become important, favoring WZ formation at very small diameters. This is a complex situation since the NW diameter decreases with increasing V/III ratio and temperature. Accordingly, while higher V/III ratio and temperature promote small NW diameter and ZB formation, a very small diameter itself may favor TP (or WZ) formation. Thus, it is possible that at some point the NW becomes so thin that the diameter effect will dominate and the TP probability will increase with V/III ratio. The density of TPs (or WZ structure) is generally observed to be highest in the beginning and in the end of the growth, consistent with previous reports on the structural distribution along the NWs.13,36 To explain this, we have to look at the adatom kinetics. AV,s can safely be neglected because As sticking on the oxide is zero, and AV,w is extremely small at the growth temperatures used.39,40 Thus if the NWs are not tapered, Al is approximately constant during growth, and so is IV. This means, according to eq 2, that the change in xV will be due to a change in IIII and possibly τV,l at constant incoming fluxes, fi. However, a change in τV,l would be due to a change in IIII through a change in the liquid chemical potential, µl. But since µl is an increasing function of xV (for xV < 0.3, which is within the range we expect it to be), it is reasonable to analyze the change in IIII using eq 3 in order to understand whether xV is increasing or decreasing. In the high-temperature regime, Ga desorption from the oxide is high and the contribution from AIII,s will not be as dominant, while AIII,w will consist of both direct and secondary adsorption. As the NW grows longer AIII,s will decrease due to the longer path to the growth interface and AIII,w will increase due to the increasing NW length for L < λIII,w. AIII,s will be dominant in the beginning of growth, therefore implying an increase in xV as the NW grows. However, AIII,w will after some length start to increase IIII and xV will at this point decrease, as the NW grows. These predictions are consistent with the structural distribution seen along the NWs, if TP formation is favored at low xV. We propose a procedure that has turned out to be adequate for obtaining vertical and nontapered NWs with a high aspect ratio and a pure ZB structure without TPs. (1) First, reduce the NW diameter and the fraction of nonvertical wires by minimizing fIII until there is not enough excess Ga on the surface to nucleate NW growth. Then increase fIII so as to obtain a sufficient density of wires. (2) Maximizing the substrate temperature until GaAs sticking approaches zero. Choose the threshold value still producing a reasonable NW density. A transmission electron microscopy (TEM) image of a NW obtained after this step is shown in Figure 5, the diameters of the wires are d ∼ 45-50 nm, and the lengths are L ∼ 14-15 µm for a growth time of 45 min. TPs are present at the top couple of micrometers

(2)

Equation 2 is sufficient to explain qualitatively how a change in xV depends on the sorption currents, which are given by

Ii(T, L) ) fi

∑ Ki,j Ai,j(T, L)

(3)

j

Here the sum over j accounts for the different kinetic properties of the substrate (s), the NW sidewalls (w) and the liquid (l), with Ki,j being the adsorption coefficient and Ai,j(T,L) the respective effective collection area of the ith element (L is the NW length). Based on the simple geometrical considerations illustrated in Figure 4, the three different effective collection areas for direct adsorption events of group i adatoms are given as Ai,s(T,L ) 0) ) πλ2i,s (T) cos(φi), Ai,w(T,L ) λi,w) ) dλi,w (T) sin(φi), and Al ) (πd2)/4, where λi,s and λi,w are the effective diffusion lengths on the substrate and NW sidefacets, respectively. φi is the angle of the incoming beam. However, growth on the oxide surface increases not only the desorption of As adatoms but also Ga adatoms, and AIII,w becomes important after the NW has reached a certain length due to secondary adsorption.35 A recent study by Spirkoska et al.36 showed that increasing fV while keeping other parameters constant (which apparently means increasing ∆µ) resulted in an increase in the fraction of ZB structure. This is in contrast to previous predictions using the highly approximate classical twodimensional (2D) nucleation theory,31,37,38 which has been used to describe the phase transition at the growth interface for Au-assisted NWs. Spirkoska et al. argued that a possible reason for their findings could be that the vapor pressure has a large effect on the effective edge energies of the 2D nucleus that could affect the WZ and ZB nucleation rates and therefore possibly favor ZB at large ∆µ. They also point out the possibility that because the critical nuclei become smaller with larger supersaturation, it implies a relatively larger area for ZB nucleation events than WZ, since ZB is favored for nucleation away from the triple phase line (TPL) and WZ at the TPL. The most obvious difference from the well studied Au-assisted system is the interfacial energy values associated with the liquid droplet,31,34 which has a significant effect on the structural phase probability as a function of ∆µ. Even though the sharp edge 2D nucleation theory approach is approximate in the way it is used and with the lack of knowledge of the systems specific energy values, we have plotted its structural phase prediction as function of xV in the Supporting Information, Figure S3a. The theory suggests that above a critical value of xV, the TP probability decreases with © 2010 American Chemical Society

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FIGURE 6. Long and thin NWs with ZB structure is achieved after optimizing the Ga flux, temperature, and As flux, grown with the final parameters, fGa ) 9.25 × 10-8 Torr, fAs2 ) 1.0 × 10-5 Torr, T ) 660 °C. See discussion for an explanation regarding the twin planes at the bottom of the NW.

FIGURE 7. (a) A TEM image of a GaAs NW grown with a V/III ratio of 20 and a temperature of 620 °C. The convergent beam electron diffraction patterns in the insets below stems from nanosize regions along the wire as shown. The relatively large ZBa and ZBb parts at the NW root are separated by a twin plane, which rotates the structure by 180° around the wire axis. The pattern corresponding to the ZBa/ZBb region reveals closely separated ZBa and ZBb parts and the density of twin planes is seen to increase as we move toward the tip. Finally, the structure keeps rotating at every monolayer that produces the small segment with the WZ structure. At the given growth conditions, the effective V/III ratio at the growth interface is most likely increasing as the NW grows longer (see the discussion). (b) A top view SEM image of the same growth illustrating the massive chunk like bulk growth in between the NWs.

as well as one or two of them occasionally along the clean part of the NW. (3) Finally, and surprisingly, we observed that the few TPs vanished by lowering fV slightly. This resulted in a decrease in the NW length and slight increase in the diameter. A TEM image of a NW obtained after this step is shown in Figure 6. The growth time was 45 min and the diameters for this sample are d ∼ 55-60 nm and the lengths are L ∼ 8-9 µm. There can be different explanations of why fewer TPs are observed at lower fV in this regime. A likely explanation for this could be that the driving force is approaching the limit for 2D nucleation where the growth interface will progres© 2010 American Chemical Society

sively roughen due to kinetic reasons, leading to a more diffuse interface and lower step edge energy.41 Other explanations could be due to the nucleation area size effect (Supporting Information, Figure S3b) or possibly also due to a higher temperature at the growth interface (Figure 3a). Finally we should mention the possibility that the 2D WZ nucleus probability could be increasing again as a function of xV due to a change in surface energies. An advantage obtained by approaching the extremes of the growth window is the low NW density that assures no interaction between the wires and therefore provides compatibility with processes requiring external coating of the NWs. 4480

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ported by the Israeli Science Foundation Grant 530/08 and the Israel Ministry of Science Grant no. 3-66799.

To remove the TPs appearing at the root of the NWs shown in Figure 6 and 7 would require a reduction of AIII,s which again would either require a higher temperature or a lower fGa. This is a difficult task since these parameters are already close to the threshold value of initiating NW growth, however it should in principle be possible. λIII,s ) (DIIIτIII)1/2 is the key parameter in controlling the TP concentration at the beginning through AIII,s at all temperatures. DIII and τIII are the effective diffusivity and lifetime of the Ga adatoms on the substrate, respectively, where the latter is given by 1/τIII ) 1/τIII,des + 1/τIII,inc with τIII,des and τIII,inc being the respective lifetimes before desorption or incorporation. It is reasonable to assume that τIII = τIII,des at high temperatures, but τIII,inc becomes important (i.e., small) at low temperatures since GaAs chunk nucleation occurs everywhere on the oxide surface, see Figure 7b. The contribution from AIII,s is therefore small at low temperatures due to a small τIII,inc. Because AIII,w increases as the NW grows longer it begins to increase IIII after a certain length, and so does the TP probability. This structural distribution is seen in Figure 7a where the density of TPs increases toward the nanowire tip. Looking at the structural change for the relatively thick NW in this figure and comparing it with the temperature decrease shown in Figure 3a, it seems reasonable to neglect the liquid undercooling dependence on the structural outcome due to a temperature gradient. In conclusion, this study has yielded well-controlled growth of pure ZB GaAs nanowires with a large aspect ratio without using gold as a catalyst. A study of the mechanisms controlling the MBE growth of self-catalyzed GaAs NWs has been presented. By looking at the structural change, we explain how the growth conditions naturally change as the NWs grow longer. We conclude that the primary effect for this change is the temperature related compositional change in the liquid during growth, while the mere temperature change along the NW length only plays a minor role for thin and long wires.

Supporting Information Available. Information regarding obtaining vertical wires, obtaining uniform and thin wires, nucleation model, specific energy values used in the paper, and additional references. This material is available free of charge via the Internet at http://pubs.acs.org. REFERENCES AND NOTES (1) (2) (3) (4) (5) (6) (7) (8) (9) (10) (11)

(12) (13) (14) (15) (16)

Acknowledgment. We are grateful to Moty Heiblum, head of the Braun Center for submicron research, for making the NW research at the Braun Center possible. We are thankful to Perla Kacman for fruitful discussions and to Martin Aagesen and SunFlake A/S for supporting this project. Claus B. Sørensen and Michael Fourmansky are thanked for assistance on the MBE systems at the Niels Bohr and the Weizmann Institute of Science, respectively. The electron microscopy studies were conducted at the Irving and Cherna Moskowitz Center for Nano and Bio-Nano Imaging of the Weizmann Institute of Science, and at the Nano-Science Center, Niels Bohr Institute. We acknowledge HyungKook Choi for helping with SEM imaging. We thank Andrey Kretinin for continuous discussions and support in the NWs project at the Braun Center and Stefano Curiotto, for help in finding thermodynamic data on specific energies. This work was supported by the Danish National Advanced Technology Foundation through Project 002-20091, and the Danish Stategic Research Council through project 09-065736. The work at the Braun Center was partially sup© 2010 American Chemical Society

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DOI: 10.1021/nl102308k | Nano Lett. 2010, 10, 4475-–4482