Self-Assisted Nucleation and Vapor–Solid Growth of InAs Nanowires

ARTICLE pubs.acs.org/crystal

Self-Assisted Nucleation and Vapor
Solid Growth of InAs Nanowires on Bare Si(111) Emmanouil Dimakis,* Jonas L€ahnemann, Uwe Jahn, Steffen Breuer, Maria Hilse, Lutz Geelhaar, and Henning Riechert Paul-Drude-Institut f€ur Festk€orperelektronik, Hausvogteiplatz 5-7, 10117 Berlin, Germany

bS Supporting Information ABSTRACT: The nucleation and growth of InAs nanowires on bare Si(111) has been investigated by molecular beam epitaxy. Nontapered InAs nanowires with high aspect ratio were grown perpendicular to the substrate without the use of catalyst particles, surface oxide, or other substrate mask. The nucleation of InAs takes place in In-rich areas forming spontaneously on the substrate in the beginning of the growth process. As the nucleation proceeds, the local stoichiometry on the growth interface changes from In-rich to As-rich, and the growth continues in a vapor
solid mode. This transition to As-rich conditions is correlated with the evolution of nanowire morphology, that is, with the growth becoming strictly uniaxial and with well-defined vertical sidewalls forming. The diameter, the number density, and the axial growth rate of the nanowires were found to depend exclusively on the surface diffusivity of In adatoms on the substrate.

’ INTRODUCTION The epitaxial growth of III
V nanowires (NWs) on Si substrates is of great interest, because the two complementary technologies can thus be integrated on single multifunctional chips combining the superior electronic and optoelectronic properties of the former with the mature CMOS (complementary metal oxide semiconductor) technology of the latter.1,2 In particular, InAs is considered as a promising material for highspeed electronics due to its high electron mobility, as high as ∼30 000 cm2/(V s) at 300 K.3 Furthermore, the same material system is very important to study because it is one of the constituent binary alloys of InxGa1
xAs, which is technologically attractive for NW photonic applications due to its direct band gap energy, which can be engineered to be anywhere between 0.36 and 1.42 eV at 300 K. Several prototype devices based on III-As NWs on Si have already been demonstrated,4
9 but still, most of the research efforts have been focusing on understanding the growth mechanisms of GaAs and InAs NWs on Si(111). The Au-assisted vapor
liquid
solid (VLS) growth10
13 is a well-established and well-controlled technique that has been widely used for the preparation of NWs in various material systems, as well as in III-As.13
19 Nevertheless, the use of Au nanoparticles is not suitable for the growth of NWs on a CMOS platform, since Au is incompatible with the CMOS processing standards.20,21 In addition, the use of any foreign nanoparticles may result in the degradation of the NW material itself.22,23 Instead, III-As NWs can be grown in a self-induced manner without using any catalysts.4,24
38 Satisfactory results have been obtained mostly when a Si-oxide (SiOx) layer with small holes masks the substrate. These holes are spontaneously or lithographically r 2011 American Chemical Society

opened, and the NWs nucleate inside them. In the case of GaAs, the presence of Ga droplets at the NW tips clearly evidence a Gadroplet-mediated VLS growth mode. In the case of InAs, In droplets have never been observed at the NW tips after the end of the growth. However, it is not straightforward to determine whether In droplets exist during growth because they may easily be desorbed or consumed by excess As during termination of growth. The contradictory reports found in the literature reflect the fact that the growth mechanism of InAs NWs is not well understood: it is not clear whether the formation of InAs NWs results from an In-droplet-mediated VLS34 or a droplet-free vapor
solid (VS)4,36 growth mode. Obviously, the type of growth mode significantly influences the properties of the resulting NWs and by what means NW position, diameter, and crystal structure can be controlled. Therefore, the goal of this work is to develop a better understanding of the growth mechanism in order to improve the growth control and the crystal quality of the InAs NWs on Si. Aiming to elucidate the Au-free growth mechanism of InAs NWs, we present growth studies conducted on bare Si(111) substrates by molecular beam epitaxy (MBE), using neither any catalyst particles nor SiOx or other coating layers. In this way, NW growth is investigated under the least extrinsic restrictions possible. As a matter of fact, just a very limited number of reports on the growth of III-As NWs on oxide-free substrates exist in the literature,4,32 and in any case, the growth has not been investigated systematically. In the present work, we find that two Received: May 4, 2011 Revised: June 22, 2011 Published: June 24, 2011 4001

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distinct mechanisms describe the nucleation and the growth of InAs NWs on Si; the nucleation is associated with the formation of In-rich areas, while the subsequent NW formation and growth proceeds under As-rich conditions in a VS mode. It is also found that the surface diffusivity of In adatoms on Si is of major importance, since it determines the number density, the diameter, and the axial growth rate of the NWs.

’ EXPERIMENTAL DETAILS To identify the growth mechanism of InAs NWs on Si, the role of the most decisive growth parameters, namely, the growth temperature (Tsub), the atomic flux of In (FIn), the flux ratio (FAs/FIn), and the growth duration (tGR), was studied independently in a systematic way. To that end, we have chosen to grow four series of samples (A, B, C, and D). Series A consists of samples grown at various Tsub in the range from 425 to 475
C, using FIn = 100 nm/h and FAs/FIn = 120 (the corresponding flux of As, FAs, was equal to ∼12 μm/h). The impinging fluxes are given as equivalent nominal growth rates of planar InAs layers. For the effusion cell of As, the cracker zone was kept at low enough temperatures where no cracking of As4 molecules occurs. Series B consists of samples grown at 450
C, varying FIn in the range from 13 to 190 nm/h. FAs was adjusted accordingly, so FAs/FIn was approximately equal to 120 in all cases. Series C consists of samples grown at 450
C, using FIn = 100 nm/h and values of FAs in the range from 3.6 to 31 μm/h. The resulting FAs/FIn varied in the range from 36 to 310. The tGR value for each sample of series A, B, and C was adjusted in a way that the total amount of deposited In was the same (equivalent to a 200 nm thick InAs planar layer). Finally, series D consists of samples grown at 450
C with FIn = 100 nm/h and FAs/FIn = 120 for various tGR, from 5 s to 2 h. For all samples, the growth was started by first opening the shutter of the As effusion cell and then that of In. The growth was ended by closing both As and In shutters simultaneously; tGR corresponds to the time that the shutter of the In effusion cell was open. The Si(111) substrates were etched in 5% HF for 1 min to remove the native oxide, in H2SO4/H2O2 (3:1) solution for 10 min to remove the organic residuals, in 40% NH4F for 10 min to etch the oxide formed during the previous step and create atomically flat (111) terraces, and in 5% HF for 10 s to passivate the surface with hydrogen (H-terminated).39 After each step, the substrates were rinsed with deionized H2O. At the end, they were dried with N2 and loaded into the MBE system within 10 min to avoid reoxidation of the surface. Prior to the growth, the substrate temperature was first increased in situ to more than 500
C and then adjusted to the appropriate growth temperature in the range 425
475
C. The pregrowth annealing was proven to be essential for the subsequent successful growth of NWs, since otherwise (if the substrate is set directly to the growth temperature) a continuous polycrystalline InAs layer would always form on the substrate apart from the very few NW-like structures. We believe that the annealing to more than 500
C is necessary for the removal of hydrides from the Si surface40 that obstruct the epitaxial overgrowth. The strong effect of hydrides on the epitaxial growth of InAs and its suppression by annealing the substrate to more than 500
C serve as an indirect confirmation that the Si substrate surface is still H-terminated after its transfer into the growth chamber, and no SiOx layer exists. Tsub was measured using a pyrometer, while reflection high-energy electron diffraction (RHEED) oscillations on InAs(001) were employed to calibrate FAs and FIn with respect to the beam equivalent pressures of As (PAs) and In (PIn), respectively. As a reference, a planar InAs layer was grown on InAs substrate with a rate of 100 nm/h at Tsub < 400
C when either PIn = 1.5  10
7 mbar was used under As-rich conditions or PAs = 3.1  10
7 mbar was used under stoichiometric conditions. Here, FAs/FIn is calculated as the ratio of FAs and FIn at low Tsub, but its actual value is expected to be reduced when Tsub > 400
C due to the sharp

Figure 1. SEM images of InAs NWs grown on Si(111) at 450
C with FIn = 190 nm/h and FAs/FIn =120. (a) Side-view image showing that nontapered NWs grew perpendicular to the substrate. (b) Side-view close-up on the NW base showing no diameter widening close to the interface with the substrate. (c) Top-view image showing the hexagonal cross section of NWs with side facets parallel to the {110} planes of the Si substrate. (d) Side-view close-up on the NW tips showing flat top facets parallel to the (111) plane of the Si substrate and the absence of In droplets. The scale bar is 1 μm for panel a and 200 nm for panels b
d. increase of the desorption of As species from the InAs surface. The desorption of As was evaluated based on the RHEED pattern transition from the (2  4) As-stable surface reconstruction to the (4  2) In-stable one.41 This concerned desorption from an InAs surface, and potential differences may apply for the Si(111) surface. The morphological features of NW ensembles were characterized by scanning electron microscopy (SEM). The number density (NNW) and the diameter (DNW) were measured for InAs NW ensembles contained in (25  18) μm2 top-view images, while the length (LNW) was measured from 25 μm-wide side-view images. Fluctuations in LNW due to different nucleation times were taken into account using Gaussian distribution curves for the statistical analysis. The axial growth rate (GRNW) for long NWs was calculated as their total length divided by the total growth duration (LNW/tGR). SEM electron backscatter diffraction (SEM-EBSD) was employed to determine the crystal structure of the NWs and the crystallographic orientation of the side facets.42 For that purpose, the NWs were mechanically dispersed onto a Au-covered Si wafer to avoid shadowing effects that are typically present when measuring NW ensembles. The amorphous 50 nm-thick Au layer was sputtered on Si to prevent diffraction from the latter and to limit the sample charging. The electron beam was accelerated to 15 keV and the incidence angle was 70
.

’ NANOWIRE MORPHOLOGY AND CRYSTAL STRUCTURE Columnar InAs structures were obtained under all growth conditions used in series A, B, and C. The side-view SEM image of a typical sample of NWs with high aspect ratio is presented in Figure 1a. As seen, all NWs grew perpendicular to the substrate surface and, apart from a small number of structures with lower aspect ratio (islands), InAs did not grow in any other form 4002

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Figure 2. (a) SEM image of a NW that was transferred onto a Au-coated Si wafer and measured by EBSD. The scale bar is 200 nm. (b) EBSD pattern acquired from the location marked with “” in panel a and fitted using Kikuchi lines for WZ InAs. (c) EBSD pattern as corrected for the 70
incident angle, revealing {2110} sidewalls.

between the NWs. This is observed in the close-up SEM image on the NW base in Figure 1b. Even though the presence of a very thin planar InAs layer that covers the substrate surface uniformly cannot be ruled out just based on the SEM images, the absence of respective streaky diffraction features in the RHEED pattern confirms that this is not the case at least during the first several minutes of growth. Diffraction features from a planar InAs layer should have appeared after the first ∼10 s of growth, since the nominal growth rate was one monolayer every 10 s as defined by the In flux used in most of our experiments. In contrast, diffraction from the uncovered Si substrate was sustained until the latter was completely shadowed by the growing NWs. Coming back to the NW morphology, DNW of each NW is observed to be constant throughout the entire NW length. As it will be shown later, this is a result of the efficient surface diffusion of In adatoms from the substrate, along the side facets, to the top of the growing NWs. All NWs have a hexagonal cross section as evidenced from the top-view SEM image in Figure 1c, where the side facets are parallel to the {110} planes of the Si substrate. Finally, as seen in Figure 1d, the NWs end with a flat top facet, which is parallel to the (111) substrate surface, and no In droplets exist at the NW tips after removal from the growth system. The crystal structure of our NWs was identified by SEMEBSD to be wurtzite (WZ). An example of a NW that was measured to that end is depicted in Figure 2a. The EBSD pattern shown in Figure 2b was acquired from the central part of one of the side facets, as indicated by the symbol “” in Figure 2a. The pattern was successfully fitted using the Kikuchi lines of a WZ InAs crystal, and the diffraction spots were indexed correspondingly (not shown here). After correcting the Kikuchi pattern for the tilt of 70
, it was revealed that the NW side facets have the Æ2110æ crystallographic orientation (Figure 2c). Our results differ from the zinc-blende (ZB) structure with {110} side facets typically reported for self-induced InAs NWs grown either on bare Si(111)4 or on SiOx/Si(111).34 Up to now, InAs NWs with large WZ segments have only been observed for Au-assisted NW growth and then always with {1100} side facets.16,43 The {2110} side facets are scarcely reported in the literature.44 The origin of the formation of {2110} side facets in our case is not clear, but we assume it is related to the MBE conditions and the oxide-free substrate surface used. For a consistency check, the complete surface area of the NW side facet was scanned by SEM-EBSD with a spatial resolution of 20 nm, showing that the crystal structure is WZ throughout the entire length of the NW. Nevertheless, inclusions of small ZB

)

)

bands, narrower than 20 nm, cannot be excluded. Combining the SEM and the SEM-EBSD results, the epitaxial relationship between the InAs NWs and the Si substrate is determined to be {0001}InAs {111}Si and {2110}InAs {110}Si.

’ VAPOR
SOLID GROWTH MODE The NNW, DNW, and GRNW values exhibited strong dependence on the growth conditions as shown in Figure 3, where they are plotted as a function of Tsub for series A, FIn for series B, and FAs/FIn for series C. For series A (Figure 3a), when higher Tsub was used (in the range from 425 to 475
C), NNW decreased monotonically by 2 orders of magnitude, while both DNW and GRNW increased by a factor of 2.5
3.0. For series B (Figure 3b), when higher FIn was used (in the range from 13 to 190 nm/h), NNW increased by 2 orders of magnitude, DNW decreased by a factor of 2.5, and GRNW increased by a factor of 3. In principle, for diffusion-limited growth mechanisms, the use of higher growth temperature or lower growth rate leads to nucleation of fewer, but larger, structures. This has been theoretically predicted and experimentally observed for various Inrelated systems, for example, the formation of liquid In droplets on foreign substrates45,46 or the nucleation of InN on GaN(0001).47,48 Evidently, the nucleation of InAs NWs on Si is also limited by the diffusivity of In adatoms on the substrate surface. The importance of the diffusivity of In on Si, as well as on the NW side facets, is also suggested by the fact that GRNW is always higher than FIn but lower than FAs. After all, it is implied that in order to grow InAs NWs with small diameter, growth conditions that limit the surface diffusivity of In adatoms on Si need to be employed. Still, the diffusivity along the side facets needs to remain efficient, and typically this is true for Tsub > 400
C and up to the longest NWs of the present work (∼ 6 μm), as it will be shown in the next section. For sample series C (Figure 3c), almost no changes were observed when FAs/FIn was increased from 120 (FAs = 12 μm/h) to 310 (FAs = 31 μm/h). Only when a lower FAs of 3.6 μm/h (FAs/FIn = 36) was used, NNW decreased by 2 orders of magnitude and DNW increased by a factor of 2.5. Apparently, the diffusivity of In adatoms on Si becomes so high when FAs/FIn 2 min. The high GRNW that exceeds FAs in the beginning of the growth manifests that the local V/III ratio at the growth interface of the just nucleated NWs is significantly modified in favor of In (In-rich conditions), as will be explained in the following. The high GRNW requires increased amount of both In and As species arriving on the nucleating NWs for a given growth duration. The increased amount of In adatoms can be understood as the result of their efficient diffusivity from the substrate surface onto the nucleating NWs. On the other hand, the increased amount of As has two possible explanations. Similarly to In adatoms, the first explanation is the significant diffusion of As species from the substrate surface onto the NWs. However, the pre-existence of excess In atoms on the NWs (In-terminated growth interface) is a prerequisite for GRNW to exceed FAs, given that As does not agglomerate at as high Tsub as the ones used here. The second possible explanation is the extended supply of As, since As4 is present in the growth chamber even after terminating the growth process by closing the shutters of the effusion cells. In that case, it is implied that a significant excess of In (possibly in a form of a droplet) is present on the nucleating NW, which is consumed by As after terminating the growth process. No matter which one of the two explanations for the increased amount of As is valid, the nucleation of InAs on Si(111) evidently occurs under In-rich conditions. The interplay between the NW nucleation and the concentration of In adatoms has already been pointed out by the results described in Figure 3a,b; growth conditions that limit the diffusivity of In adatoms on the Si substrate led to nucleation of a large number of thin NWs. The nucleation of InAs on Si(111) under In-rich conditions is in contrast to the subsequent VS growth under As-rich conditions that was pointed out in the previous section. This transition of the local V/III ratio on the growth interface from In-rich to Asrich is evidenced by the fast decrease of GRNW to a value lower than FAs, which completes within the first 40
50 s of growth (Figure 4b). Furthermore, this is the point in time when the

nucleated NWs start developing well-defined vertical sidewalls (Figure 4a, 43 s), and their growth becomes strictly uniaxial along the substrate normal (no significant change in DNW is observed after the first ∼1 min of growth). Thus, it is suggested that the characteristic NW morphology develops only after the local conditions on the NW growth interface change from In-rich to As-rich. The transition in local V/III ratio, evidenced by the decrease of GRNW with tGR (alternatively, the sublinear increase of LNW with tGR), can be due to several reasons. The most obvious one is the simultaneous increase of DNW with tGR; given that the arrival rate of the constituent species on each InAs NW does not increase with tGR, GRNW is expected to decrease as long as DNW increases. Indeed, this is in agreement with our results shown in Figure 4b. The second reason for the sublinear increase of LNW with tGR is the simultaneous increase of NNW. That is, fewer In adatoms arrive on each NW as new NWs that compete for the available diffusing adatoms nucleate on the substrate. The decrease of the arrival rate of In adatoms on each NW as NNW increases is also evidenced by the corresponding decrease of the volume growth rate per NW (Supporting Information, Figure S4). When the arriving flux is eventually stabilized, the increase of LNW with tGR becomes linear (and the volume growth rate per NW constant). Thus, the primary reason of the transition from In-rich to As-rich conditions on the NW growth interface is the decreasing arrival of In adatoms from the substrate due to the ongoing nucleation of new NWs. In contrast, in ref 50, the sublinear increase of LNW with tGR in the beginning of NW growth has been attributed to the lower diffusivity of the constituent species along the NW sidewalls than on the substrate surface, so that fewer adatoms diffuse all the way from the substrate to the NW tips as the NWs grow longer. Finally, after the NWs have grown longer than the diffusion length on the NW sidewalls, only the species impinging close to the NW tip contribute to the axial growth, stabilizing in that way the axial growth rate. If this model was valid in our case, then the diffusion length on the sidewalls, Ldiff, should be equal to LNW at 4005

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the time when its temporal dependence transforms from a sublinear to a linear one (at tGR ≈ 2 min), that is, Ldiff ≈ 0.4 μm. To test whether this can be the case, we calculate Ldiff of In adatoms in an alternative way that should lead to the same result. We assume that only the amount of In impinging directly on the sidewalls and within one diffusion length away from the NW tip can contribute to the axial growth, as schematically depicted in Figure 4c. The projection of this effective sidewall area, Aeff, on the plane normal to the In beam is given by Aeff ¼ 2RNW Ldiff cos θ sin j

ð1Þ

where RNW is the NW radius (RNW = DNW/2), θ is the in-plane angle of the sidewall normal with respect to the beam direction, and j is the angle of incidence of the beam on the substrate. The value of θ ranges from 0 to 30
due to the in-plane 6-fold symmetry of NWs, while j = 30
as defined by the design geometry of our growth chamber. The impinging flux of In on the effective sidewall area corresponds to an equivalent volume growth rate, VGRsidewalls, which is given by the expression in eq 1 multiplied by the impinging rate Fimp (Fimp = FIn/cos j, where FIn is the equivalent growth rate of a planar layer): VGR sidewalls ¼ 2RNW Ldiff cos θ tan j FIn

ð2Þ

Assuming that all In adatoms diffuse from the effective sidewall area to the tip, VGRsidewalls must be equal to the volume growth rate by which NWs grow axially, VGRaxial. The volume growth rate due to direct impingement of In from the vapor onto the top NW facet is negligible (∼23 times lower) compared with VGRaxial. The latter is given by ! pffiffiffi 3 3 2 RNW VGR axial ¼ GR NW ð3Þ 2 where the factor in the parenthesis corresponds to the surface area of the top facet. Equating eq 2 with eq 3, Ldiff is found to be pffiffiffi 3 3GR NW RNW ð4Þ Ldiff ¼ 4 cos θ tan j FIn For GRNW = 2.3 μm/h, RNW = 0.1 μm, FIn = 0.1 μm/h, and 0 e θ e 30
, Ldiff ranges from 5.2 to 6.0 μm. This means that on the basis of a limited-sidewall-diffusion model, an axial growth rate of 2.3 μm/h can only be obtained for NWs longer than 5.2
6 μm. This is much different from the results shown in Figure 4b, where an axial growth rate of 2.3 μm/h was already obtained for 0.4 μm long NWs. Such high axial growth rate for such short NWs clearly suggests that it is not only the In adatoms impinging directly on the NW sidewalls that contribute to the axial growth of NWs. A more significant contribution exists and cannot be other than the efficient diffusion of In adatoms from the substrate surface (after impingement from the vapor) along the NW sidewalls to the NW tip, where they eventually incorporate in the InAs crystal. According to this explanation, the diffusivity of In adatoms along the NW sidewall certainly cannot be limited. Similar calculations were repeated for all samples of series A, B, and C, and in all cases Ldiff was found to be almost equal to or higher than the final NW length. This means that for any of the growth conditions used in this work, the diffusivity of In along the NW sidewalls never limited the axial growth rate. Focusing again on the initial In-rich local conditions, it is still not clear in which phase the excess In that promotes the nucleation of InAs NWs on Si exists. As discussed earlier, this

Figure 5. RHEED pattern evolution for InAs NWs on Si(111) after 0, 5, 40, and 200 s of growth. Before the growth starts (0 s), diffraction only from the Si substrate is visible. Diffraction spots from InAs, typical of WZ crystal structure, emerge after initiating the growth (5 s). After 40 s, streaky diffraction features from InAs also appear, indicating the formation of a high density of stacking faults. After 200 s, diffraction only from InAs is visible, and the pattern remains unchanged for the rest of the growth time.

could be either an In-terminated growth interface or an In droplet. Even though In droplets are never observed at the NW tips by SEM after growth, even when tGR < 2 min, their presence in the NW nucleation stage cannot be ruled out. This is because the In droplets may have been consumed by the As4 still being present in the growth chamber even after terminating the growth by simultaneously closing In and As shutters. Such consumption of the In droplets would result in the effective lengthening of the NWs after closing the In and As shutters. Since this phenomenon is more significant for shorter NWs, one could also consider it as a potential source of the increased GRNW for tGR < 2 min. Nevertheless, this is not in quantitative agreement with our results. In particular, the effective lengthening for tGR ≈ 2 min should be insignificant because GRNW stabilizes at that point (2.3 μm/h). However, after 2 min of growth the NWs are still much longer (∼400 nm) than what would be expected for NWs growing with a constant rate of 2.3 μm/h (∼80 nm). Therefore, the most important factor causing the significantly higher initial GRNW is the initially higher arrival rate of In adatoms, and on the basis of the present data, we cannot conclude whether In droplets exist during the nucleation stage. Additional information about the nucleation mechanism is retrieved from the temporal evolution of the RHEED pattern, with the e-beam along the Æ110æ crystallographic orientation of the Si substrate, as shown in Figure 5. As soon as the nucleation of InAs begins, a typical WZ diffraction pattern appears, while diffraction from the Si substrate is still present (Figure 5, 0 and 5 s). As the growth proceeds, streaky diffraction features from the InAs crystal also appear in the pattern (Figure 5, 40 and 200 s), 4006

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forming spontaneously on Si in the beginning of the growth process mediate the nucleation of InAs. As the nucleation continues and more InAs structures form on the substrate, the concentration of arriving In adatoms on each one of them decreases rapidly, and soon As-rich conditions are established at the growth interface. From that point on, the already nucleated InAs structures continue growing uniaxially in the VS mode, developing the characteristic NW morphology.

Figure 6. Schematic representation of the growth model for InAs NWs on bare Si(111). Temporal evolution from panel a
d: (a) Formation of In-rich areas (possibly In droplets) on Si. (b) Excess-In-mediated nucleation of InAs. (c) The decreasing flux of arriving In adatoms at the growth interface, due to the ongoing formation of new InAs nuclei on Si, causes the fast transition to As-rich conditions. (d) Under As-rich conditions, the growth of InAs proceeds uniaxially in VS mode, and vertical sidewalls develop.

evidencing the formation of a high density of planar defects (stacking faults) oriented perpendicular to the growth direction. The existence of a high density of stacking faults was also observed by X-ray diffraction measurements that will be published elsewhere. The transformation of the diffraction pattern lasts for the first 40
50 s of growth, coinciding with the period of time during which GRNW exceeds FAs. Thus, the onset of stacking fault formation is attributed to the transition of the local stoichiometry in the beginning of the growth from In-rich to As-rich. Certainly, this has to be investigated in more detail using high-resolution transmission electron microscopy (HR-TEM) on single NWs. The fact that In-rich conditions result in the formation of WZ InAs with fewer stacking faults was also observed for a sample grown in two steps. The nucleation was performed under typical conditions (Tsub = 450
C, FIn = 100 nm/h, FAs/FIn = 120) for the first 2 min, while almost equal fluxes of In and As4 were used (by decreasing the latter) for the subsequent 1 h of growth. This change of flux ratio was intended to create local In-rich conditions on the already nucleated NWs. Indeed, the WZ diffraction pattern recovered soon after changing the fluxes (the streaky diffraction features from InAs disappeared, Supporting Information, Figure S5a) and was preserved throughout the growth. It should be also mentioned that the NW morphology was dramatically affected by using In-rich conditions, resulting in severe widening of the bottom part of the initial InAs NWs and degradation of their sidewalls (Supporting Information, Figure S5b). This is an additional indication that the development of NW morphology is a result of the transition to As-rich conditions. According to a theoretical description of the self-assisted growth of GaAs NWs by Krogstrup et al.,51,52 the WZ phase is thermodynamically favored for low supersaturation values of the Ga droplet with As, that is, for low atomic fraction of As in the Ga droplet. Departure from the pure WZ phase can be due to either the increase of the supersaturation or the shrinkage of the liquid droplet (detachment of the droplet from the edges of the top facet). Both cases can be interpreted as transitions from groupIII-rich to As-rich conditions, in agreement with our understanding of the nucleation of InAs NWs. To summarize our model, the nucleation and the growth of InAs NWs on Si(111) are described by two distinct mechanisms, as illustrated in Figure 6. In-rich areas (potentially In droplets)

’ CONCLUSIONS Wurtzite InAs NWs with Æ2110æ oriented sidewalls were grown perpendicular to the bare Si(111) substrate in the temperature range from 425 to 475
C, only under As-rich conditions. Their nucleation occurs in In-rich areas spontaneously formed on Si, while their growth proceeds in VS mode under As-rich conditions. The diffusivity of In adatoms on Si governs the nucleation and the growth kinetics, limiting the axial growth rate, the number density, and the diameter of the NWs. The ability to grow InAs NWs without the presence of In droplets on the NW tips is expected to be advantageous in the realization of axial heterostructures, where abrupt compositional and doping changes are desirable. ’ ASSOCIATED CONTENT

bS

Supporting Information. Additional experiments aiming to confirm the validity of the VS growth mode of InAs NWs on Si(111), as well as the interplay between the In-rich conditions and the formation of WZ structure, and calculation of the volume growth rate of InAs NWs as a function of tGR for sample series D. This material is available free of charge via the Internet at http:// pubs.acs.org.

’ AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected].

’ ACKNOWLEDGMENT We thank C. Herrmann and A.-K. Bluhm for technical and SEM support. This work was partly supported by Deutsche Forschungsgemeinschaft (DFG) under Grant Ge 2224/2. ’ REFERENCES (1) Bakkers, E. P. A. M.; Borgstr€om, M. T.; Verheijen, M. A. MRS Bull. 2007, 32, 117. (2) Yang, P.; Yan, R.; Fardy, M. Nano Lett. 2010, 10, 1529. (3) Rode, D. L. In Semiconductors and Semimetals; Willardson, R. K., Beer, A. C., Eds.; Academic Press: New York, 1975; Vol. 10, p 1. (4) Wei, W.; Bao, X. Y.; Soci, C.; Ding, Y.; Wang, Z.
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