Simultaneous Growth of Pure Wurtzite and Zinc Blende Nanowires

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Simultaneous growth of pure wurtzite and zinc blende nanowires Sebastian Lehmann, Jesper Wallentin, Erik K. Mårtensson, Martin Ek Rosén, Knut Deppert, Kimberly A. Dick, and Magnus T Borgström Nano Lett., Just Accepted Manuscript • Publication Date (Web): 19 Mar 2019 Downloaded from http://pubs.acs.org on March 19, 2019

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Simultaneous growth of pure wurtzite and zinc blende nanowires Sebastian Lehmann1,*, Jesper Wallentin1,3, Erik Mårtensson1, Martin Ek2, Knut Deppert1, Kimberly A. Dick1,2, Magnus T. Borgström1

1

2

Solid State Physics and NanoLund, Lund University, Box 118, S-221 00 Lund, Sweden,

Centre for Analysis and Synthesis, Lund University, Box 124, 221 00, Lund, Sweden, 3

Synchrotron Radiation Research and NanoLund, Box 118, S-221 00 Lund, Sweden

ABSTRACT: The opportunity to engineer III-V nanowires in wurtzite and zinc blende crystal structure allows for exploring properties not conventionally available in the bulk form as well as opening up the opportunity for use of additional degrees of freedom in device fabrication. However, the fundamental understanding of the nature of polytypism in III-V nanowire growth is still lacking key ingredients to be able to connect the results of modelling and experiments. Here we show InP nanowires of both pure wurtzite and pure zinc blende grown simultaneously on the same InP [100]-oriented substrate. We find wurtzite nanowires to grow along 〈111〉B and zinc blende counterparts along 〈111〉A. Further, we discuss the nucleation, growth, and polytypism ACS Paragon Plus Environment

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of our nanowires against the background of existing theory. Our results demonstrate firstly, that the crystal growth conditions for wurtzite and zinc blende nanowire growth are not mutually exclusive, and secondly, that the interface energies predominantly determine the crystal structure of the nanowires.

KEYWORDS Semiconductor, nanowire, wurtzite, zinc blende, transmission electron microscopy, polytypism

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Nanowires can nowadays be successfully implemented in semiconductor devices as has been exemplarily proven including for example in optoelectronics1–3 and sensors4–6. III-V nanowires usually show two common characteristics, which were both recognized already in pioneering works two decades ago7,8. First, the nanowires preferentially grow along 〈111〉B, that is, the group V- terminated surfaces. Second, the nanowires frequently adopt a polytypic behaviour, i.e., a mixed crystal structure of wurtzite and zinc blende stacking. In this paper, we show that both of these features are related to the interfacial energies by using the A- and B-polar directions on a [100]-oriented substrate. The natural tendency for III-V nanowires to exhibit a mixture of wurtzite and zinc blende stacking represents both a challenge for control over synthesis as well as an opportunity for designing homomaterial heterostructures. Precise tailoring of either wurtzite or zinc blende phase allows for an additional degree of freedom in III-V technology as these crystal phases adopt different properties like band gap energy9–12, electrical properties13, and morphologies14–17. Different nanowire terminating facets are connected to these morphologies which is in turn of high relevance for synthesis of core-shell architectures17–21. An axial stacking of alternating wurtzite and zinc blende segments allows controlled tailoring of e.g. crystal phase quantum dots22–26. On the other hand, random crystal phase mixing is typically undesirable since stacking defects can adversely affect the nanowire device performance27–30. The polytypic behaviour was reported for III-V nanowires already in early works7. In this and other studies it was recognized that growth along 〈100〉 or 〈111〉A led to pure zinc blende formation31–33; however the yield of straight nanowires was much lower in other directions than 〈111〉B8,34,35. Today, the control over crystal phases in III-V nanowires has matured, with examples of high crystal quality single phase36 and refs. therein as well as heterostructured configurations i.e. the tailored axial stacking of different crystal phase in a homomaterial system25,37–40. Various parameters have been modified in order to control the crystal structure of resulting nanowires like growth temperature and nominal [V]/[III]-ratio of the incoming precursor fluxes41,42, seed particle diameter43, total precursor flows, and in-situ doping13,38,44–46. ACS Paragon Plus Environment

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Despite the success of

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crystal phase control in III-V nanowires, enabled by in-situ growth

experiments47,48, still key growth experiments are lacking contributing to a generalized picture of the polytypic nature of III-V nanowires; a picture which should be capable of consolidating existing experimental works with detailed theoretical understanding49. The experimental results have been interpreted in different ways, which can be grouped into effects predominantly related to supersaturatione.g.50 – with higher values of supersaturation to promote wurtzite formation – and interfacial energiese.g.44,51 – with an increasing ratio of liquid-solid/liquid-vapour interface energies to promote zinc blende formation. The key limitation of previous experiments has been the indirect relationship between the theoretical models and the employed experimental parameters, where for instance a variation in growth temperature affects both interfacial energies and supersaturation, alternatively growth was not carried out simultaneously on identical substrates52. Here we report on overcoming this limitation by simultaneous, particle-assisted epitaxial growth of pure wurtzite and zinc blende nanowires on the same substrate, using monodisperse Au seed particles. To further evaluate the general character of our findings, we have carried out experiments with and without the presence of hydrogen sulphide (H2S), which is known to strongly influence the nanowire crystal structure46. In both cases, with and without the presence of H2S, both wurtzite and zinc blende nanowires grow under identical conditions on the same substrate, only differing in the polarity of the growth direction and thus the liquid-solid interfacial energy. The simplicity of the experiment allows us to conclusively attribute the nanowire crystal structure to differences in interfacial energies under these parameters. Pure wurtzite and pure zinc blende nanowires – the two structural extremes of III-V nanowires – were prepared in the same growth run, i.e., under identical growth conditions, on the same [100]-oriented InP substrates using metal-organic vapour phase epitaxy (MOVPE). Scanning electron microscopy (SEM) revealed that these nanowires grew in four different directions with 90° rotational symmetry in top view (Fig. 1a,b) i.e. corresponding to a projection onto the [100]-oriented substrate. Further, they adopted an inclination angle of about 35° (Fig. 1d,e) with respect to the {100}-type substrate surface8 which is additionally indicated by the corresponding stereographic projection in the Wulff plot in Fig. 1c. Thus, ACS Paragon Plus Environment

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we conclude that the nanowires grow simultaneously in both the 〈111〉- and 〈111〉-type directions i.e. the A- and B-polar crystallographic directions labelled 〈111〉A and 〈111〉B. In addition, Fig. 1b shows two nanowires of different diameter, although monodisperse Au seed particles were used. This can occur when seed particles become mobile and merge during the high-temperature stage before growth53. We did not observe any correlation between nanowire diameter and crystal structure.

Figure 1. InP:S nanowire morphologies and growth directions. A top view image (a) and a higher magnification inset (b) reveals the perpendicular orientation between the two kinds of nanowire types and proves the growth in all four upward directed 〈111〉/〈111〉-type directions. (c) displays a Wulff plot including the corresponding crystallographic directions of nanowire growth directions and substrate orientation as revealed by TEM investigations (see Fig. 2). Two different nanowire morphologies with smooth side facets and alternating segmented side facets (a, b) are visible. Cross-sectional (2° tilted) SEM micrographs recorded perpendicular to [011] (d) and [011] (e) directions, respectively. A ~35° inclination angle between nanowire growth direction and the substrate surface is indicated. ACS Paragon Plus Environment

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The SEM investigation also revealed two clearly distinct NW morphologies. One type adopted flat side facets (Fig. 1b, d) while the other one adopted periodically alternating side facets (Fig. 1b, e). The nature of the two distinct morphologies is correlated to the crystal structure as revealed by transmission electron microscopy (Fig. 2a,b). Wurtzite crystal phase nanowires adopt smooth side facets, while zinc blende nanowires adopt a periodically alternating morphology, in line with previous work38. In addition, the zinc blende nanowires also show stacking defects in the radial growth on the side facets, which forms in a vapor-solid mode after the axial vapor-liquid-solid growth. Convergent beam electron diffraction (CBED) was applied in order to determine the growth directions for wurtzite (Fig. 2c) and zinc blende (Fig. 2d) nanowires. Comparison with simulated data (Fig. 2e) showed that wurtzite nanowires grew along the Bpolar 〈111〉B-type directions54 and zinc blende nanowires along the A-polar 〈111〉A-type directions of the substrate. For clarity we will denote all nanowire growth directions according to the zincblende crystallographic notation of the substrate, e.g. 〈111〉B-type will be used also when discussing the wurtzite nanowires with a [0001]B growth direction in the wurtzite crystallographic notation. TEM micrographs are however labelled in accordance with the crystallography of the nanowires themselves. Here we would like to note that there is no clear difference in length, and therefore growth rates, between the two types of nanowiresIn order to deepen the understanding of the polytypic nature of III-V nanowire growth, we will in the following interpret our findings in the framework of existing models for polytypism in nanowires49.


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Figure 2. Crystal structure and polarity determination of InP:S nanowires, overview, high resolution images, and selected area diffraction patterns (SADPs) of wurtzite (a) and zinc blende (b) InP:S nanowires recorded by transmission electron microscopy (TEM). (c) and (d) show convergent beam electron diffraction (CBED) patterns recorded in a 〈112〉 direction from nanowires of the same type as the ones shown in (a) and (b), respectively. A corresponding simulation for 54 nm specimen thickness is given in (e) indicating characteristic intensity differences between the polar 111A/0002A and 111B/0002B discs. Scale bar of the CBED patterns is 2 nm-1 and insets in (c) and (d) show the orientation of the nanowires for the CBED acquisition.

Before discussing the actual nanowire growth, the interaction between seed and substrate will be elucidated, as it is a crucial part of successful straight, epitaxial nanowire growth. We performed experiments where we annealed samples with Au seed particles, without initiating InP growth. Upon preconditioning at 550°C in a PH3/H2 atmosphere, the Au seed particle alloys with the [100]-oriented InP substrates to form an AuxIn1-x-alloy55 (Fig. 3a,b). This process requires the dissolution of InP and results ACS Paragon Plus Environment

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in the formation of a pit in the substrate31. Together, the alloy particle and the pit represent the equilibrium situation as defined by the equilibrium crystal shape of the substrate material in combination with the surface tension of the particle under the applied conditions. It is the lowest interfacial energy configuration which is formed at the particle/substrate interface56 under the used experimental conditions. Most of the particles adopted an asymmetric shape with an elongation along the [011]/[011] directions of the substrate (Fig. 3a,b) – according to the major and minor flat directions of the wafer used. The asymmetric particle shape is stemming from smaller {111}A facets compared to the {111}B counterparts56 i.e. the areal A{111}A

facet ratio being smaller than 1, A{1 1 1} < 1. In accordance with theoretical studies57,58, this indicates that B

the liquid-solid interface energy of the {111}A interface is higher than that of the corresponding {111}B interface: 𝛾ALS > 𝛾BLS.

Figure 3. Schematics of nanowire nucleation and growth with (a) elongated shapes of AuxIn1x-alloy

particles on a [100]-oriented InP substrate after PH3/H2 annealing are visible in the SEM

top view image. (b) A schematic figure of a half-cut AuxIn1-x-alloy particle on a [100]-oriented substrate indicating the underlying {111}A- (green) and {111}B-type (red) facets formed due to ACS Paragon Plus Environment

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Au particle alloying with the substrate. (c) Schematic view of larger contact angles arising for nanowire growth in 〈111〉A-compared to 〈111〉B-directions due to different interface energetics, with (d) a more detailed description of the involved vapour-solid, vapour-liquid, and liquid-solid interface energies including (e) the nucleus formation at the vapour-liquid-solid triple-phase line.

Next, we discuss the yield of straight nanowire growth in the 〈111〉A and 〈111〉B directions. Please note that Fig. 1a was chosen to present the growth in all 4 out of plane 〈111〉/〈111〉-type directions, rather than giving a representative image of the statistics for 〈111〉A- vs. 〈111〉B-oriented growth. In our experiment we counted an overall yield of 50% straight but inclined nanowires, with the remaining 50% of nanowires being kinked and curled (see supporting information Fig. 2a). Out of the fraction of straight nanowires, about 93% were wurtzite nanowires oriented along 〈111〉B. The remaining 7% of initially straight nanowires were zinc blende 〈111〉A-grown nanowires, for which we observed a kinking rate of about two thirds. We attribute this to the particle exceeding the critical contact angle, as will be discussed in detail in the following paragraphs (see e.g. Fig. 1a and supporting information S2). In contrast, not a single kinked nanowire was observed for growth along 〈111〉B. On the other hand, the strong preference for growth along 〈111〉A/〈111〉B-type directions as opposed to growth in other crystallographic directions, especially along 〈100〉, reflects the unfavourable interface energetics of {100}-type or other facets with the alloy particle under these conditions. However, tailored pre-growth annealing has been reported as a suitable option to suppress growth along 〈111〉A/〈111〉B directions and instead promote vertical [100]oriented growth31,33,54. The low density of straight 〈111〉A nanowires in our experiments, similar to previous studies reporting on the challenge of growing along 〈111〉A34, could be due to two scenarios which will be discussed in the following: First, an unstable seed particle and secondly, a low probability of nucleation.


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Firstly, the seed particle stability is described by the so-called Nebol’sin criterion, which determines the range of interfacial energies which can lead to stable, steady-state nanowire growth59. In general, it is experimentally observed that, in combination with the commonly found AuxIII1-x-alloys, {111}B-type interfaces normally result in reasonable liquid-solid interface energies 𝛾LS and contact angles 𝛽 which enable nanowire growth. However, due to their higher free surface energies, {111}A-type facets have higher 𝛾LS and contact angles 𝛽, which under comparable growth conditions result in unfavourable interface energetics. The higher 𝛾LS makes it comparatively more favourable for the seed to wet the side facets59 and leads to crawling/kinking.44,60 This we observed for all growth conditions (Fig. 1a) and was enhanced at the non-optimized conditions (Supporting Information S1, S2). Post-growth contact angles of 𝛽ZB = 153° ± 10° were determined for 〈111〉A-grown nanowires (Supporting Information S2). This finding supports the assumption that the contact angle during growth might have adopted values close to its critical value61 accounting for the effects of increased wetting at elevated temperatures62 and potential material consumption from the particle upon neck formation during the cool-down procedure. Secondly, we will focus on the probabilities of nanowire nucleation in the 〈111〉A- and 〈111〉B-substrate related directions commencing with the discussion of the interface energetics. We consider a model for particle/nanowire interface energetics, 44,49 as shown in Fig. 3c-e, since we only characterize successfully nucleated and grown nanowires. At the time of nucleation of the nanowire, the seed particle has a single supersaturation and is in contact with two {111}A and two {111}B interfaces. From the discussion above, it is clear that the interfacial energy of {111}A is higher compared to {111}B58: 𝛾ALS > 𝛾BLS. Further, we assume that the liquid-vapour interface energy 𝛾LV is identical for seed particles on nanowires grown along both polarities

(𝛾ALV = 𝛾BLV),

since with an identical surrounding vapour phase it is

thermodynamically supported that all seed particles will have similar compositions. Following Young’s 𝛾LS

equation for the alloy particle we get cos𝛽 = ― 𝛾LV

63

and hence, larger contact angles are expected for

wires grown in 〈111〉A direction with respect to the 〈111〉B counterparts (𝛽A > 𝛽B). This agrees with our


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post-growth observations of βAZB = 153° ± 10° and βBWZ = 113° ± 10° (see Figs. 1c, 2a,b and Supporting Information S2). Further, the increase in surface energy of a nucleus, as given according to Ref.49 and simplified for the specific case,44 is: Γ = (1 ― 𝑥)𝛾lLS ― 𝑥𝛾LVsin 𝛽 + 𝑥𝛾SV, where x is the area of the nucleus in contact with the vapour and is quite small. In addition, the lateral liquid-solid interfacial energies, 𝛾lLS, are assumed identical49 for the two different polarities and growth directions. For the contact angles we have 𝛽𝐴 > 𝛽𝐵 > 90°, which means that the second term is smaller for A than for B. Therefore we find for the nuclei surface energies that Γ𝐴 > Γ𝐵 and it is hence more favourable to nucleate nanowires in 〈111〉B direction than in 〈111〉A. The probability for successfully nucleating nanowires in 〈111〉A- vs. 〈111〉B-directions is further reduced by the geometrical constriction A{111}A

of the ratio between the {111}A and {111}B facet areas (A) formed under the particle, A{111} < 1. With an B

experimentally determined ratio of 0.93 for 〈111〉B- vs. 〈111〉A-oriented nanowires (Supporting Information S2) we find a good agreement to our interpretation which additionally coincides with a previous report on GaAs nanowire growth in a similar experiment on [100]-oriented GaAs substrate34. To conclude this section, a combination of the alloy particle and {111}B-type surfaces seems to be the favourable configuration for stable straight nanowire growth since it represents the lowest-energy liquidsolid interface. We speculate that it is the first, namely the unfavourable wetting of the {111}A facets, which represents the dominating effect for the low yield of straight zinc blende 〈111〉A-oriented nanowires. This represents the first of the presented explanations, which all stem from the difference in interface energies. Our speculation is further supported by the fact that a significant fraction of the zinc blende nanowires kinked after having initially grown straight (Supporting Information S2) which implies that some of the observed kinked/curled events might simply stem from failed growth along 〈111〉Adirections.


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After having discussed the nucleation probabilities for the different polarities, we now focus on the observed striking difference in structural properties i.e. wurtzite in B-polar and zinc blende in A-polar directions. We assume that the polarity of the growth direction for a given nanowire is fixed during the first nucleation, as discussed in the previous section, but that the crystal structure is governed by each nucleation of a new layer. In general, the ratio between the nuclei surface energies for wurtzite and zinc blende is given by44,49: 𝜂=

ΓWZ

=

ΓZB

(1 ― 𝑥)𝛾lLS ― 𝑥𝛾LVsin 𝛽 + τ𝑥𝛾ZB 𝑆𝑉 (1 ― 𝑥)𝛾lLS ― 𝑥𝛾LVsin 𝛽 + 𝑥𝛾ZB 𝑆𝑉

Again, we assume that the lateral γlLS , unlike the horizontal one, is independent of the polarity A/B, just like τ, which is the ratio between the solid-vapor surface energies for wurtzite and zinc blende (τ 𝛽𝐵, as discussed in the previous section, which means that the magnitude of the second term in both surface energies is lower for the A polarity. Since the other terms are independent of polarity, the surface energies for both wurtzite and zinc blende are larger for A polarity with a similar ratio of

γWZ LS γZB LS

for both polarities64. This we interpret

as an indication that 𝜂A > 𝜂B (𝜂A closer to 1). Next, the ratio (𝜉) between the nucleation barriers for wurtzite (𝐺WZ) and zinc blende (𝐺ZB) can be written44: 𝜉=

𝐺WZ 𝐺ZB

=

Δ𝜇LS𝜂2 Δ𝜇LS ― ΨWZ

With 𝜂A > 𝜂B, for a similar supersaturation in the difference of the liquid-solid chemical potential (Δ𝜇LS) and a fixed difference for the wurtzite cohesive energy (ΨWZ)65, this increases: 𝜉A > 𝜉B. That is, for the


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same Δ𝜇LS, there will be less wurtzite and more zinc blende in 〈111〉A and vice versa for growth along

〈111〉B i.e. more wurtzite and less zinc blende, which agrees very well with our observations. An alternative explanation could be that the supersaturation was higher during growth of the B polarity nanowires compared to the A counterparts

50,66.

Previous experiments have tuned growth parameters to

control the crystal structure36 and refs. therein, but a key conceptual disadvantage of such an approach is that changing the growth parameters can affect both the supersaturation and the interfacial energies. The unique advantage of our experiment is that both polarities grow under exactly the same conditions on the identical substrate. Differences in surface diffusion on WZ and ZB nanowires, due to the different side facets, could change the supply of In to the liquid particle for 〈111〉A- and 〈111〉B-oriented nanowires, which could change crystal structure. However, this would also be expected to change growth rate, while our results show similar growth rates for 〈111〉A- and 〈111〉B-grown nanowires (Fig. 1) and phase purity for both polarities. Another alternative explanation could be the position of nucleation. While nucleation at the triple phase line is generally understood to be the reason why wurtzite can form, zinc blende formation can occur via centre nucleation in addition to triple phase line nucleation50,52,67. However, we would expect to observe stacking defect free zinc blende for growth along 〈111〉A52,68 rather than closeto periodic twinning as in our case(s). We thus think that our findings support the conclusion that it is the change in the interface energetics of the vapour-liquid-solid system at the growth interface of the nanowire that induces the different crystal phases in the〈111〉A- and 〈111〉B-oriented nanowires. In the results presented up to now we have discussed nanowires grown by in-situ use of H2S. However, sulphur is known to decrease both 𝛾LV and 𝛾LS for liquid metals68 which could result but does not necessarily have to result in a lower contact angle, since it depends on the ratio 𝛾LV/𝛾LS46. An additional aspect of using sulphur doping is that the {111}A and {111}B growth facets are polar, i.e. they have different terminations, which could potentially result in a different susceptibility to S-passivation and hence to a different effect on the respective 𝛾A/B LS . Likely, the In-terminated (cation) {111}A facets are more probable to interact strongly with S. In order to generalize our interpretation of the interfacial ACS Paragon Plus Environment

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energies being the determining factor for the crystal structure of the grown nanowires, we carried out similar experiments under sulphur-free conditions. We used a growth temperature of 480 °C and otherwise identical conditions as reported for optimized phase pure wurtzite InP 〈111〉B-oriented growth51 – instead of 400 °C as used for the InP:S growth presented here. I.e. even for growth conditions which originally were optimized to result in pure wurtzite for growth on 〈111〉B-oriented substrates, we observed identical results regarding the crystal structural and morphological properties of the InP nanowires grown on [100]oriented substrates i.e. wurtzite for 〈111〉B- and zinc blende for 〈111〉A-grown nanowires (Figs. 4a,b and Supporting Information S3). Note that the yield of successful nanowire growth was lower than in the InP:S case which we attribute mainly to increased wetting on 〈100〉-type substrates upon the increased growth temperature62 favouring crawling of the Au particle alloy rather than nanowire growth since the Nebol’sin condition59 is not fulfilled (Supporting Information S3). In addition a zinc blende neck can be noticed directly under the particle in Fig. 4 a.. This indicates residual axial nanowire formation during the cooling down procedure for this experiment, which thus reflect conditions not representative of the growth itself.

Figure 4. Crystal structure and polarity determination of sulphur-free InP nanowires. Overview, high resolution images, and selected area diffraction patterns (SADPs) of wurtzite (a) and zinc


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blende (b) InP nanowires recorded by transmission electron microscopy (TEM). (c) and (d) show convergent electron beam diffraction (CBED) patterns recorded in a 〈112〉 direction from the same type of nanowires as shown in (a) and (b), respectively. A corresponding simulation for 54 nm specimen thickness is given in (e) indicating characteristic intensity differences between the polar 111A/0002A and 111B/0002B discs. Scale bar of the CBED patterns is 2nm-1 and insets in (c) and (d) show the orientation of the nanowires for the CBED acquisition.

In view of our results it seems not coincidental that first, Au is the prevailing seed material for nanowire growth and second, polytypism is frequently reported in Au-seeded III-V nanowire growth along 〈111〉B. The specific surface tension of gold68 in combination with the low surface energy of the {111}B-type facets of III-V compound materials57,58 allows for stable nanowire growth. This combination of Au and III-V material preferentially promotes growth along 〈111〉B which represents a local thermodynamic minimum. At the same time this materials combination covers a broad range of possible interface energetics/stable contact angles59,61 to allow for the observed polytypic behaviour. For alternative seed particle materials and/or growth directions, stable growth and polytypic behaviour might be limited due to unfavourable interface energetics as compared to the case of using Au-seeds. In summary, we show that pure wurtzite and pure zinc blende crystal phases – the two structural extremes of III-V nanowires – can be prepared simultaneously under identical growth conditions. Our results demonstrate that interface energies strongly influence the crystal structure properties of III-V nanowires and that the experimental conditions for wurtzite and zinc blende InP nanowire growth are not mutually exclusive. We anticipate that the observed mechanism is general for particle assisted nanowire growth of different materials, as indicated by previous experiments51.


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METHODS Nanowire growth. InP nanowires were grown by metal-organic vapour phase epitaxy (MOVPE) on Fe-doped semi-insulating [100]-oriented InP wafers manufactured by Wafer Technology. Nanowire growth followed the particle-assisted mode using Au particles which were deposited by aerosol technique69. The nominal particle diameters and aerial densities were 80 nm/0.2 m-2 for InP:S growth (50 nm/0.5 m-2 and 80 nm/4 m-2 for selected experiments shown and indicated in the supplementary information) and 20 nm, 40 nm, 50 nm, 60 nm, 90 nm at a total density of 1.0 m-2 for sulphur-free growth, respectively. Sulphur-containing growth was carried out in an EpiQuip system while sulphur-free growth was carried out in an AIXTRON 200/4 setup at total gas flows of 6 slm and 13 slm, respectively. In order to remove surface oxides and allow proper substrate preconditioning a 10 min annealing step was carried out prior to growth above 550 °C, 100 mbar, and in a PH3/H2 atmosphere. After that step the temperature was set to the growth temperature – 400 °C for InP:S growth and 480 °C for sulphur-free InP nanowire growth. After thermal stabilization the precursors were introduced to initiate growth. The molar fractions of the incoming precursors were set to χTMIn= 3.8x10-6, χPH3= 1.2x10-2, and -6 χH2S= 6.5x10 for sulphur-containing conditions with a 30 sec sulphur-free nucleation step and


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to χTMIn= 7.0x10-6 and χPH3= 5.4x10-5 for sulphur-free growth conditions. Samples were cooled in a PH3/H2 mixture after growth. Microscopy. SEM characterization was carried out in a ZEISS Leo Gemini 1560 equipped with a field emission gun (FEG) and operated at 15 kV. For structural characterization by transmission electron microscopy (TEM), nanowires were placed on copper grids covered with a lacey carbon layer and investigated in a JEOL-3000F equipped with a FEG and operated at 300 kV. The polarity of the growth directions were determined from comparisons of recorded and calculated ⟨112⟩ convergent beam electron diffraction (CBED) patterns. In this orientation the ZB and WZ patterns are identical. A 3 mrad convergence semi-angle was used in both cases. The simulations were carried out in JEMS (Electron Microscopy Software, Java version) using the Bloch wave method in thickness increments of 5 nm.

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ASSOCIATED CONTENT Supporting Information The Supporting Information is available free of charge on the ACS Publications website at DOI: xxx. 

Effect of PH3- and TMIn-variation on InP:S nanowire growth on [100]-oriented InP



Nanowire yield and ex-situ contact angle analysis



InP nanowire growth without H2S on [100]-oriented InP.

AUTHOR INFORMATION Corresponding author * E-mail: [email protected] Author contributions S.L and J.W. made the growth experiments. S.L and M.E. performed the electron microscopy and the data analysis. S.L., E.K.M., and J.W. discussed the results. S.L and J.W. wrote the paper with support of all the other authors. K.D., K.A.D. and M.T.B. provided guidance through the project. All authors have given approval to the final version of the manuscript. ORCID Sebastian Lehmann 0000-0002-4091-905X Jesper Wallentin 0000-0001-5909-0483 Erik K. Mårtensson 0000-0002-2181-2177


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Martin Ek 0000-0002-5705-8495 Knut Deppert 0000-0002-0471-951X Kimberly A. Dick 0000-0003-4125-2039 Magnus T. Borgström 0000-0001-8061-0746 Notes The authors declare no competing interests.

Acknowledgments S.L. gratefully acknowledges the support by a fellowship within the Postdoc-Programme of the German Academic Exchange Service (DAAD). The authors acknowledge financial support from NanoLund, the Swedish Research Council (VR), the Swedish Foundation for Strategic Research (SSF), the Swedish energy agency, and the Knut and Alice Wallenberg Foundation (KAW).

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Simultaneous Growth of Pure Wurtzite and Zinc Blende Nanowires

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