Kinetic Effects in InP Nanowire Growth and Stacking Fault Formation

Apr 18, 2011 - Impact of nucleation conditions on diameter modulation of GaAs nanowires. Samuel C Crawford , Sema Ermez , Georg Haberfehlner , Eric J ...
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LETTER pubs.acs.org/NanoLett

Kinetic Effects in InP Nanowire Growth and Stacking Fault Formation: The Role of Interface Roughening Thalita Chiaramonte,† Luiz H. G. Tizei,†,‡ Daniel Ugarte,† and Mo^nica A. Cotta*,† † ‡

Instituto de Física“GlebWataghin”, Universidade Estadual de Campinas, UNICAMP, 13083-859, Campinas, SP, Brazil Laboratório Nacional de Luz Síncrotron, C P 6192, 13083-970 Campinas, SP, Brazil

bS Supporting Information ABSTRACT: InP nanowire polytypic growth was thoroughly studied using electron microscopy techniques as a function of the In precursor flow. The dominant InP crystal structure is wurtzite, and growth parameters determine the density of stacking faults (SF) and zinc blende segments along the nanowires (NWs). Our results show that SF formation in InP NWs cannot be univocally attributed to the droplet supersaturation, if we assume this variable to be proportional to the ex situ In atomic concentration at the catalyst particle. An imbalance between this concentration and the axial growth rate was detected for growth conditions associated with larger SF densities along the NWs, suggesting a different route of precursor incorporation at the triple phase line in that case. The formation of SFs can be further enhanced by varying the In supply during growth and is suppressed for small diameter NWs grown under the same conditions. We attribute the observed behaviors to kinetically driven roughening of the semiconductor/metal interface. The consequent deformation of the triple phase line increases the probability of a phase change at the growth interface in an effort to reach local minima of system interface and surface energy. KEYWORDS: Nanowire, InP, crystal structure, zinc blende, wurtzite, stacking fault, electron microscopy

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emiconductors nanowires (NWs) have been extensively studied due to their potential applications in the new generation of electronics, photonics and sensing devices.1 4 As for optoelectronics, III V NWs present many advantages regarding emission spectra range and carrier mobilities. These NWs have been successfully synthesized by the vapor liquid solid (VLS)5 mechanism using Au nanoparticles (NPs) as catalysts.6 While bulk materials and thin epitaxial films present zinc blende (ZB) as the stable crystallographic phase, III V NWs usually present an structure where crystalline planes can be stacked as ZB or wurtzite (WZ), even if a ZB substrate is used.7 9 The formation of WZ NWs is favorable when surface energy outweighs the volume cohesive energies; this leads to a critical diameter below which the WZ phase is dominant.10 In general, the transition from WZ to ZB creates stacking faults (SFs) in NWs grown along the [111] direction. The energy barrier for the WZ/ ZB transition is rather low, estimated as 3.4 meV/atom for InP NWs with diameters greater than 20 nm.10 Recently, a satisfactory control of both SF occurrence and the formation of pure and mixed WZ/ZB segments has been achieved by changing growth parameters and nanoparticle diameter.11,12 By varying the crystal structure between ZB and WZ, it is possible to develop band gap engineering in a single material along the NW axial length.12 Despite these very promising results, the physical mechanisms behind the growth of a determined crystal structure are not yet fully understood. Moreover, polytypic behavior shows striking similarities for several III V materials despite the use of different substrates and growth techniques.7 9,13 15 Recently, it has been r 2011 American Chemical Society

shown that the SFs density and polytypism along the NW depend on growth conditions, such as temperature, V/III ratio, and NP diameter.12,16 18 For example, Paiman et al.19 reported that the InP NWs grown by metalorganic metal vapor deposition (MOCVD) are preferably formed in WZ phase with increasing V/III ratios (under fixed In precursor flow) or decreasing NP diameter. Also, Dick et al.,20 using metalorganic vapor phase epitaxy (MOVPE), have observed a transition from WZ at low TMI molar fraction to a mixture of WZ and ZB at either very high molar fraction or reduced V/III ratio. These results show that surface kinetics plays a dominant role in polytypism even though detailed surface conditions depend on the employed growth technique. On the basis of nucleation kinetic models, the growth of different crystal structures relies on the small energy difference between the two critical nuclei. At higher growth temperatures or supersaturation conditions, larger fluctuations are expected; these factors can lead for example to the formation of both WZ and ZB sections in NWs grown by VLS. Glas et al.21 have proposed a nucleation-based model where the nucleus is formed at the triple phase line (TPL); under certain supersaturation conditions, the WZ phase could be favored over ZB. More complete models also take into account the distortion of the NP due to the nucleus formation at the TPL and estimate the different energy barriers for nucleation associated with the NP distortion.11,17 It must be noted that most models assume a Received: January 8, 2011 Revised: April 1, 2011 Published: April 18, 2011 1934

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Nano Letters monolayer nucleation and subsequent layer-by-layer growth mode taking place at the NP/NW interface. Among different possibilities to control NW VLS growth, the variations of the vapor supply flow should provide a much better tool since faster system responses can be achieved. This avoids growth interruptions, usually associated with enhanced contamination and generation of wider interfaces in epitaxial thin films. In this work, we present an analysis of the morphology and crystal structure of InP NWs grown by chemical beam epitaxy (CBE). In particular, we have investigated the effects of variations in the vapor supply of the In precursor. Our results suggest different dominant routes for precursor incorporation at the TPL as a function of flow which can lead to interface roughening and defect formation. The model proposed from these results may be exploited to fine control NW crystal structure. InP NWs were grown on GaAs (100) substrates in a CBE chamber, using colloidal Au NPs as the catalyst material. The substrate native oxide was not desorbed prior to the InP growth. Two different particle sizes were used as catalysts (with 10 or 25 nm average diameters, hereafter called small and large NPs, respectively). Trimethylindium (TMI) diluted with hydrogen carrier gas and thermally decomposed phosphine (PH3) were used as group III and V sources, respectively. Growth temperature was chosen as 420 °C in order to provide WZ NWs for the NP sizes used here.22 Two types of samples were grown. For A-type NWs, all growth parameters were kept constant throughout the run. For B-type NWs, growth was carried out under low TMI flow (0.15 sccm) conditions except for short segments where TMI flow was stepped up or down (with a 12 s transition period) to reach the intended value (1.2 sccm) without growth interruption. Two sets of B-type samples were grown, with different numbers of flow transitions. For both types of NWs, PH3 flow was kept at 15 sccm, providing the dominant background pressure for CBE growth. The corresponding P2 overpressure warrants good crystallinity in epitaxial film growth for all TMI flows used here. Samples were cooled down in a PH3 atmosphere in order to minimize any possible effects due to a residual P concentration at the NP after growth, as suggested by previous results.22,23 Several A-type samples were grown with TMI flow between 0.15 and 1.2 sccm; deposition time was varied inversely with TMI flow in order to keep approximately constant the total amount of material supplied for each sample. The samples were subsequently exposed to the environment and prepared for morphological and structural analysis by electron microscopy. The wires were observed with a scanning electron microscope (SEM, model JSM 6330F) to obtain information on NW shape and length. The crystal structure and chemical composition of the NWs were investigated with electron diffraction and high-resolution transmission electron microscopy (HRTEM) as well as Energy Dispersive Spectroscopy (EDS), using a JEM 3010 URP operated at 300 kV and a JEM 2100 ARP operated at 200 kV. The NWs axial growth rate was estimated from SEM images considering only the longer NWs in each sample. The rate was assumed linear and calculated by dividing the measured NW length by the growth time. SEM and TEM images were used to check for the presence of constrictions and defective regions (indicative of the flow transitions, as explained further in the text) as well as the axial length between these regions in B-type NWs. The growth rates thus obtained correlate quite well with those obtained from SEM images for corresponding A-type samples. Axial length values are ∼50 times larger than the thickness of the

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Figure 1. (a, c) Typical TEM images of A-type InP NWs grown under 0.15 and 1.2 sccm TMI flow, respectively (see text for explanations), for large (25 nm) NPs. Scale bars represent 1 μm. (c) Arrows indicate larger variations in diameter along the wires. (b, d) Atomic resolution images of the NWs. A high density of stacking faults is observed for the wire grown under larger TMI flow (d). The plot in (e) shows the average SF density obtained from HRTEM images for the TMI flow condition used.

InP film simultaneously deposited on the GaAs substrate and ∼10 times larger than thicknesses usually found at 500 °C for planar growth using the same TMI flow. Figure 1 illustrates TEM images of A-type InP NWs grown under low (0.15 sccm) and high (1.2 sccm) TMI flows for the large NPs. Electron diffraction indicates that these two samples exhibited WZ crystal structure with growth axis along the [0001] direction.22 The wires grown under low flow show an almost perfect WZ structure, a slight tapering and a smooth lateral surface. In contrast, the wires generated under high TMI flow display large diameter variations along the NW length (arrows in Figure.1c); HRTEM images (Figure.1d) reveal the formation of a high density of SFs, as well as few-monolayer-wide ZB segments along the growth direction. A quantitative analysis shows that SF densities are not statistically significant for the lower TMI flows but reach up to approximately 300 μm 1 for NWs grown under 1.2 sccm, as shown in Figure 1e. Figure 2 shows electron images of wires from B-type samples, for which TMI flow conditions vary between those used for the NWs in Figure 1. These wires also exhibit a WZ structure and a tapered shape; in particular, the wires show large diameter variations in the form of constrictions (or necks) along their length. The axial 1935

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Figure 3. Quantitative results for In atomic concentration at the NP measured by EDS and the NW axial growth rate obtained from SEM statistical analysis considering only the longer NWs in each sample (and with the same orientation). Measured lengths are projections of the actual values for the average NP diameter on our sample. The EDS measurements were carried out on NWs with lengths compatible with the average value from SEM analysis. In practice it has been rather difficult to find long NWs conserving the NP on the TEM grid; thus three nanowires were analyzed in each sample. The lines in the figure are guides for the eye.

Figure 2. TEM images of B-type InP NWs where strong TMI flow variations were intentionally produced: (a) nanowire constriction showing a high density of SFs in the high flow region (inset shows HRTEM of the same region, scale bar = 2 nm); (b) typical NW morphology displaying diameter variations. Scale bar = 2 μm.

distance between these constrictions indicates that they are correlated with the TMI flow variations during growth. Careful analysis of HRTEM images shows that SFs are formed mainly in the wire segments close to the constrictions (see Figure 2a and inset). In particular, these highly defective regions are associated with segments grown under larger TMI flow, in agreement with observations for A-type NWs (Figure 1). The analysis of different B-type NWs reveals that the length and shape of the constricted regions depend on the wire diameter (which is a function of both NP size and lateral growth). Moreover, larger SF densities can consistently be observed when the diameter increases at these regions (see examples in Supporting Information, Figure S1). Another common feature is the correlation of small variations in diameter and SF occurrence. Indeed, this is commonly observed in NWs of many III V materials when ZB/WZ segments are intercalated in the wires.11,12,19,20 Contrarily, NWs grown under low TMI flow usually present diameter variations near the base (Figure 1a). These variations are most likely related to radial growth provided by mass transport from the surface which is enhanced under this particular condition (notice Supporting Information, Figure S2). It is usually observed that defects are generated in NWs at the regions between NW and the metal NP formed during sample cool down;9,20 this is attributed to a reduction of the NP supersaturation after the metal precursor is turned off. Moreover, for different III V materials, the NW/NP neck region usually exhibits a crystal structure opposite to that of the NW; thus WZ

NWs show predominantly ZB necks and vice versa.20 Both kinds of defects, those formed close to the NW/NP interface and those observed in the B-type NWs (Figure 2) occur in a situation where the In supply is varying. Thus, it is interesting to note that NWs grown under constant precursor conditions (or constant NP supersaturation, Figure 1c,d) share similar structural characteristics (i.e., SF formation). We must note that when the TMI flow is increased, the supersaturation level in the whole growth system raises; however, this does not straightforwardly imply that the NP supersaturation has been changed. In order to gather more information on the metal NP for different growth conditions, we have analyzed ex situ the average In atomic concentration at the NP by EDS (Figure 3). The ex situ measurement will certainly provide information that differs from the actual In concentration during growth (we must keep in mind that samples were cooled down under P2 atmosphere). In our case, the set of analyzed samples was grown at the same temperature, with similar cool down conditions. Thus, we expect the postgrowth In concentration at the NP to be directly related to supersaturation levels during NW growth. For NWs grown under constant precursor conditions (Figure 1), EDS shows statistically similar NP supersaturation levels despite the very different TMI flows (or very distinct NW structures) between the two samples shown in Figure.1. This indicates that the increase in SF density for high TMI flow cannot be purely associated with supersaturation levels at the NP. The relationship between NW growth and NP supersaturation is indicated in Figure 3, where we have also plotted the estimated NW axial growth rate. When lower TMI flows (0.1 0.5 sccm) are used for growth, both the growth rate and the In concentration at the NP show similar trends, increasing monotonically. Further increase of the TMI flow over 0.5 sccm generates a significant decrease of both parameters, what clearly defines a maximum in both curves. For TMI flows over 0.7 0.8 sccm, however, the trend of axial growth rate dissociates from the In concentration in the NP. While growth rate shows a clear 1936

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Nano Letters enhancement and larger dispersion with TMI flow increase, the NP In content remains constant (∼25%). The data in Figure 3 clearly indicate that, for large TMI flows (>0.8 sccm), a different route of In precursor incorporation into the NW should become dominant, distinct from the usually accepted Au In alloy generated at the NP, since both NW and NP diameters seem not to be affected in this case (see Supporting Information, Figure S2). The formation of constricted (neck) regions in B-type NWs can be partially explained on the basis of the growth rate and NP composition changes presented in Figure 3. When the TMI flow is raised during growth, the axial growth rate will temporarily increase as well; however, the lateral growth rate at the NW sidewall should not significantly change. As a result, the local tapering should be diminished since the ratio between axial and lateral growth rates will be temporarily larger. As the substrate surface acts as a large material reservoir, the tapering angle will also vary as a function of heights along the NW. Furthermore, the In Au alloy NP composition changes can induce modifications of the contact angle between semiconductor and eutectic droplet. In fact, both effects can lead to the formation of the constricted regions shown in Figure 2; however, only a contact angle change can account for the subsequent NW enlargement. Indeed, this enlargement has been observed even in the absence of a long neck (see Supporting Information, Figure S1). A qualitative evaluation of the relation between lateral and axial growth in our samples is provided by the average base and apex diameters of NWs grown under constant precursor conditions, according to SEM imaging (see Supporting Information, Figure S2). As expected, the contribution of surface material for growth (and tapering) is reduced under larger flows.24 In contrast, for the lower TMI flows used in our study, the base diameters are considerably larger than the apex. Moreover, in this case measurable incubation times are necessary to drive the supersaturation in the dispersed NPs, similarly to the behavior reported by Froberg et al.25 These two observations suggest that the TMI flow in this case is barely enough to maintain the minimum NP supersaturation for NW growth. Also, the contribution of material coming from the substrate should maintain NP supersaturation conditions, since the finite precursor mass transported by the vapor phase should in fact act as a limiting factor for growth. At this point, we must consider that our results indicate that an intermediate TMI flow of 0.45 sccm allows a maximum In supersaturation in the NP, whose concentration agrees with previous results for CBE growth.25 Moreover, for this particular TMI flow a maximum axial growth rate is observed (see Figure 3), while SFs are still kept at low density. Further increase in TMI flow to ∼0.7 sccm induces rather unexpected results: both the In concentration at the NP and the axial growth rate fall down, but the density of SFs increases significantly (Figure 1e). As discussed above, the correlation between ex situ NP In concentration and axial growth rates can thus provide a more complete scenario of what happens during NW growth. We must note that the data included in Figure 3 also imply that supersaturation does not completely determine the structural properties of InP NWs. Our TEM analysis on A-type wires indicated a monotonous increase of SF density with In precursor supply, for samples grown with TMI flow larger than 0.6 sccm. Therefore, similar NP In concentration values (as those obtained at the lower and upper limits of the TMI flow range) can yield completely different SF densities. At the same time, WZ NWs with negligible SF densities can be achieved with different supersaturations (in the case of TMI flows 0.15 and 0.45 sccm).

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In vapor phase techniques such as CBE and MOCVD, variations of the NW axial growth rate must be related to the surface kinetics of metalorganics catalysis.26 28 Indeed, our results for InP NWs show a similar trend to InAs NW grown by MOVPE.29 Attaining a rather large growth rate requires both group III and V atom availability. Even for a liquid NP with a nearly perfect accommodation coefficient, however, catalytic activity could be partially suppressed at temperatures commonly used for NW growth, as observed for TMI pyrolysis on epitaxial substrate surfaces.28 In that case the supersaturation which can be attained within the NP is strongly dependent on the NP surface area. Moreover, at In-rich growth conditions, we should consider possible kinetic limitations imposed on the axial growth rate by group V atoms transport during growth. Thus the drop in axial growth rate at ∼0.5 sccm TMI flow (Figure.3) can be related to reduced precursor availability. In these terms, it is important to consider that the actual source of group V atoms during growth is still under discussion in literature.9,23,30 In fact, several phenomena influencing precursor availability may occur simultaneously during growth, such as limited surface diffusion of species toward the NP, issues on NP size and solubility limits, or the eventual saturation of the metal surface catalytic activity. We would thus expect the NW growth rate to continue to fall down with increasing TMI flow. In contrast, a further increase in TMI flow (>0.8 sccm) induces an increase of growth rate, attaining values even higher than the previous maximum value at ∼0.5 sccm. Surprisingly, despite the significant variations in growth rate, our EDS analysis of In concentration at the NP remains essentially constant for TMI flows >0.8 sccm. A natural question arises: how can we provide more In for NW growth if the NP contribution seems to be saturated? If the metal catalytic activity is diminished or can be no longer increased, In adatoms on the NW sidewalls (which represent the major adatom capture area at shorter diffusion lengths) can contribute to solid semiconductor growth by either directly reaching the TPL enhancing axial growth or incorporating at the sidewalls for lateral growth. The well accepted NW growth scenario supposes axial growth through a single step propagation along the NP/NW interface.17,31 If a competition between different nucleation sites occurs at TPL through both NP vapor and NP semiconductor phases, the growth scenario can be drastically different. In fact, roughening of the interface can take place due to the simultaneous nucleation of different monolayers around the TPL. The relatively long sample cool down can hinder the ex situ observation of such anomalous NP/semiconductor interfaces generated during the dynamic nuclei creation. Nevertheless, our extensive TEM study of NW/NP interfaces has revealed that very different NP/NW interfaces are observed when NWs show very different SF densities. For the WZ material grown at low TMI flow (0.15 sccm, Figure.4a), the HRTEM image shows sharply defined interfaces where incomplete monolayers can be observed at the edges of the NW (inset); this nucleation causes a small distortion of the TPL with no visible phase change. In contrast, NWs grown under large TMI flow (1.2 sccm, Figure.4b), show usually lesswell-behaved interfaces (see also Supporting Information, Figure.S3). In some cases, large distortions and a clear curvature of the interface are observed, as well as different faceting for predominantly ZB regions. The results discussed above suggest that even under constant growth conditions, the NP/NW interface can fluctuate to a much larger level than what is usually assumed for the propagation of a single nuclei along the semiconductor/metal interface. This 1937

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Figure 4. HRTEM images (with same magnification) of A-type NP/ semiconductor interfaces for (a) pure WZ and (b) WZ/ZB NW, grown under 0.15 and 1.2 sccm TMI flow conditions, respectively. Large (25 nm) NPs were used for the growth.

Figure 5. (a) TEM image of B-type InP NW grown with small (10 nm) NP with no constrictions along its length (scale bar = 1 μm). (b) HRTEM image of the marked region shows diameter variations (arrow) along the wire with no associated SFs.

hypothesis is not precluded from reported in situ TEM observations since they are mainly based on single element NW growth31,32 and under supply conditions limited by the microscope environment. On the other hand, Joanny and de Gennes33 have discussed the effect of roughness of a solid surface on wettability by introducing the concept of TPL elasticity and studying its response to external forces, either arbitrary or localized. These authors have predicted that the irregularities found in solid surfaces create random fluctuations on the work of adhesion; these fluctuations act as a force that moves the contact line. In particular, this phenomenological model reveals an equivalence for the two most likely found experimental types of defects: contaminants present on the surface and surface roughness. For the latter case, the variation of work of adhesion with regard to the undisturbed case depends both on the actual contact angle and on the local slope of the surface near the TPL.33 Moreover, if the droplet is at the nanoscale, a correction to the contact angle is induced by TPL tension34 which increases with the length/surface ratio. In these terms, the variations observed in growth rate and the NP In content as a function of TMI flow (Figure 3) result from the competition between different nuclei to incorporate In and P atoms provided from different routes, (i.e., through NP alloy and surface diffusion). As energy considerations do still hold, it must be physically expected that nucleation should occur at TPL and growth proceed by step propagation along the metal/semiconductor interface. However, if kinetic

conditions prevail such that the step propagation rate is smaller than the nucleation rate at TPL, roughness should build up in the region close to the TPL. A similar scenario has been recently suggested for the growth of Si NWs under large disilane pressures using Cu as catalyst.35 The kinetic origin of TPL fluctuations further suggests a behavior which is dependent on the NP size: NPs of smaller diameters should be less susceptible to show interface roughening. In fact, small NPs should attain larger supersaturations for lower precursor flow conditions and the time necessary to complete a monolayer at the NW growth front should be much shorter. Moreover, for the cases of both low P solubility in the NP and diffusion of P from the TPL along the NP/semiconductor interface, we would expect a faster P incorporation at the growth front for smaller NPs. Such effects prevent multiple nucleus formation and certainly contribute to keep a smoother NW/NP interface. Also, more restrictive kinetic conditions should be necessary in order to change the length of the TPL when using small NPs. Our experimental observations confirm these deductions. Indeed, TEM images of B-type InP NWs grown with our smaller NP diameters (10 nm), and under identical conditions as to those in Figures 2 and S1, show pure WZ structures. No statistically relevant SF density could be obtained for observations carried out along the whole length of the NWs; diameter variations can seldom be detected, but they are not associated with SF occurrence (Figure 5). The measured In concentration 1938

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Nano Letters at the NP, though, is similar to those found for our large NPs despite the larger growth rate. It must be emphasized that our TEM results also agree very well with previous reports on the defect density dependence on the ratio of V/III elements. So far, the WZ ZB transition in III V NWs has been mainly attributed to changes in the temperature or V/III ratio.12,19,20 Modifications of these parameters induce a variation of the vapor solid nucleus surface energy, which may, for example, lead to different surface reconstructions on the NW facets. However, these phenomena should not solely account for the rather homogeneous response concerning morphologies and crystal structures of NWs generated with different synthesis techniques based on very dissimilar growth environments. When fluctuations of the TPL occur and roughness builds up, the distortions created at the NP interface should lead to the formation of fault planes in order to reach local minima of system interface and surface energy. In the case of WZ NWs, contact angle variation effects may be minimized and the growth front stability may be maintained by the nucleation of {111} facets which are tilted toward or away from the wire axis. In the particular case of WZ InP NWs studied here, these phenomena must be related to the modulation of NW diameter (Figure 2b) since contact angle hysteresis is expected to occur for strong inhomogeneities causing TPL distortion.33,36 Finally, in our interpretation, crystal phase changes happen as a response to kinetically driven roughening, then it naturally fulfills Poisson statistics for defect formation as observed for III V nanowires.37 It is also interesting to notice that TPL modifications can be triggered by surface contamination.33 For example, Algra et al.11 have created InP twinning superlattices using Zn as impurity dopant. The WZ to ZB transition was attributed to the slightly different solid liquid surface free energy upon zinc addition. More recently, Wallentin et al.38 suggested increased contact angles upon addition of Zn precursor during growth; the changes in wetting were correlated to changes in surface energies which may affect the crystal structure of the NWs, increasing the barrier for WZ nucleation. In these terms, we must analyze our experiments carefully, as a similar role could be assigned to excess carbon atoms under the large TMI flows used here. However, electron energy loss spectroscopy experiments have not revealed the presence of C above detection level. In summary, on the basis of ex situ In atomic concentration at the catalyst particle, we have shown that SF formation in InP NWs is not univocally attributed to the NP supersaturation. Moreover, we have detected that when the wires are generated under growth conditions associated with larger SF densities (for example, under large TMI supply rate) there is a lack of relation between NP In concentration and NW axial growth rate. Our study has shown that SF formation is strongly enhanced when In supply (and NW diameter) variations occur, either from NP depletion or from competition between different precursors incorporation routes. In contrast, for NWs grown under the same conditions with our small NPs, defect formation is suppressed. The ensemble of experimentally observed behaviors can be explained if we take into account kinetically driven roughening and deformation of the TPL due to a competition between different nuclei to incorporate III V atoms from different routes, such as through NP or surface diffusion. This increases the probability of a phase change in the system attempt to minimize planes at the growth interface in order to reach local minima of system interface and surface energy.

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’ ASSOCIATED CONTENT

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Supporting Information. TEM images of nanowires and NP/NW interfaces (S1 and S3, respectively) and measurements of NW base and apex diameters (S2). This material is available free of charge via the Internet at http://pubs.acs.org.

’ AUTHOR INFORMATION Corresponding Author

*E-mail: monica@ifi.unicamp.br.

’ ACKNOWLEDGMENT The authors acknowledge Daniela Zanchet for discussions on NP synthesis and H. T. Obata for technical assistance. SEM and TEM measurements were performed at the Electron Microscopy Facility (LME), Brazilian Synchrotron Laboratory (LNLS). This work was funded by FAPESP and CNPq. ’ REFERENCES (1) Cui, Y.; Lieber, C. M. Science 2001, 291, 851–853. (2) Bj€ork, M. T.; Ohlsson, B. J.; Sass, T.; Persson, A. I.; Thelander, C.; Magnusson, M. H.; Deppert, K.; Wallenberg, L. R.; Samuelson, L. Appl. Phys. Lett. 2002, 80, 1058–1060. (3) Gudiksen, M. S.; Lauhon, L. J.; Wang, J.; Smith, D. C.; Lieber, C. M. Nature 2002, 415, 617–620. (4) Lu, W.; Xiang, J.; Timko, B. P.; Wu, Y.; Lieber, C. M. Proc. Natl. Acad. Sci. U.S.A. 2005, 102, 10046–10051. (5) Wagner, R. S.; Ellis, W. C. Appl. Phys. Lett. 1964, 4, 89–90. (6) Hiruma, K.; Katsuyama, T.; Ogawa, K.; Morgan, G. P.; Koguchi, M.; Kakibayashi, H. Appl. Phys. Lett. 1991, 59, 431–433. (7) Koguchi, M.; Kakibayashi, H.; Yazawa, M.; Hiruma, K.; Katsuyama, T. Jpn. J. Appl. Phys. 1992, 31, 2061–2065. (8) Harmand, J. C.; Patriarche, G.; Pere-Laperne, N.; Merat-Combes, M.-N.; Travers, L.; Glas, F. Appl. Phys. Lett. 2005, 87, 203101. (9) Persson, A. I.; Larsson, M. W.; Stenstr€om, S.; Ohlsson, B. J.; Samuelson, L.; Wallenberg, L. R. Nat. Mater. 2004, 3, 678–681. (10) Akiyama, T.; Sano, K.; Nakamura, K.; Ito, T. Jpn. J. Appl. Phys. Part 2 2006, 45, L275–L278. (11) Algra, R. E.; Verheijen, M. A.; Borgstr€om, M. T.; Feiner, Lou-Fe; Immink, G.; van Enckevort, W. J. P.; Vlieg, E.; Bakkers, E. P. A. M. Nature 2008, 456, 369–372. (12) Caroff, P.; Dick, K. A.; Johansson, J.; Messing, M. E.; Deppert, K.; Samuelson, L. Nat. Nanotechnol. 2009, 4, 50. (13) Krishnamachari, U.; Borgstrom, M.; Ohlsson, B. J.; Panev, N.; Samuelson, L.; Seifert, W. Appl. Phys. Lett. 2004, 85, 2077–2079. (14) Cornet, D. M.; Mazzetti, V. G. M.; LaPierre, R. R. Appl. Phys. Lett. 2007, 90, 013116. (15) Xiong, Q.; Wang, J.; Eklund, P. C. Nano Lett. 2006, 6, 2736–2742. (16) Joyce, H. J.; Gao, Q.; Tan, H. H.; Jagadish, C.; Kim, Y.; Zhang, X.; Guo, Y.; Zou, J. Nano Lett. 2007, 7, 921. (17) Joyce, H. J.; Leung, J. W.; Gao, Q.; Tan, H. H.; Jagadish, C. Nano Lett. 2010, 10, 908. (18) Johansson, J.; Wacaser, B. A.; Dick, K. A.; Seifert, W. Nanotechnology 2006, 17, S355–S361. (19) Paiman, S.; Gao, Q.; Tan, H. H.; Jagadish, C.; Pemasiri, K.; Montazeri, M.; Jackson, H. E.; Smith, L. M.; Yarrison-Rice, J. M.; Zhang, X.; Zou, J. Nanotechnology 2009, 20, 225606. (20) Dick, K. A.; Caroff, P.; Bolinsson, J.; Messing, M. E.; Johansson, J.; Deppert, K.; Wallenberg, L. R.; Samuelson, L. Semicond. Sci. Technol. 2010, 25, 024009. (21) Glas, F.; Harmand, J.-C.; Patriarche, G. Phys. Rev. Lett. 2007, 99, 146101. (22) Tizei, L. H. G.; Chiaramonte, T.; Ugarte, D.; Cotta, M. A. Nanotechnology 2009, 20, 275604. 1939

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dx.doi.org/10.1021/nl200083f |Nano Lett. 2011, 11, 1934–1940